Preface

Chapter 7 **Selective Laser Sintering of Nanoparticles 147**

**Functional Materials 165** Wallace Matizamhuka

Chapter 8 **High-Pressure High-Temperature (HPHT) Synthesis of**

Sukjoon Hong

**VI** Contents

Powder-based materials and treatment technologies rank high in contemporary scientifictechnical progress due to their numerous significant technoeconomic qualities. Sintering of such materials allows saving on materials and lowering the cost price of the product, as well as manufacturing complex composite materials with unique combinations of qualities. Ma‐ terials of record high values of some physic-mechanical and also biochemical characteristics can be obtained owing to structural peculiarities of super dispersed condition.

Sintering of functional materials for innovative perspectives in automotive and aeronautical engineering, space technology, lightweight construction, mechanical engineering, modern design, and many other applications requires established relationship in the materials-proc‐ ess-properties system. Therefore, the industry being interested in understanding theoretical modeling and control over behavior of such powdered materials has promoted the research activities of this manuscript's authors.

The authors of the chapters hope that this book will serve as introduction into the world of sintering of functional materials for the graduate and postgraduate students in materials sci‐ ence and related disciplines.

> **Prof. Igor Shishkovsky** Lebedev Physical Institute of the Russian Academy of Sciences, Samara Branch, Samara, Russian Federation

**Section 1**

**Sintering of Functional Materials**

**Sintering of Functional Materials**

**Chapter 1**

**Provisional chapter**

**Two-Step Sintering of Ceramics**

Murugathas Thanihaichelvan and

**Two-Step Sintering of Ceramics**

Murugathas Thanihaichelvan and Ramesh Singh

DOI: 10.5772/68083

Sintering is a critical phase in the production of ceramic bodies. By controlling the density and microstructure formation, sintering now emerged as a processing technology of ceramic materials. Tailoring the structural, mechanical, electrical, magnetic and optical properties is widening the application of ceramics in various fields. Recently, many advanced sintering methods have reported to fabricate ceramic materials with controlled properties. Two‐stage sintering (TSS) is one of the simple and cost‐effective methods to obtain near‐theoretical density materials with controlled grain growth without adding any dopants. Many recent works have reported the use of TSS as a processing method to fabricate nanoceramics for various applications. With this background, this chapter reviews the advantages of TSS in ceramic preparation based on properties and materials and explores the future directions.

**Keywords:** two‐step sintering, grain growth, ceramic properties, densification,

© 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.

Highly dense ceramics with smaller grain size are widely used in high‐performance applica‐ tions in extreme conditions. Sintering is the responsible step for densification of ceramic bod‐ ies, and due to its influence on the properties of the material, sintering is also emerging as a new fabrication method. Controlling the powder size [1], use of sintering additives [2], green body density [3–5] and sintering environment [6, 7] and using new sintering methods such as microwave sintering [8, 9], pressure‐assisted sintering [10], spark plasma sintering [11] and field‐assisted sintering [3] are used for fabrication of dense and fine‐grained ceramics. But these may destroy the unique properties of ceramics [12]. Also the applications of new

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

Ubenthiran Sutharsini,

Ramesh Singh

Ubenthiran Sutharsini,

http://dx.doi.org/10.5772/68083

sintering mechanism

**1. Introduction**

**Abstract**

## **Chapter 1**

**Provisional chapter**

## **Two-Step Sintering of Ceramics**

Murugathas Thanihaichelvan and

**Two-Step Sintering of Ceramics**

Ubenthiran Sutharsini, Ramesh Singh

Ubenthiran Sutharsini,

Murugathas Thanihaichelvan and Ramesh Singh

Additional information is available at the end of the chapter Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/68083

#### **Abstract**

Sintering is a critical phase in the production of ceramic bodies. By controlling the density and microstructure formation, sintering now emerged as a processing technology of ceramic materials. Tailoring the structural, mechanical, electrical, magnetic and optical properties is widening the application of ceramics in various fields. Recently, many advanced sintering methods have reported to fabricate ceramic materials with controlled properties. Two‐stage sintering (TSS) is one of the simple and cost‐effective methods to obtain near‐theoretical density materials with controlled grain growth without adding any dopants. Many recent works have reported the use of TSS as a processing method to fabricate nanoceramics for various applications. With this background, this chapter reviews the advantages of TSS in ceramic preparation based on properties and materials and explores the future directions.

DOI: 10.5772/68083

**Keywords:** two‐step sintering, grain growth, ceramic properties, densification, sintering mechanism

### **1. Introduction**

Highly dense ceramics with smaller grain size are widely used in high‐performance applica‐ tions in extreme conditions. Sintering is the responsible step for densification of ceramic bod‐ ies, and due to its influence on the properties of the material, sintering is also emerging as a new fabrication method. Controlling the powder size [1], use of sintering additives [2], green body density [3–5] and sintering environment [6, 7] and using new sintering methods such as microwave sintering [8, 9], pressure‐assisted sintering [10], spark plasma sintering [11] and field‐assisted sintering [3] are used for fabrication of dense and fine‐grained ceramics. But these may destroy the unique properties of ceramics [12]. Also the applications of new

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. © 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.

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons

sintering techniques are limited by low mass productivity, and they are not economically feasible. Two‐stage sintering (TSS) is an effective way to achieve fine‐grained microstructured ceramics with high densification and relatively low cost. TSS method is successfully applied for all types of ceramics such as structural ceramics, bioceramics, ferrites, piezoelectric ceram‐ ics and electrolyte ceramics. Most of the ceramics exhibit controlled or no grain growth in the final stage of sintering and achieved near‐theoretical densities. The fine‐grained micro‐ structure enhances the mechanical, electrical, magnetic as well as piezoelectric properties of ceramics which widen the applications of ceramics.

**2. TSS with high first‐step sintering temperature**

studies, and others hold for a few minutes at temperature TP

The successful TSS method using profile 1 was first introduced for Y<sup>2</sup>

O3

firstly proposed by Chen and Wang [15], for a TSS study on Y<sup>2</sup>

0.05K0.5Na0.5NbO3

• Holding time (t1

ics such as Mg, Nb‐doped Y<sup>2</sup>

times and sintering atmospheres.

accepted and verified for other ceramics.

**2.1. Sintering mechanism**

0.06BaTiO3‐

0.89Bi0.5Na0.5TiO3‐

TSS with a high first‐step sintering temperature is widely used to obtain fully dense ceram‐ ics with controlled grain growth due to lower second‐step sintering temperature. The sin‐ tering profile for TSS with higher first‐step sintering temperature is depicted in **Figure 1(a)**.

• Cooling rate (β): Few studies assumed sample rapidly cooled from first sintering tempera‐ ture (TP) to second‐step sintering temperate (TH), and others used controlled cooling rates.

[15]. In the TSS method, the ceramic samples were heated to a higher temperature to achieve critical density and then immediately cooled to lower temperature and held at that tem‐ perature for long holding time to achieve full densification. Density of the sintered sample increased with increasing grain size during the first‐step sintering. However, grain growth is controlled at the second stage of sintering, and grain size versus density graph is horizontal at the final stage of the sintering [15]. Similar relationship was observed for other ceram‐

However, the most crucial task in this method is to identify suitable sintering parameters such as heating and cooling rate, the first‐ and second‐step sintering temperatures, holding

TSS with a high first‐step sintering temperature is widely used to obtain fully dense ceram‐ ics with controlled grain growth due to lower second‐step sintering temperature. The mech‐ anism for controlled grain growth in TSS with higher first‐step sintering temperature was

The general mechanisms that are responsible for densification during sintering are grain boundary migration and grain boundary diffusion. Grain boundary migration is responsible for the rapid grain growth in the final stage of conventional sintering. The densification of ceramics with grain growth suppression at the second‐step sintering can be explained by the absence of grain boundary migration during the second‐step sintering as described in **Figure 2**. In conventional single‐step sintering (SSS), the grain growth is accelerated due to grain boundary migration and grain boundary diffusion in the final stage of sintering. Rapid cooling before the second stage of sintering freezes the microstructure by immobilizing the tri‐ ple‐point junctions and continues densification will be achieved by grain boundary diffusion. For a successful TSS profile with higher first‐step sintering temperature, few conditions should be achieved at the end of first‐step sintering. The sample must be reaching a critical density at

) is assumed to be zero in few

Two-Step Sintering of Ceramics http://dx.doi.org/10.5772/68083 5

by Chen and Wang

ceramics, and it is widely

[17] and

O3

lead free antiferroelectric ceramics [18].

O3

.

[15], ZnO [16], Ni‐Cu‐Zn ferrite ceramics, BaTiO3

Sintering profile 1 also has few categories based on other sintering parameters:

): The first‐step sintering holding time (t<sup>1</sup>

TSS consists of heating the samples in two stages. Different sintering profiles were applied in TSS. Generalized diagrams of TSS are shown in **Figure 1(a)** and **(b)**. In the sintering profile 1, the first‐step sintering temperature is higher than the second‐step sintering temperature, and the first‐step holding time is lower than the second‐step holding time. Sintering profile 2 is the other way around. The first‐step sintering temperature is lower than the second‐step sintering temperature, and the first‐step holding time is normally higher than the second‐step holding time.

In addition, a modified two‐step sintering profile is also reported with a cooling step (to room temperature) in between first and second step sintering [2, 13, 14]. Especially, this method used different sintering methods at first and second steps [2, 14].

This chapter mainly focuses on TSS with high first‐step sintering temperature. The mecha‐ nism of densification with controlled grain growth is explained briefly, and the extended applications of TSS on different ceramics are outlined. Finally the chapter concludes the cur‐ rent trends and challenges of TSS as a fabrication method.

**Figure 1.** Different sintering profiles used in TSS (a) with high first‐step sintering temperature and (b) with a low first‐ step sintering temperature.

## **2. TSS with high first‐step sintering temperature**

TSS with a high first‐step sintering temperature is widely used to obtain fully dense ceram‐ ics with controlled grain growth due to lower second‐step sintering temperature. The sin‐ tering profile for TSS with higher first‐step sintering temperature is depicted in **Figure 1(a)**. Sintering profile 1 also has few categories based on other sintering parameters:


The successful TSS method using profile 1 was first introduced for Y<sup>2</sup> O3 by Chen and Wang [15]. In the TSS method, the ceramic samples were heated to a higher temperature to achieve critical density and then immediately cooled to lower temperature and held at that tem‐ perature for long holding time to achieve full densification. Density of the sintered sample increased with increasing grain size during the first‐step sintering. However, grain growth is controlled at the second stage of sintering, and grain size versus density graph is horizontal at the final stage of the sintering [15]. Similar relationship was observed for other ceram‐ ics such as Mg, Nb‐doped Y<sup>2</sup> O3 [15], ZnO [16], Ni‐Cu‐Zn ferrite ceramics, BaTiO3 [17] and 0.89Bi0.5Na0.5TiO3‐ 0.06BaTiO3‐ 0.05K0.5Na0.5NbO3 lead free antiferroelectric ceramics [18]. However, the most crucial task in this method is to identify suitable sintering parameters such as heating and cooling rate, the first‐ and second‐step sintering temperatures, holding times and sintering atmospheres.

#### **2.1. Sintering mechanism**

sintering techniques are limited by low mass productivity, and they are not economically feasible. Two‐stage sintering (TSS) is an effective way to achieve fine‐grained microstructured ceramics with high densification and relatively low cost. TSS method is successfully applied for all types of ceramics such as structural ceramics, bioceramics, ferrites, piezoelectric ceram‐ ics and electrolyte ceramics. Most of the ceramics exhibit controlled or no grain growth in the final stage of sintering and achieved near‐theoretical densities. The fine‐grained micro‐ structure enhances the mechanical, electrical, magnetic as well as piezoelectric properties of

TSS consists of heating the samples in two stages. Different sintering profiles were applied in TSS. Generalized diagrams of TSS are shown in **Figure 1(a)** and **(b)**. In the sintering profile 1, the first‐step sintering temperature is higher than the second‐step sintering temperature, and the first‐step holding time is lower than the second‐step holding time. Sintering profile 2 is the other way around. The first‐step sintering temperature is lower than the second‐step sintering temperature, and the first‐step holding time is normally higher than the second‐step

In addition, a modified two‐step sintering profile is also reported with a cooling step (to room temperature) in between first and second step sintering [2, 13, 14]. Especially, this method

This chapter mainly focuses on TSS with high first‐step sintering temperature. The mecha‐ nism of densification with controlled grain growth is explained briefly, and the extended applications of TSS on different ceramics are outlined. Finally the chapter concludes the cur‐

**Figure 1.** Different sintering profiles used in TSS (a) with high first‐step sintering temperature and (b) with a low first‐

ceramics which widen the applications of ceramics.

used different sintering methods at first and second steps [2, 14].

rent trends and challenges of TSS as a fabrication method.

holding time.

4 Sintering of Functional Materials

step sintering temperature.

TSS with a high first‐step sintering temperature is widely used to obtain fully dense ceram‐ ics with controlled grain growth due to lower second‐step sintering temperature. The mech‐ anism for controlled grain growth in TSS with higher first‐step sintering temperature was firstly proposed by Chen and Wang [15], for a TSS study on Y<sup>2</sup> O3 ceramics, and it is widely accepted and verified for other ceramics.

The general mechanisms that are responsible for densification during sintering are grain boundary migration and grain boundary diffusion. Grain boundary migration is responsible for the rapid grain growth in the final stage of conventional sintering. The densification of ceramics with grain growth suppression at the second‐step sintering can be explained by the absence of grain boundary migration during the second‐step sintering as described in **Figure 2**. In conventional single‐step sintering (SSS), the grain growth is accelerated due to grain boundary migration and grain boundary diffusion in the final stage of sintering. Rapid cooling before the second stage of sintering freezes the microstructure by immobilizing the tri‐ ple‐point junctions and continues densification will be achieved by grain boundary diffusion.

For a successful TSS profile with higher first‐step sintering temperature, few conditions should be achieved at the end of first‐step sintering. The sample must be reaching a critical density at

**Figure 2.** Schematic of densification of ceramics with grain growth during conventional single‐step sintering (SSS) and densification without grain growth at lower second‐step temperature during TSS with higher first‐step temperature.

**Figure 3.** Kinetic window for pure, Mg‐ and Nb‐doped Y<sup>2</sup>

magnified images of grain boundaries marked in circles at (a) [22].

size) [15].

O3

**Figure 4.** TEM images of the ZnO samples sintered (a) using TSS profile with the first‐step sintering temperature of 800°C for 60 s and a second‐step sintering temperature of 750°C for 72,000 s, (b) conventionally at 1100°C and (a1)–(a10)

(T is the first‐step sintering temperature and G is grain

Two-Step Sintering of Ceramics http://dx.doi.org/10.5772/68083 7

the end of the first‐step sintering so that the densification is active in the final sintering step. This critical density depends on the material. As reported, the critical density to be achieved at the end of the first‐step sintering for Y<sup>2</sup> O3 is 75% [15], BaTiO3 is 73% [17] and Ni‐Cu‐Zn ferrite ceramic is 76% [17]. The critical density is essential to ensure that the pores in the material are subcritical and unstable against shrinkage which can be filled by grain boundary diffusion in the second‐step sintering.

A kinetic window will be used to identify the optimum range of the first‐step sintering tem‐ perature for successful TSS. **Figure 3** illustrates the kinetic window for pure and Mg‐ and Nb‐doped Y<sup>2</sup> O3 . The filled squares between the lower and upper limits of the first‐step sinter‐ ing represent the successful second‐step sintering. The open squares above the upper limit of the first‐step sintering show grain growth during the second step and below the lower limit represents the failed attempts to achieve full densification. It can also be concluded that the kinetic window can be shifted up and down along the first‐step temperature axis with the addition of dopants. So far, the kinetic windows have been proposed to Y<sup>2</sup> O3 [15, 19], Mg, Nb‐doped Y<sup>2</sup> O3 [15, 19], ZnO [20], Ni‐Cu‐Zn ferrite [17], BaTiO3 [17] and (1−*x*) BiScO3−*<sup>x</sup>* PbTiO3 (BSPT) [21].

Another important study on exploring the mechanism of grain growth suppression in TSS with higher first‐step sintering temperature was reported on ZnO ceramics. The transmis‐ sion electron microscope (TEM) images of a TSS sample at a first‐step sintering temperature of 800°C for 60 s and a second‐step sintering temperature of 750°C for 20 h (**Figure 4(a)**) and a conventionally sintered ZnO at 1100°C (**Figure 4(b)**) were examined for the evidences of grain boundary migration. Ten different grain boundary zones were analysed from the TEM

**Figure 3.** Kinetic window for pure, Mg‐ and Nb‐doped Y<sup>2</sup> O3 (T is the first‐step sintering temperature and G is grain size) [15].

the end of the first‐step sintering so that the densification is active in the final sintering step. This critical density depends on the material. As reported, the critical density to be achieved at

**Figure 2.** Schematic of densification of ceramics with grain growth during conventional single‐step sintering (SSS) and densification without grain growth at lower second‐step temperature during TSS with higher first‐step temperature.

ceramic is 76% [17]. The critical density is essential to ensure that the pores in the material are subcritical and unstable against shrinkage which can be filled by grain boundary diffusion in

A kinetic window will be used to identify the optimum range of the first‐step sintering tem‐ perature for successful TSS. **Figure 3** illustrates the kinetic window for pure and Mg‐ and

ing represent the successful second‐step sintering. The open squares above the upper limit of the first‐step sintering show grain growth during the second step and below the lower limit represents the failed attempts to achieve full densification. It can also be concluded that the kinetic window can be shifted up and down along the first‐step temperature axis with the

Another important study on exploring the mechanism of grain growth suppression in TSS with higher first‐step sintering temperature was reported on ZnO ceramics. The transmis‐ sion electron microscope (TEM) images of a TSS sample at a first‐step sintering temperature of 800°C for 60 s and a second‐step sintering temperature of 750°C for 20 h (**Figure 4(a)**) and a conventionally sintered ZnO at 1100°C (**Figure 4(b)**) were examined for the evidences of grain boundary migration. Ten different grain boundary zones were analysed from the TEM

is 75% [15], BaTiO3

. The filled squares between the lower and upper limits of the first‐step sinter‐

is 73% [17] and Ni‐Cu‐Zn ferrite

O3

[17] and (1−*x*) BiScO3−*<sup>x</sup>*

[15, 19], Mg,

PbTiO3

O3

addition of dopants. So far, the kinetic windows have been proposed to Y<sup>2</sup>

[15, 19], ZnO [20], Ni‐Cu‐Zn ferrite [17], BaTiO3

the end of the first‐step sintering for Y<sup>2</sup>

the second‐step sintering.

6 Sintering of Functional Materials

O3

O3

Nb‐doped Y<sup>2</sup>

Nb‐doped Y<sup>2</sup>

(BSPT) [21].

**Figure 4.** TEM images of the ZnO samples sintered (a) using TSS profile with the first‐step sintering temperature of 800°C for 60 s and a second‐step sintering temperature of 750°C for 72,000 s, (b) conventionally at 1100°C and (a1)–(a10) magnified images of grain boundaries marked in circles at (a) [22].

image of TSSZnO sample. As can be seen in **Figure 4(a1–a10)**, junctions seem to have pinned the boundaries of growing grains, and the curvature of these boundaries resulted from the mentioned immobilized triple points. However, no similar zones are observed in the TEM images from conventionally sintered samples (**Figure 4(b)**).

Despite the fulfilment of zirconia in a wide variety of application, it suffers from low‐tem‐ perature degradation in the humid atmosphere around 65–500°C, with hinder biomedical application. Recently Sutharsini et al. [62] sintered fully dense (99% of TD) 3Y‐TZP via TSS with average grain size 290 nm. The sintered ceramics exhibited better hydrothermal ageing

Two-Step Sintering of Ceramics http://dx.doi.org/10.5772/68083 9

8YSZ ceramic is a promising candidate as a solid electrolyte for fuel cell application. Maca et al. [5] reported that efficiency of TSS depends on the crystal structure of the ceramics. The authors claimed that cubic structure has higher efficiency in controlling growth than hexagonal and tetragonal structures. Hence, TSS on 8YSZ controls the grain growth compared to 3Y‐TZP due to its cubic crystal structure. TSS has been widely applied to 8YSZ to improve mechani‐ cal, electrical as well as the gas permeance via controlling the grain growth [3, 9, 24, 32–34]. It enhance the hardness (∼13 GPa) [9, 32, 63], fracture toughness (3 MPam1/2) [9, 32, 63] and ionic conductivity (0.3 Scm−1) [9]. In addition to the above property, it is necessary to control grain growth to achieve optimum ratio of grain size to electrolyte layer thickness (0.1 < z < 0.3) to control gas permeance value for SOFC [64]. TSS effectively controls gas permeance value [26]. Hesabi et al. [9] reported that the TSS is efficient in obtaining fine grain 8YSZ when com‐ pared to conventional and microwave sintering with nearly full density. The grain size and conductivity of conventionally sintered sample at 1500°C with a heating rate of 5°C/min, microwave‐sintered samples at 1500°C with a heating rate of 50°C/min and TSS with the first‐step sintering temperature of 1250°C with no holding time and the second‐step sinter‐ ing temperature of 1050°C with 20 h holding time were compared. The grain sizes were reported as 2.15 μm, 0.9 μm and 295 nm, and the conductivity at 1000°C were 255.4, 322.6 and 398.6 mS/cm for conventional, microwave and two‐step sintered samples, respectively.

Alumina‐based ceramics are widely used for optical and biomedical applications. It is also used as filler for plastic and cutting inserts. However, the brittle nature of alumina limits the application. Ultrafine grain microstructure is crucial to enhance the mechanical properties such as hardness, wear resistance and strength. TSS successfully applied doped and pure alumina [5, 24, 25, 35–41]. Bodišová et al. [36] and others [5, 25, 35, 37, 39] successfully sintered fully dense alumina (98%) with controlled grain growth (grain growth below five times than that of powder size). Furthermore, the sample sintered with TSS exhibits excellent hardness

It is also reported that the room temperature cooling between the first and second steps of sin‐ tering also affects the grain growth suppression of corundum abrasives. The TSS with 1250°C first‐step temperature and 1350°C with a second‐step temperature of 1150°C for 5 h yielded fine‐ grained corundum abrasives with grain sizes of 65 and 80 nm, respectively. The samples sin‐ tered at the same sintering profiles with a room temperature cooling yielded grain sizes of 400 and 560 nm, respectively. The conventional sintering at 1300°C for 2 h resulted to a final grain

resistance against the conventionally sintered 3Y‐TZP ceramics.

*2.2.1.2. Sintering of 8YSZ*

*2.2.2. Sintering of aluminium‐based ceramics*

(18 GPa) and fracture toughness (4 MPam1/2) [35].

#### **2.2. Effect of TSS with high first‐step temperature on properties of ceramics**

By successfully controlling the grain growth, the TSS with higher first‐step sintering tem‐ perature is used to fabricate many ceramic materials with enhanced properties that are used for advanced applications [23]. So far TSS is applied for 3 mol% yttria‐stabilized tetragonal zirconia polycrystals (3Y‐TZP) [1, 4, 5, 8, 24–31], 8 mol% yttria‐stabilized zir‐ conia (8YSZ) [3, 5, 9, 24, 32–34], Al2 O3 [5, 24, 25, 35–43], doped alumina [37, 41], yttrium aluminium garnet (YAG) [44, 45], magnesium aluminium silicate [46], corundum [42, 43], hydroxyapatite [14, 47–50], forsterite (Mg2 SiO4 ) [51], TiO2 [11, 52–54], SrTiO3 [25] (K*<sup>x</sup>* Na1−*<sup>x</sup>* ) NbO3 (KNN) ceramic [12, 55–57], SiC [58–60] and Si3N4 [2] and used for several applica‐ tions. This section outlined the changes in properties of various ceramics sintered using different TSS profiles.

#### *2.2.1. Sintering of zirconia ceramics*

Pure zirconia has three crystallographic structures, monoclinic, tetragonal and cubic. At room temperature monoclinic is stable. In pure form, zirconia has very low appeal for use as engineering ceramic due to reverse transformation on cooling resulting in severe cracking associated with volume expansion (∼3%) to the monoclinic phase. In order to overcome this problem, stabilizers such as MgO, CaO and Y<sup>2</sup> O3 are added. Stabilizers reduce the change of chemical free energy and stabilize tetragonal or cubic phase at room temperature.

#### *2.2.1.1. 3Y‐TZP ceramics*

3Y‐TZP is also known as ceramic steel, which exhibits excellent mechanical properties. It is found in many applications, such as cutting tools, grinding media for powders, extrusion dies and biomedical application. Fully dense ceramics with uniform microstructure and fine grain size is essential to stabilize the tetragonal phase as well as to improve the mechanical properties such as hardness and toughness. There are many TSS profiles conducted to control the grain growth of Y‐TZP ceramics [1, 4, 5, 8, 24–31, 61]. Only few researchers successfully obtained fully dense ceramics with controlled grain growth (<5 times of powder size) [1, 4, 5, 25, 28, 29, 31, 62]. Among these studies, Binner et al. [31] and others [1, 29] successfully sintered nanostructured zirconia ceramics. Binner et al. [31] achieved fully dense (99% of TD) nanoceramics with the application of hybrid radiant and microwave sintering. However, the second‐step grain growth was not entirely suppressed in both radiant and microwave sintering, but the rapid growth observed in conventionally sintered samples was controlled. Application of TSS also enhances the hardness (12–14 GPa), fracture toughness 5–9 MPam1/2 and bending strength (900–1100) MPa [61, 62].

Despite the fulfilment of zirconia in a wide variety of application, it suffers from low‐tem‐ perature degradation in the humid atmosphere around 65–500°C, with hinder biomedical application. Recently Sutharsini et al. [62] sintered fully dense (99% of TD) 3Y‐TZP via TSS with average grain size 290 nm. The sintered ceramics exhibited better hydrothermal ageing resistance against the conventionally sintered 3Y‐TZP ceramics.

#### *2.2.1.2. Sintering of 8YSZ*

image of TSSZnO sample. As can be seen in **Figure 4(a1–a10)**, junctions seem to have pinned the boundaries of growing grains, and the curvature of these boundaries resulted from the mentioned immobilized triple points. However, no similar zones are observed in the TEM

By successfully controlling the grain growth, the TSS with higher first‐step sintering tem‐ perature is used to fabricate many ceramic materials with enhanced properties that are used for advanced applications [23]. So far TSS is applied for 3 mol% yttria‐stabilized tetragonal zirconia polycrystals (3Y‐TZP) [1, 4, 5, 8, 24–31], 8 mol% yttria‐stabilized zir‐

aluminium garnet (YAG) [44, 45], magnesium aluminium silicate [46], corundum [42, 43],

Pure zirconia has three crystallographic structures, monoclinic, tetragonal and cubic. At room temperature monoclinic is stable. In pure form, zirconia has very low appeal for use as engineering ceramic due to reverse transformation on cooling resulting in severe cracking associated with volume expansion (∼3%) to the monoclinic phase. In order to

reduce the change of chemical free energy and stabilize tetragonal or cubic phase at room

3Y‐TZP is also known as ceramic steel, which exhibits excellent mechanical properties. It is found in many applications, such as cutting tools, grinding media for powders, extrusion dies and biomedical application. Fully dense ceramics with uniform microstructure and fine grain size is essential to stabilize the tetragonal phase as well as to improve the mechanical properties such as hardness and toughness. There are many TSS profiles conducted to control the grain growth of Y‐TZP ceramics [1, 4, 5, 8, 24–31, 61]. Only few researchers successfully obtained fully dense ceramics with controlled grain growth (<5 times of powder size) [1, 4, 5, 25, 28, 29, 31, 62]. Among these studies, Binner et al. [31] and others [1, 29] successfully sintered nanostructured zirconia ceramics. Binner et al. [31] achieved fully dense (99% of TD) nanoceramics with the application of hybrid radiant and microwave sintering. However, the second‐step grain growth was not entirely suppressed in both radiant and microwave sintering, but the rapid growth observed in conventionally sintered samples was controlled. Application of TSS also enhances the hardness (12–14 GPa), fracture toughness 5–9 MPam1/2

 (KNN) ceramic [12, 55–57], SiC [58–60] and Si3N4 [2] and used for several applica‐ tions. This section outlined the changes in properties of various ceramics sintered using

) [51], TiO2

SiO4

[5, 24, 25, 35–43], doped alumina [37, 41], yttrium

[11, 52–54], SrTiO3

O3

[25] (K*<sup>x</sup>*

are added. Stabilizers

Na1−*<sup>x</sup>* )

images from conventionally sintered samples (**Figure 4(b)**).

conia (8YSZ) [3, 5, 9, 24, 32–34], Al2

*2.2.1. Sintering of zirconia ceramics*

NbO3

different TSS profiles.

8 Sintering of Functional Materials

temperature.

*2.2.1.1. 3Y‐TZP ceramics*

hydroxyapatite [14, 47–50], forsterite (Mg2

and bending strength (900–1100) MPa [61, 62].

**2.2. Effect of TSS with high first‐step temperature on properties of ceramics**

O3

overcome this problem, stabilizers such as MgO, CaO and Y<sup>2</sup>

8YSZ ceramic is a promising candidate as a solid electrolyte for fuel cell application. Maca et al. [5] reported that efficiency of TSS depends on the crystal structure of the ceramics. The authors claimed that cubic structure has higher efficiency in controlling growth than hexagonal and tetragonal structures. Hence, TSS on 8YSZ controls the grain growth compared to 3Y‐TZP due to its cubic crystal structure. TSS has been widely applied to 8YSZ to improve mechani‐ cal, electrical as well as the gas permeance via controlling the grain growth [3, 9, 24, 32–34]. It enhance the hardness (∼13 GPa) [9, 32, 63], fracture toughness (3 MPam1/2) [9, 32, 63] and ionic conductivity (0.3 Scm−1) [9]. In addition to the above property, it is necessary to control grain growth to achieve optimum ratio of grain size to electrolyte layer thickness (0.1 < z < 0.3) to control gas permeance value for SOFC [64]. TSS effectively controls gas permeance value [26].

Hesabi et al. [9] reported that the TSS is efficient in obtaining fine grain 8YSZ when com‐ pared to conventional and microwave sintering with nearly full density. The grain size and conductivity of conventionally sintered sample at 1500°C with a heating rate of 5°C/min, microwave‐sintered samples at 1500°C with a heating rate of 50°C/min and TSS with the first‐step sintering temperature of 1250°C with no holding time and the second‐step sinter‐ ing temperature of 1050°C with 20 h holding time were compared. The grain sizes were reported as 2.15 μm, 0.9 μm and 295 nm, and the conductivity at 1000°C were 255.4, 322.6 and 398.6 mS/cm for conventional, microwave and two‐step sintered samples, respectively.

#### *2.2.2. Sintering of aluminium‐based ceramics*

Alumina‐based ceramics are widely used for optical and biomedical applications. It is also used as filler for plastic and cutting inserts. However, the brittle nature of alumina limits the application. Ultrafine grain microstructure is crucial to enhance the mechanical properties such as hardness, wear resistance and strength. TSS successfully applied doped and pure alumina [5, 24, 25, 35–41]. Bodišová et al. [36] and others [5, 25, 35, 37, 39] successfully sintered fully dense alumina (98%) with controlled grain growth (grain growth below five times than that of powder size). Furthermore, the sample sintered with TSS exhibits excellent hardness (18 GPa) and fracture toughness (4 MPam1/2) [35].

It is also reported that the room temperature cooling between the first and second steps of sin‐ tering also affects the grain growth suppression of corundum abrasives. The TSS with 1250°C first‐step temperature and 1350°C with a second‐step temperature of 1150°C for 5 h yielded fine‐ grained corundum abrasives with grain sizes of 65 and 80 nm, respectively. The samples sin‐ tered at the same sintering profiles with a room temperature cooling yielded grain sizes of 400 and 560 nm, respectively. The conventional sintering at 1300°C for 2 h resulted to a final grain size of 800 nm [43]. In addition, the two‐step sintered corundum abrasive exhibited excellent hardness (22 GPa), fracture toughness (3 MPam1/2) and wear resistance (<2 × 10−7 mm3 /Nm) [43].

growth during the sintering in the presence of liquid phase is much more significant than that of solid‐state sintering. Therefore, it is practically impossible to obtain nanostructured ceramics by conventional single‐step liquid‐phase sintering. TSS was successfully applied in liquid‐phase‐sintered SiC ceramics, and a fully dense nanostructured SiC ceramic with a grain size of ~40 nm has been obtained [58, 59] in argon atmosphere. Magnani et al. [60, 70] also successfully sintered doped SiC via TSS with enhanced mechanical properties. The sin‐ tered samples exhibited excellent hardness (24 GPa), fracture toughness (3 MPam1/2), Young

Similar to SiC, silicon nitride, also non‐oxide ceramics, exhibits high hardness strength and thermal shock resistance. It is widely applied to automotive engine wear parts due to its outstanding mechanical properties and wear resistance. Bimodal microstructure of silicon

Bimodal microstructure enhances strength and toughness of the ceramic via crack bridging toughness mechanism [72]. Barium aluminosilicate‐reinforced silicon nitride sintered via TSS also exhibited higher flexural strength (565 MPa) and fracture toughness (7 MPam1/2). The obtained composite exhibits excellent mechanical properties compared to unreinforced bar‐

sintered with the first‐step sintering temperature 1300°C and the second‐step sintering tem‐ perature at 750°C for 15 h revealed high density (98.5%) with grain size 300 nm. Furthermore it exhibited fracture toughness of 3.61 MPam1/2. Compared with hydroxyapatite ceramics, for‐ sterite shows a significant improvement in the fracture toughness. Authors suggested that the two‐step sintering method can be used to fabricate improved forsterite dense ceramics with desired bioactivity and mechanical properties that might be suitable for hard tissue repair and

Environmental friendly lead‐free alkaline‐based niobate ceramics exhibited excellent piezoelec‐

successfully sintered by using TSS, and they exhibited excellent dielectric [12, 55, 73], ferroelec‐ tric [12, 55, 56, 73] and piezoelectric properties [12, 21, 55, 56]. Furthermore (K0.4425Na0.52Li0.0375)

Zinc oxide has been widely applied to electronic and optical devices. Furthermore, alumina‐ doped ZnO is used as an alternative to indium‐doped tin oxide (ITO) as a transparent con‐ ductive electrode in photovoltaic devices and displays. Electrical and optical properties of ZnO are mainly influenced by grain size. Grain growth of ZnO was successfully controlled using TSS [16, 20, 22, 74–78]. Zhang et al. [20] and others [16] successfully sintered fully dense

) ceramic is a new bioceramic with good biocompatibility. Forsterite

exhibited excellent temperature stability over a wide range of tempera‐

N4

ceramics. Alkaline‐based niobate (KNN) ceramics are

seed crystal [10, 71].

Two-Step Sintering of Ceramics http://dx.doi.org/10.5772/68083 11

modulus (400 Gpa) and flexural strength (500 MPa).

ium aluminosilicate matrix [10].

biomedical applications [51].

(Nb0.8925Sb0.07Ta0.0375)O3

*2.2.8. Sintering of zinc oxide*

tric properties compared to Pb(Zr,Ti)O3

SiO4

The forsterite (Mg2

nitride was successfully sintered by using TSS without using β‐Si<sup>3</sup>

*2.2.7. Sintering of alkaline niobate‐based lead‐free piezoelectric ceramics*

ture, which is attractive for piezoelectric applications [12].

#### *2.2.3. Sintering of YAG (Y3 Al5 O12)*

YAG is a familiar ceramic material for luminescent materials. Presently, single crystalline YAG is applied in lasers pumped by solid‐state LEDs, scintillators, and infrared windows. Generally, YAG is sintered for high temperature and long holding time which leads to abnor‐ mal grain growth. TSS is an efficient and economic method, which control the abnormal grain growth and improved the transmittance of YAG [44, 45]. YAG sintered with the first‐step sin‐ tering temperature at 1800°C and the second‐step sintering temperature 1550–1600°C revealed more than 40% transmission [44, 45]. Furthermore, neodymium‐doped YAG (Nd:YAG) sin‐ tered via TSS exhibited excellent transmittance (85%) [65].

#### *2.2.4. Sintering of hydroxyapatite (HA)*

Hydroxyapatite is a bioceramic that is used as tissue implants due to its excellent biocom‐ patibility. However, low toughness hinders application of artificial bone and teeth implants. Furthermore, the major drawback of HA is that it decomposed into secondary phases (α‐ or β‐tricalcium phosphate). In order to avoid such decomposition, TSS has been applied to HA [14, 47–51, 66, 67]. Feng et al. [14] and others [14, 49, 50, 66, 67] successfully sintered monophase HA without decomposition. Furthermore, TSS improved mechanical properties. Mazaheri et al. [50] achieved highest hardness (7.8 GPa) and fracture toughness (1.9 MPam1/2) via TSS.

#### *2.2.5. Sintering of Ni‐Cu‐Zn ferrite*

Ni‐Cu‐Zn ferrite ceramics received special attention due to its low cost, excellent heat and cor‐ rosion resistance, high magnetic permeability and low magnetic loss. It is used in many elec‐ tronic devices such as multilayer capacitor, sensors, antennas and broadband transformers. The electromagnetic properties of Ni‐Cu‐Zn ferrite are controlled by its microstructure and densification. Wang et al. [17] and Su et al. [68, 69] successfully sintered Ni‐Cu‐Zn ferrite by using TSS. Wang et al. [17] proposed kinetic window for successful TSS. Ni‐Cu‐Zn ferrite sin‐ tered by using TSS exhibited excellent magnetic properties [68, 69]. Magneto‐dielectric materi‐ als with matched permeability and permittivity are promising candidates as loading materials to reduce the physical dimensions of low‐frequency antennas. Ni‐Cu‐Zn ferrite sintered via TSS revealed almost equal permeability and permittivity of around 11.8. And the magnetic and dielectric loss tangents were lower than 0.015 in a frequency range from 10 to 100 MHz. These properties make the material useful to the design of miniaturized antennas [69].

#### *2.2.6. Sintering of Si‐based ceramics*

Silicon carbide is widely used for abrasives and refractories due to its high strength, hard‐ ness and excellent thermal shock resistance. In conventional single‐step sintering, abnor‐ mal grain growth is progressed due to its high sintering temperature. Generally, the grain growth during the sintering in the presence of liquid phase is much more significant than that of solid‐state sintering. Therefore, it is practically impossible to obtain nanostructured ceramics by conventional single‐step liquid‐phase sintering. TSS was successfully applied in liquid‐phase‐sintered SiC ceramics, and a fully dense nanostructured SiC ceramic with a grain size of ~40 nm has been obtained [58, 59] in argon atmosphere. Magnani et al. [60, 70] also successfully sintered doped SiC via TSS with enhanced mechanical properties. The sin‐ tered samples exhibited excellent hardness (24 GPa), fracture toughness (3 MPam1/2), Young modulus (400 Gpa) and flexural strength (500 MPa).

Similar to SiC, silicon nitride, also non‐oxide ceramics, exhibits high hardness strength and thermal shock resistance. It is widely applied to automotive engine wear parts due to its outstanding mechanical properties and wear resistance. Bimodal microstructure of silicon nitride was successfully sintered by using TSS without using β‐Si<sup>3</sup> N4 seed crystal [10, 71]. Bimodal microstructure enhances strength and toughness of the ceramic via crack bridging toughness mechanism [72]. Barium aluminosilicate‐reinforced silicon nitride sintered via TSS also exhibited higher flexural strength (565 MPa) and fracture toughness (7 MPam1/2). The obtained composite exhibits excellent mechanical properties compared to unreinforced bar‐ ium aluminosilicate matrix [10].

The forsterite (Mg2 SiO4 ) ceramic is a new bioceramic with good biocompatibility. Forsterite sintered with the first‐step sintering temperature 1300°C and the second‐step sintering tem‐ perature at 750°C for 15 h revealed high density (98.5%) with grain size 300 nm. Furthermore it exhibited fracture toughness of 3.61 MPam1/2. Compared with hydroxyapatite ceramics, for‐ sterite shows a significant improvement in the fracture toughness. Authors suggested that the two‐step sintering method can be used to fabricate improved forsterite dense ceramics with desired bioactivity and mechanical properties that might be suitable for hard tissue repair and biomedical applications [51].

#### *2.2.7. Sintering of alkaline niobate‐based lead‐free piezoelectric ceramics*

Environmental friendly lead‐free alkaline‐based niobate ceramics exhibited excellent piezoelec‐ tric properties compared to Pb(Zr,Ti)O3 ceramics. Alkaline‐based niobate (KNN) ceramics are successfully sintered by using TSS, and they exhibited excellent dielectric [12, 55, 73], ferroelec‐ tric [12, 55, 56, 73] and piezoelectric properties [12, 21, 55, 56]. Furthermore (K0.4425Na0.52Li0.0375) (Nb0.8925Sb0.07Ta0.0375)O3 exhibited excellent temperature stability over a wide range of tempera‐ ture, which is attractive for piezoelectric applications [12].

### *2.2.8. Sintering of zinc oxide*

size of 800 nm [43]. In addition, the two‐step sintered corundum abrasive exhibited excellent

YAG is a familiar ceramic material for luminescent materials. Presently, single crystalline YAG is applied in lasers pumped by solid‐state LEDs, scintillators, and infrared windows. Generally, YAG is sintered for high temperature and long holding time which leads to abnor‐ mal grain growth. TSS is an efficient and economic method, which control the abnormal grain growth and improved the transmittance of YAG [44, 45]. YAG sintered with the first‐step sin‐ tering temperature at 1800°C and the second‐step sintering temperature 1550–1600°C revealed more than 40% transmission [44, 45]. Furthermore, neodymium‐doped YAG (Nd:YAG) sin‐

Hydroxyapatite is a bioceramic that is used as tissue implants due to its excellent biocom‐ patibility. However, low toughness hinders application of artificial bone and teeth implants. Furthermore, the major drawback of HA is that it decomposed into secondary phases (α‐ or β‐tricalcium phosphate). In order to avoid such decomposition, TSS has been applied to HA [14, 47–51, 66, 67]. Feng et al. [14] and others [14, 49, 50, 66, 67] successfully sintered monophase HA without decomposition. Furthermore, TSS improved mechanical properties. Mazaheri et al. [50] achieved highest hardness (7.8 GPa) and fracture toughness (1.9 MPam1/2) via TSS.

Ni‐Cu‐Zn ferrite ceramics received special attention due to its low cost, excellent heat and cor‐ rosion resistance, high magnetic permeability and low magnetic loss. It is used in many elec‐ tronic devices such as multilayer capacitor, sensors, antennas and broadband transformers. The electromagnetic properties of Ni‐Cu‐Zn ferrite are controlled by its microstructure and densification. Wang et al. [17] and Su et al. [68, 69] successfully sintered Ni‐Cu‐Zn ferrite by using TSS. Wang et al. [17] proposed kinetic window for successful TSS. Ni‐Cu‐Zn ferrite sin‐ tered by using TSS exhibited excellent magnetic properties [68, 69]. Magneto‐dielectric materi‐ als with matched permeability and permittivity are promising candidates as loading materials to reduce the physical dimensions of low‐frequency antennas. Ni‐Cu‐Zn ferrite sintered via TSS revealed almost equal permeability and permittivity of around 11.8. And the magnetic and dielectric loss tangents were lower than 0.015 in a frequency range from 10 to 100 MHz.

These properties make the material useful to the design of miniaturized antennas [69].

Silicon carbide is widely used for abrasives and refractories due to its high strength, hard‐ ness and excellent thermal shock resistance. In conventional single‐step sintering, abnor‐ mal grain growth is progressed due to its high sintering temperature. Generally, the grain

/Nm) [43].

hardness (22 GPa), fracture toughness (3 MPam1/2) and wear resistance (<2 × 10−7 mm3

*2.2.3. Sintering of YAG (Y3*

10 Sintering of Functional Materials

*2.2.4. Sintering of hydroxyapatite (HA)*

*2.2.5. Sintering of Ni‐Cu‐Zn ferrite*

*2.2.6. Sintering of Si‐based ceramics*

*Al5 O12)*

tered via TSS exhibited excellent transmittance (85%) [65].

Zinc oxide has been widely applied to electronic and optical devices. Furthermore, alumina‐ doped ZnO is used as an alternative to indium‐doped tin oxide (ITO) as a transparent con‐ ductive electrode in photovoltaic devices and displays. Electrical and optical properties of ZnO are mainly influenced by grain size. Grain growth of ZnO was successfully controlled using TSS [16, 20, 22, 74–78]. Zhang et al. [20] and others [16] successfully sintered fully dense

ZnO without grain growth at the final stage of sintering. Furthermore Zhang et al. proposed kinetic window for successful TSS profile. Mazaheri et al. [22] confirmed the triple‐point drag mechanism for controlled grain growth at the second‐step sintering proposed by Chen and Wang [15] by using TEM image of two‐step sintered ZnO compacts. ZnO sintered via TSS exhibited excellent I–V characteristics [16, 74, 76]. ZnO varistors sintered via TSS exhibited higher breakdown field of 6–8 kVmm−1 and nonlinear coefficient of over 270 due to fine grain size and high concentration of ZnO‐ZnO grain contacts [16].

Al2 O3

ventionally sintered sample [40].

**4. Conclusion**

**Author details**

Ubenthiran Sutharsini1

Kuala Lumpur, Malaysia

1468‐6996/10/2/025004/meta

**References**

 ceramics are also known as a better translucent with gas‐impermeable properties which is suitable for high‐pressure lamp envelopes. The high optical transmittance requires spe‐ cial efforts to eliminate any light scattering centres such as residual pores, grain boundaries, secondary phases and rough surfaces in the material. MgO‐doped alumina ceramic sintered by TSS with lower first‐step sintering exhibited improved the transmittance compared to con‐

Despite long holding times, TSS with higher first‐step sintering temperature is convenient to achieve fully dense and fine‐grained microstructured ceramics with improved properties. The TSS is successfully applied to a range of ceramic materials, and their application is broad‐ ened. TSS also helped the emergence of sintering as a fabrication technique. Tailoring the TSS conditions and theoretical studies on TSS mechanisms will make TSS a cost‐effective method to fabricate advanced ceramics. TSS can also be studied by using different sintering methods

and Ramesh Singh2

ceramics by colloidal processing.

Two-Step Sintering of Ceramics http://dx.doi.org/10.5772/68083 13

and sintering environments in the first and second steps of sintering.

\*Address all correspondence to: ubsutharsini@gmail.com

powders on the sintering of nanostructured ZrO2

2002;**85**:245‐252. DOI: 10.1111/j.1151‐2916.2002.tb00073.x

Society. 2013;**33**:637‐641. DOI: 10.1016/j.jeurceramsoc.2012.10.002

\*, Murugathas Thanihaichelvan1

1 Department of Physics, Faculty of Science, University of Jaffna, Jaffna, Sri Lanka

2 Department of Mechanical Engineering, Faculty of Engineering, University of Malaya,

[1] Suárez G, Sakka Y, Suzuki TS, Uchikoshi T, Zhu X, Aglietti EF. Effect of starting

[2] Kim HD, Han BD, Park DS, Lee BT, Becher PF. Novel two‐step sintering process to obtain a bimodal microstructure in silicon nitride. Journal of the American Ceramic Society.

[3] Schwarz S, Guillon O. Two step sintering of cubic yttria stabilized zirconia using Field Assisted Sintering Technique/Spark Plasma Sintering. Journal of the European Ceramic

Science and Technology of Advanced Materials. 2009;**10**:025004(1‐8). DOI: 10.1088/

## *2.2.9. Sintering of BaTiO3*

Barium titanate (BaTiO3 ) is a polycrystalline piezoelectric ceramic. It is widely used to piezo‐ electric transducers, sensors and actuators. Many TSS studies have been conducted on BaTiO3 [17, 79–90]. Barium titanate ceramic is widely applied to multilayered ceramic capacitors (MLCC), transducers and pyroelectric detectors due to its dielectric, ferroelectric and piezo‐ electric properties. Wang et al. [17, 83] successfully sintered fully dense nanostructured ceramic and proposed kinetic window for successful TSS. TSS not only improved the densifi‐ cation and grain growth but also enhanced dielectric and piezoelectric properties [79, 81, 82, 85–88, 91, 92]. TSS samples exhibited excellent piezoelectric constant 519 pN/C and relative permittivity of 6079 [82, 92]. Tian et al. [86] reported that TSS revealed excellent dielectric con‐ stant of 2400 at room temperature, low dielectric loss (<1%) and high insulation resistivity of 1012 Ωcm, which could be beneficial for multilayer capacitor application.

## **3. TSS with low first‐step sintering temperature**

In few TSS studies, samples were initially heated to lower temperatures and then to higher temperature as shown in **Figure 1(b)**. Here the first step is normally a pre‐coarsening step that is performed for several purposes including removal of volatile fraction and smoothing the pore channels. This method was successful to prepare fully dense nano‐sized pure ZrO2 , fully dense (>98%) alumina [38, 40] and alumina‐doped zirconia (7.5Al2 O3 .92.5ZrO2 ) (vol.%) [8].

Sintering of pure zirconia has major drawback due to its reversible tetragonal‐to‐monoclinic phase transformation associated with shape deformation. Tartaj and Tartaj [93] applied TSS for pure zirconia below the phase transition temperature (<1150°C). In the first step, the com‐ pact allowed to achieve 96% of the density at 950°C for 10 h, and then the second‐stage sinter‐ ing temperature increased to 1050°C and achieved fully dense crack‐free pure zirconia with grain size less than 200 nm.

Fully dense (99%) alumina‐doped zirconia (7.5Al<sup>2</sup> O3 –92.5ZrO2 ) (vol.%) is also successfully sin‐ tered by using microwave TSS by using lower first‐step sintering temperature. Furthermore, microwave‐assisted TSS revealed higher density (99%), hardness (13 GPa), fracture tough‐ ness (12 MPam1/2) and bending strength (750 MPa) than the conventional single‐step sintering. Alumina‐toughened zirconia is widely applied to dental implant due to its excellent biocom‐ patibility and hardness [8].

Al2 O3 ceramics are also known as a better translucent with gas‐impermeable properties which is suitable for high‐pressure lamp envelopes. The high optical transmittance requires spe‐ cial efforts to eliminate any light scattering centres such as residual pores, grain boundaries, secondary phases and rough surfaces in the material. MgO‐doped alumina ceramic sintered by TSS with lower first‐step sintering exhibited improved the transmittance compared to con‐ ventionally sintered sample [40].

## **4. Conclusion**

ZnO without grain growth at the final stage of sintering. Furthermore Zhang et al. proposed kinetic window for successful TSS profile. Mazaheri et al. [22] confirmed the triple‐point drag mechanism for controlled grain growth at the second‐step sintering proposed by Chen and Wang [15] by using TEM image of two‐step sintered ZnO compacts. ZnO sintered via TSS exhibited excellent I–V characteristics [16, 74, 76]. ZnO varistors sintered via TSS exhibited higher breakdown field of 6–8 kVmm−1 and nonlinear coefficient of over 270 due to fine grain

electric transducers, sensors and actuators. Many TSS studies have been conducted on BaTiO3 [17, 79–90]. Barium titanate ceramic is widely applied to multilayered ceramic capacitors (MLCC), transducers and pyroelectric detectors due to its dielectric, ferroelectric and piezo‐ electric properties. Wang et al. [17, 83] successfully sintered fully dense nanostructured ceramic and proposed kinetic window for successful TSS. TSS not only improved the densifi‐ cation and grain growth but also enhanced dielectric and piezoelectric properties [79, 81, 82, 85–88, 91, 92]. TSS samples exhibited excellent piezoelectric constant 519 pN/C and relative permittivity of 6079 [82, 92]. Tian et al. [86] reported that TSS revealed excellent dielectric con‐ stant of 2400 at room temperature, low dielectric loss (<1%) and high insulation resistivity of

In few TSS studies, samples were initially heated to lower temperatures and then to higher temperature as shown in **Figure 1(b)**. Here the first step is normally a pre‐coarsening step that is performed for several purposes including removal of volatile fraction and smoothing the pore channels. This method was successful to prepare fully dense nano‐sized pure ZrO2

Sintering of pure zirconia has major drawback due to its reversible tetragonal‐to‐monoclinic phase transformation associated with shape deformation. Tartaj and Tartaj [93] applied TSS for pure zirconia below the phase transition temperature (<1150°C). In the first step, the com‐ pact allowed to achieve 96% of the density at 950°C for 10 h, and then the second‐stage sinter‐ ing temperature increased to 1050°C and achieved fully dense crack‐free pure zirconia with

O3

tered by using microwave TSS by using lower first‐step sintering temperature. Furthermore, microwave‐assisted TSS revealed higher density (99%), hardness (13 GPa), fracture tough‐ ness (12 MPam1/2) and bending strength (750 MPa) than the conventional single‐step sintering. Alumina‐toughened zirconia is widely applied to dental implant due to its excellent biocom‐

–92.5ZrO2

) is a polycrystalline piezoelectric ceramic. It is widely used to piezo‐

O3

.92.5ZrO2

) (vol.%) is also successfully sin‐

size and high concentration of ZnO‐ZnO grain contacts [16].

1012 Ωcm, which could be beneficial for multilayer capacitor application.

**3. TSS with low first‐step sintering temperature**

dense (>98%) alumina [38, 40] and alumina‐doped zirconia (7.5Al2

*2.2.9. Sintering of BaTiO3*

12 Sintering of Functional Materials

Barium titanate (BaTiO3

grain size less than 200 nm.

patibility and hardness [8].

Fully dense (99%) alumina‐doped zirconia (7.5Al<sup>2</sup>

Despite long holding times, TSS with higher first‐step sintering temperature is convenient to achieve fully dense and fine‐grained microstructured ceramics with improved properties. The TSS is successfully applied to a range of ceramic materials, and their application is broad‐ ened. TSS also helped the emergence of sintering as a fabrication technique. Tailoring the TSS conditions and theoretical studies on TSS mechanisms will make TSS a cost‐effective method to fabricate advanced ceramics. TSS can also be studied by using different sintering methods and sintering environments in the first and second steps of sintering.

## **Author details**

Ubenthiran Sutharsini1 \*, Murugathas Thanihaichelvan1 and Ramesh Singh2

\*Address all correspondence to: ubsutharsini@gmail.com

1 Department of Physics, Faculty of Science, University of Jaffna, Jaffna, Sri Lanka

2 Department of Mechanical Engineering, Faculty of Engineering, University of Malaya, Kuala Lumpur, Malaysia

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**Chapter 2**

**Provisional chapter**

**Development of Metal Matrix Composites Using**

**Development of Metal Matrix Composites Using** 

DOI: 10.5772/68081

In this book chapter, aluminum (Al)-based metal matrix composites (AMMCs) with

microwave sintering and subsequent hot extrusion processes. The role of various nano/ micro-sized reinforcements in altering the structural, mechanical, and thermal properties of the microwave-extruded composites was systematically studied. The X-ray diffraction (XRD) patterns indicated that the main components were Al, SiC, Si<sup>3</sup>

, and Al-Al2

electron microscopy (SEM) and energy dispersive spectroscopy (EDS) elemental mapping confirm the homogeneous distribution of reinforcing particles in the Al matrix.

rior hardness, ultimate compression/tensile strength, and Young's modulus, while having a lower coefficient of thermal expansion compared to other studied Al composites. Findings presented are expected to pave the way to design, develop, and synthesize other aluminum-based metal matrix composites for automotive and industrial applications.

**Keywords:** Al matrix composites, ceramic reinforcements, microwave sintering, hot extrusion, mechanical properties, thermal properties, fracture behavior

O3

N4

N4

, and Al2

O3

N4 metal matrix composite exhibited supe-

composites, respectively. Scanning

, were produced by

N4 , and

© 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.

Microwaves occupy the portion of electromagnetic radiation spectrum between 300 MHz and 300 GHz with wavelengths ranging from 1 mm to 1 m in free space. Although the frequencies available for processing of materials are 24.124 GHz, 5.8 GHz, 2.45 GHz, and 915 MHz, generally it is carried out at 915 MHz and 2.45 GHz. Usually, 9.15 and 2.45 GHz are commonly

**Microwave Sintering Technique**

Penchal Reddy Matli, Rana Abdul Shakoor

**Microwave Sintering Technique**

Penchal Reddy Matli, Rana Abdul Shakoor and

Additional information is available at the end of the chapter

various reinforcing ceramic particles, such as SiC, Si3

O3 for the studied Al-SiC, Al-Si3

Mechanistic studies revealed that the Al-Si3

Additional information is available at the end of the chapter

Adel Mohamed Amer Mohamed

and Adel Mohamed Amer Mohamed

http://dx.doi.org/ 10.5772/68081

**Abstract**

Al2

**1. Introduction**

**Provisional chapter**

## **Development of Metal Matrix Composites Using Microwave Sintering Technique Microwave Sintering Technique**

**Development of Metal Matrix Composites Using** 

DOI: 10.5772/68081

Penchal Reddy Matli, Rana Abdul Shakoor and Adel Mohamed Amer Mohamed and Adel Mohamed Amer Mohamed Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

Penchal Reddy Matli, Rana Abdul Shakoor

http://dx.doi.org/ 10.5772/68081

#### **Abstract**

In this book chapter, aluminum (Al)-based metal matrix composites (AMMCs) with various reinforcing ceramic particles, such as SiC, Si3 N4 , and Al2 O3 , were produced by microwave sintering and subsequent hot extrusion processes. The role of various nano/ micro-sized reinforcements in altering the structural, mechanical, and thermal properties of the microwave-extruded composites was systematically studied. The X-ray diffraction (XRD) patterns indicated that the main components were Al, SiC, Si<sup>3</sup> N4 , and Al2 O3 for the studied Al-SiC, Al-Si3 N4 , and Al-Al2 O3 composites, respectively. Scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDS) elemental mapping confirm the homogeneous distribution of reinforcing particles in the Al matrix. Mechanistic studies revealed that the Al-Si3 N4 metal matrix composite exhibited superior hardness, ultimate compression/tensile strength, and Young's modulus, while having a lower coefficient of thermal expansion compared to other studied Al composites. Findings presented are expected to pave the way to design, develop, and synthesize other aluminum-based metal matrix composites for automotive and industrial applications.

**Keywords:** Al matrix composites, ceramic reinforcements, microwave sintering, hot extrusion, mechanical properties, thermal properties, fracture behavior

### **1. Introduction**

Microwaves occupy the portion of electromagnetic radiation spectrum between 300 MHz and 300 GHz with wavelengths ranging from 1 mm to 1 m in free space. Although the frequencies available for processing of materials are 24.124 GHz, 5.8 GHz, 2.45 GHz, and 915 MHz, generally it is carried out at 915 MHz and 2.45 GHz. Usually, 9.15 and 2.45 GHz are commonly

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. © 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.

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons

used in industrial equipment [1, 2]. Microwave cavities are of two types: single-mode resonance cavities and multimode resonance cavities. Single-mode cavities are specially designed and generally used for industrial applications. The domestic microwave ovens are multimode cavities in which multiple plane waves impinge on the load (material to be heated) from different directions. The microwave radiation and heating were applied to the fabrication of ceramic materials till the 1990s [3–5].

that cause the adherence of loose particles to each other [13]. Sintering of materials using microwaves is a newly explored method, and it has been applied successfully in processing of various materials. For sintering using microwaves, the electromagnetic contrivance of the microwaves interact directly with the materials and, magnetic and dielectric losses lead to self-heating of materials. Initially, microwave energy was used to sinter various types of ceramic materials [9, 15]. By using microwave sintering, equal or superior performance ceramic products [16–18] can be produced through shorter sintering time at a lower temperature and at low cost when compared to the conventional sintered products. At room temperature, most of the ceramics do not couple well with microwaves particularly at 2.45 GHz microwave frequency and they are not heated appreciably [19]. Their coupling efficiency can be increased by increasing the temperature, using microwave susceptor/absorber, varying their morphology, and changing the microwave frequency. In most of the hybrid sintering methods, SiC of various forms such as SiC rods [4] with SiC sample holders [20] are preferably used mainly due to its high loss factor. Most of the investigations [21–25] were conducted on the microwave sintering of semiconductors, inorganic, ceramics, and polymeric materials until 2000. Lack of research on microwave sintering of metals is based on the well-known fact that all metals reflect the microwaves causing arching during microwave heating and thus limited diffusion of the microwave waves. Later, researchers grabbed that arching phenomenon applies only for metal-based composites in the form of powder compacts [14, 26]. The idea of applying microwave energy to sinter metals and metallic materials is relatively new and limited studies on sintering of metal-based materials are available in the literature. From the literature, most of the studies are reported on the microwave sintering of iron-based materials and only a few reports are available on

Development of Metal Matrix Composites Using Microwave Sintering Technique

http://dx.doi.org/ 10.5772/68081

25

aluminum-, magnesium-, and copper-based materials/composites [14, 26–29].

Development of metal matrix composites (MMCs) has been an important innovation in materials engineering over the past three decades. Metal matrix composites offer several attractive advantages over traditional engineering materials due to their superior properties [30, 31]. Metal matrix composites can be divided into three broad categories (**Figure 2**): (i) continuous fiber-reinforced matrix composites, (ii) small fiber-reinforced matrix composites, and (iii) particulate-rein-

Long Fibre Composite Short Fibre Composite Particulate Composite

**2. Metal matrix composites**

**Figure 2.** Types of reinforcement in a composite.

#### **1.1. Characteristics and merits of microwave sintering process**

Microwave sintering has attained global acceptance because of superior benefits over the traditional sintering techniques. The characteristic (**Figure 1**) of microwave heating is basically different from conventional heating [6]. For conventional heating, an external heating element is used for heat generation and then it transferred to the test materials through convention, conduction, and radiation. In microwave heating, the heat is generated internally within the test sample by rapid oscillation of dipoles at microwave frequencies [7], instead of diffusion from external sources and hence the heating is from the core of the sample to outside. Heating is very rapid and volumetric due to energy conversion rather than energy transfer as in conventional heating.

Microwave heating has many advantages over conventional heating, including cost and energy savings, and considerable reduction in processing time [8]. By using microwave energy as a heating source, short sintering time at desired temperature offers an opportunity to control, especially the microstructure coarsening during sintering, leading to excellent mechanical properties [9]. Instead of using only microwaves as a heating source, microwave heating system through a combination of conventional conduction heating and energy conversion heating using a microwave is found to be more advantageous for heating or sintering of materials [10–12]. The advantages of microwave rapid sintering include rapid and more uniform heating, prevention of hot spot formation, and more uniform and finer microstructure leading to high-performance products [13, 14].

#### **1.2. Microwave sintering of materials**

In principle, sintering is one of the consolidation methods to make bulk objects from loose powder compacts by heating the material below its melting point. Conventionally, the green body (unsintered powder compact) is sintered using resistant heating. Since the resistive heating is the application of thermal energy, sintering process depends on the diffusion of atoms

**Figure 1.** Heat distribution within a material during conventional and microwave heating.

that cause the adherence of loose particles to each other [13]. Sintering of materials using microwaves is a newly explored method, and it has been applied successfully in processing of various materials. For sintering using microwaves, the electromagnetic contrivance of the microwaves interact directly with the materials and, magnetic and dielectric losses lead to self-heating of materials. Initially, microwave energy was used to sinter various types of ceramic materials [9, 15]. By using microwave sintering, equal or superior performance ceramic products [16–18] can be produced through shorter sintering time at a lower temperature and at low cost when compared to the conventional sintered products. At room temperature, most of the ceramics do not couple well with microwaves particularly at 2.45 GHz microwave frequency and they are not heated appreciably [19]. Their coupling efficiency can be increased by increasing the temperature, using microwave susceptor/absorber, varying their morphology, and changing the microwave frequency. In most of the hybrid sintering methods, SiC of various forms such as SiC rods [4] with SiC sample holders [20] are preferably used mainly due to its high loss factor.

Most of the investigations [21–25] were conducted on the microwave sintering of semiconductors, inorganic, ceramics, and polymeric materials until 2000. Lack of research on microwave sintering of metals is based on the well-known fact that all metals reflect the microwaves causing arching during microwave heating and thus limited diffusion of the microwave waves. Later, researchers grabbed that arching phenomenon applies only for metal-based composites in the form of powder compacts [14, 26]. The idea of applying microwave energy to sinter metals and metallic materials is relatively new and limited studies on sintering of metal-based materials are available in the literature. From the literature, most of the studies are reported on the microwave sintering of iron-based materials and only a few reports are available on aluminum-, magnesium-, and copper-based materials/composites [14, 26–29].

## **2. Metal matrix composites**

used in industrial equipment [1, 2]. Microwave cavities are of two types: single-mode resonance cavities and multimode resonance cavities. Single-mode cavities are specially designed and generally used for industrial applications. The domestic microwave ovens are multimode cavities in which multiple plane waves impinge on the load (material to be heated) from different directions. The microwave radiation and heating were applied to the fabrication of

Microwave sintering has attained global acceptance because of superior benefits over the traditional sintering techniques. The characteristic (**Figure 1**) of microwave heating is basically different from conventional heating [6]. For conventional heating, an external heating element is used for heat generation and then it transferred to the test materials through convention, conduction, and radiation. In microwave heating, the heat is generated internally within the test sample by rapid oscillation of dipoles at microwave frequencies [7], instead of diffusion from external sources and hence the heating is from the core of the sample to outside. Heating is very rapid and volumetric due to energy conversion rather than energy transfer as in conventional heating. Microwave heating has many advantages over conventional heating, including cost and energy savings, and considerable reduction in processing time [8]. By using microwave energy as a heating source, short sintering time at desired temperature offers an opportunity to control, especially the microstructure coarsening during sintering, leading to excellent mechanical properties [9]. Instead of using only microwaves as a heating source, microwave heating system through a combination of conventional conduction heating and energy conversion heating using a microwave is found to be more advantageous for heating or sintering of materials [10–12]. The advantages of microwave rapid sintering include rapid and more uniform heating, prevention of hot spot formation, and more uniform and finer microstruc-

In principle, sintering is one of the consolidation methods to make bulk objects from loose powder compacts by heating the material below its melting point. Conventionally, the green body (unsintered powder compact) is sintered using resistant heating. Since the resistive heating is the application of thermal energy, sintering process depends on the diffusion of atoms

**Figure 1.** Heat distribution within a material during conventional and microwave heating.

ceramic materials till the 1990s [3–5].

24 Sintering of Functional Materials

**1.1. Characteristics and merits of microwave sintering process**

ture leading to high-performance products [13, 14].

**1.2. Microwave sintering of materials**

Development of metal matrix composites (MMCs) has been an important innovation in materials engineering over the past three decades. Metal matrix composites offer several attractive advantages over traditional engineering materials due to their superior properties [30, 31]. Metal matrix composites can be divided into three broad categories (**Figure 2**): (i) continuous fiber-reinforced matrix composites, (ii) small fiber-reinforced matrix composites, and (iii) particulate-rein-

Long Fibre Composite Short Fibre Composite Particulate Composite

**Figure 2.** Types of reinforcement in a composite.

forced matrix composites. Among all these, particulate-reinforced metal matrix composites have gained decent interest because of their superior properties and low manufacturing expenditure. In light metal matrix composites (MMCs) [32], Al or Mg is mostly used as a base metal matrix, and ceramic particles (carbides, nitrides, and oxides) are generally used as reinforcement phase.

Nowadays, the demand of reducing energy consumption, especially in automotive industries, is becoming a critical issue. Development of lightweight aluminum-based composites can be considered as one of the promising solutions to address this issue.

The nano-sized reinforcements have a major role in improving the physical and mechanical properties which can be achieved by the addition of small volume fractions (≤2%), whereas for micron-sized particle-reinforced metal matrix composites higher volume fractions (≫10%) are essential [27]. Further addition of reinforcement will cause degradation of composite properties, which can be attributed to the possible agglomeration, clustering of reinforcement, and micro-porosity in the nanocomposites. Recently, there has been considerable interest in the production of metal matrix nanocomposites in which nanoparticulates are incorporated into the base matrix [33]. The production of nanocomposites is currently under exploration and is still at its laboratory scale research level. However, interestingly, when compared to composites with micron-sized reinforcements, nanocomposites exhibit comparable or better mechanical properties with the use of lesser amount of reinforcements [9–12]. Both casting and powder metallurgy (PM) methods can be used to fabricate metal matrix nanocomposites. Historically, PM methods have been developed successfully and commercially used by different manufactures and have also been applied in the production of MMCs for aerospace applications. As compared to casting methods, PM approach has shown its advantage to produce uniform microstructures leading to develop high-performance composite materials [34].

The blending of the mixture was carried out at room temperature using a Retsch PM400 planetary ball mill for 2 h with the milling speed of 200 rpm in order to get a homogeneous particle distribution. No balls were used in this stage. The ball milled powders were cold compacted using a uniaxial pressure of 50 tons into billets (40 mm length with 35 mm diameter). The sintering of the compacted cylindrical billets was carried out at 550°C using an innovative microwave sintering process [35], just below the melting temperature of Al. The other metal matrix

Prior to hot extrusion, the microwave sintered billets were soaked at 400°C for 1 h and then hot extruded at 350°C and 500 MPa. The extrusion ratio was ~20.25:1 to produce an 8 mm diameter extruded rod, as can be seen in **Figure 4(a)**. After extrusion, these rods were subse-

**Pure Al Al-SiC Al-Si3N4 Al-Al2O3**

O3

Development of Metal Matrix Composites Using Microwave Sintering Technique

http://dx.doi.org/ 10.5772/68081

27

) were also prepared in a similar manner.

and Al-15 vol.% Al<sup>2</sup>

composites (Al-1.5 vol.% Si<sup>3</sup>

quently used for characterization studies.

**(a)**

**(b)**

**Figure 4.** The pictures of the produced AMMCs.

**Figure 3.** Schematic flow chart of the experimental design.

N4

At present, the development of metal matrix composites with light metal matrices are gaining increasing attention due to their enhanced properties coupled with weight savings. These unique properties make them attractive for automotive and aircraft industries in which the weight reduction is the critical factor. So far, extensive studies have been done for the production of aluminum matrix composites and now these are being manufactured commercially for numerous industrial applications. The development of economical aluminum nanocomposites using cost-effective fabrication techniques will serve the requirement of the development of light-weight structural materials well suited to industrial and commercial applications. The main objective of our study was to fabricate high-performance aluminum metal matrix composites through cost-effective processing technique based on PM route incorporating microwave sintering method. A comparison of the microstructural, mechanical, and thermal properties is presented to elucidate the usefulness of the manufactured composite materials.

## **3. Fabrication of Al metal matrix composites**

In the current book chapter, the presented composite materials were synthesized through the powder metallurgy method of mixing the matrix (pure aluminum) and reinforcements (SiC, Si3 N4 , and Al2 O3 ). **Figure 3** presents the schematic flow chart of the experimental design. To produce Al-SiC nanocomposites, nano-sized SiC powder (1.5 vol.%) was added to pure Al.

**Figure 3.** Schematic flow chart of the experimental design.

forced matrix composites. Among all these, particulate-reinforced metal matrix composites have gained decent interest because of their superior properties and low manufacturing expenditure. In light metal matrix composites (MMCs) [32], Al or Mg is mostly used as a base metal matrix, and ceramic particles (carbides, nitrides, and oxides) are generally used as reinforcement phase. Nowadays, the demand of reducing energy consumption, especially in automotive industries, is becoming a critical issue. Development of lightweight aluminum-based composites

The nano-sized reinforcements have a major role in improving the physical and mechanical properties which can be achieved by the addition of small volume fractions (≤2%), whereas for micron-sized particle-reinforced metal matrix composites higher volume fractions (≫10%) are essential [27]. Further addition of reinforcement will cause degradation of composite properties, which can be attributed to the possible agglomeration, clustering of reinforcement, and micro-porosity in the nanocomposites. Recently, there has been considerable interest in the production of metal matrix nanocomposites in which nanoparticulates are incorporated into the base matrix [33]. The production of nanocomposites is currently under exploration and is still at its laboratory scale research level. However, interestingly, when compared to composites with micron-sized reinforcements, nanocomposites exhibit comparable or better mechanical properties with the use of lesser amount of reinforcements [9–12]. Both casting and powder metallurgy (PM) methods can be used to fabricate metal matrix nanocomposites. Historically, PM methods have been developed successfully and commercially used by different manufactures and have also been applied in the production of MMCs for aerospace applications. As compared to casting methods, PM approach has shown its advantage to produce uniform microstructures leading to develop high-performance composite materials [34]. At present, the development of metal matrix composites with light metal matrices are gaining increasing attention due to their enhanced properties coupled with weight savings. These unique properties make them attractive for automotive and aircraft industries in which the weight reduction is the critical factor. So far, extensive studies have been done for the production of aluminum matrix composites and now these are being manufactured commercially for numerous industrial applications. The development of economical aluminum nanocomposites using cost-effective fabrication techniques will serve the requirement of the development of light-weight structural materials well suited to industrial and commercial applications. The main objective of our study was to fabricate high-performance aluminum metal matrix composites through cost-effective processing technique based on PM route incorporating microwave sintering method. A comparison of the microstructural, mechanical, and thermal properties is presented to elucidate the usefulness of the manufactured composite materials.

In the current book chapter, the presented composite materials were synthesized through the powder metallurgy method of mixing the matrix (pure aluminum) and reinforcements (SiC,

produce Al-SiC nanocomposites, nano-sized SiC powder (1.5 vol.%) was added to pure Al.

). **Figure 3** presents the schematic flow chart of the experimental design. To

can be considered as one of the promising solutions to address this issue.

**3. Fabrication of Al metal matrix composites**

Si3 N4

, and Al2

26 Sintering of Functional Materials

O3

The blending of the mixture was carried out at room temperature using a Retsch PM400 planetary ball mill for 2 h with the milling speed of 200 rpm in order to get a homogeneous particle distribution. No balls were used in this stage. The ball milled powders were cold compacted using a uniaxial pressure of 50 tons into billets (40 mm length with 35 mm diameter). The sintering of the compacted cylindrical billets was carried out at 550°C using an innovative microwave sintering process [35], just below the melting temperature of Al. The other metal matrix composites (Al-1.5 vol.% Si<sup>3</sup> N4 and Al-15 vol.% Al<sup>2</sup> O3 ) were also prepared in a similar manner.

Prior to hot extrusion, the microwave sintered billets were soaked at 400°C for 1 h and then hot extruded at 350°C and 500 MPa. The extrusion ratio was ~20.25:1 to produce an 8 mm diameter extruded rod, as can be seen in **Figure 4(a)**. After extrusion, these rods were subsequently used for characterization studies.

**Figure 4.** The pictures of the produced AMMCs.

The phase identification of the extruded samples was carried out using X-ray powder diffractometer (PANalytical X'pert Pro) based on Cu-Kα radiation (1.541 Å) in the 2*θ* range of 30–80° at scan rate of 0.2°/min. Individual phases were identified by matching the typical X-ray diffraction (XRD) peaks against JCPDS data. Field emission scanning electron microscopy (SEM) (JEOL JSM-6010 and Hitachi FESEM-S4300) with energy dispersion spectroscopy (EDS) was used to identify the reinforcement phase and microstructure of the extruded composite samples.

The hardness testing of the pure Al and composite samples was carried out using Vicker's hardness tester with applied load of 100 gf for 15 s as per the ASTM standard E384-08. Compressive testing of the cylindrical specimens was performed at room temperature according to the procedures outlined in ASTM standard E9-89a using Universal testing machine-Lloyd. Tensile testing of the extruded pure Al and its composite samples was done using a universal testing machine-Lloyd according to the ASTM E8/E8M-15a standard at room temperature under the strain rate of 8.3 × 10−4 s−1. For tensile tests, round test specimens of 25 mm gauge length and 5 mm gauge diameter (**Figure 4(b)**) were prepared and tested on a fully automated servo-hydraulic mechanical testing machine, MTS-810. For every composition, three samples were tested to check repeatable values. The fractured surfaces of the selected compression and tensile specimens were studied by scanning electron microscope (Hitachi FESEM-S4300). Nanoindentation investigation was done using a MFP-3D Nanoidenter (head connected to AFM equipment) system equipped with standard Berkovich diamond indenter tip. The testing was performed at room temperature and the values of hardness (H) and Young's modulus (E) were directly obtained. The presented nanoindentation results are the average of six indentations values. Coefficients of thermal expansion of pure Al and developed composites were determined in the temperature range of 50–350°C using a INSEIS TMA PT 1000LT thermo-mechanical analyzer. A heating rate of 5°C/min was employed at argon flow rate of 0.1 lpm.

**4.2. SEM analysis of AMMCs**

**Figure 5.** XRD patterns for (a) pure Al, (b) Al-SiC, (c) Al-Si<sup>3</sup>

**(c) (d)**

**Figure 6.** FESEM images for (a) pure Al, (b) Al-SiC, (c) Al-Si3

**(a) (b)**

) composites.

and Al2

O3

Scanning electron microscopy (SEM) was used in order to analyze the morphology and microstructure of developed composites containing different reinforcements. **Figure 6** shows the SEM micrographs of microwave sintered-hot extruded pure Al and Al-X (X = SiC, Si3

N4

, and (d) Al-Al2

O3 composites.

N4

, and (d) Al-Al2

O3

Development of Metal Matrix Composites Using Microwave Sintering Technique

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29

composites.

N4 ,

### **4. Properties of Al metal matrix composites**

#### **4.1. X-ray diffraction analysis of AMMCs**

The X-ray diffraction (XRD) patterns for the microwave sintered and hot extruded pure Al and Al metal matrix with reinforcement particles SiC, Si3 N4 , and Al2 O3 are shown in **Figure 5**.

The diffraction peaks of Al, SiC, Si<sup>3</sup> N4 , and Al2 O3 phases can be observed. The sharp peaks representing the presence of Al phase in the XRD patterns. The presence of SiC, Si3 N4 , and Al2 O3 particles is indicated by minor peaks. The peaks of Al are indexed as (2 2 0), (3 1 1), (1 1 1), (2 0 0), and (2 2 2), whereas SiC peaks are as (1 1 1), (2 0 0); Si3 N4 peak as (1 2 0); and Al2 O3 peaks are as (0 1 2), (1 0 4), (1 1 3), (0 2 4), (1 1 6), (2 1 4), (3 0 0). During the microwave sintering and hot subsequent extrusion process, it is noted that no solid-state reaction took place between the matrix and reinforcement to form any other undesired phases. The XRD results also approve the elemental mapping results, as will be shown later in **Figure 7**, which verifies the fabrication of phase pure different ceramic-reinforced Al-composites.

**Figure 5.** XRD patterns for (a) pure Al, (b) Al-SiC, (c) Al-Si<sup>3</sup> N4 , and (d) Al-Al2 O3 composites.

#### **4.2. SEM analysis of AMMCs**

The phase identification of the extruded samples was carried out using X-ray powder diffractometer (PANalytical X'pert Pro) based on Cu-Kα radiation (1.541 Å) in the 2*θ* range of 30–80° at scan rate of 0.2°/min. Individual phases were identified by matching the typical X-ray diffraction (XRD) peaks against JCPDS data. Field emission scanning electron microscopy (SEM) (JEOL JSM-6010 and Hitachi FESEM-S4300) with energy dispersion spectroscopy (EDS) was used to identify the reinforcement phase and microstructure of the extruded com-

The hardness testing of the pure Al and composite samples was carried out using Vicker's hardness tester with applied load of 100 gf for 15 s as per the ASTM standard E384-08. Compressive testing of the cylindrical specimens was performed at room temperature according to the procedures outlined in ASTM standard E9-89a using Universal testing machine-Lloyd. Tensile testing of the extruded pure Al and its composite samples was done using a universal testing machine-Lloyd according to the ASTM E8/E8M-15a standard at room temperature under the strain rate of 8.3 × 10−4 s−1. For tensile tests, round test specimens of 25 mm gauge length and 5 mm gauge diameter (**Figure 4(b)**) were prepared and tested on a fully automated servo-hydraulic mechanical testing machine, MTS-810. For every composition, three samples were tested to check repeatable values. The fractured surfaces of the selected compression and tensile specimens were studied by scanning electron microscope (Hitachi FESEM-S4300). Nanoindentation investigation was done using a MFP-3D Nanoidenter (head connected to AFM equipment) system equipped with standard Berkovich diamond indenter tip. The testing was performed at room temperature and the values of hardness (H) and Young's modulus (E) were directly obtained. The presented nanoindentation results are the average of six indentations values. Coefficients of thermal expansion of pure Al and developed composites were determined in the temperature range of 50–350°C using a INSEIS TMA PT 1000LT thermo-mechanical analyzer. A heating rate of 5°C/min was employed at argon

The X-ray diffraction (XRD) patterns for the microwave sintered and hot extruded pure Al

O3

particles is indicated by minor peaks. The peaks of Al are indexed as (2 2 0), (3 1 1), (1 1 1), (2 0 0),

as (0 1 2), (1 0 4), (1 1 3), (0 2 4), (1 1 6), (2 1 4), (3 0 0). During the microwave sintering and hot subsequent extrusion process, it is noted that no solid-state reaction took place between the matrix and reinforcement to form any other undesired phases. The XRD results also approve the elemental mapping results, as will be shown later in **Figure 7**, which verifies the fabrication

N4

N4

, and Al2

O3

peak as (1 2 0); and Al2

phases can be observed. The sharp peaks rep-

are shown in **Figure 5**.

N4

O3

, and Al2

peaks are

O3

posite samples.

28 Sintering of Functional Materials

flow rate of 0.1 lpm.

**4. Properties of Al metal matrix composites**

and Al metal matrix with reinforcement particles SiC, Si3

and (2 2 2), whereas SiC peaks are as (1 1 1), (2 0 0); Si3

of phase pure different ceramic-reinforced Al-composites.

N4

, and Al2

resenting the presence of Al phase in the XRD patterns. The presence of SiC, Si3

**4.1. X-ray diffraction analysis of AMMCs**

The diffraction peaks of Al, SiC, Si<sup>3</sup>

Scanning electron microscopy (SEM) was used in order to analyze the morphology and microstructure of developed composites containing different reinforcements. **Figure 6** shows the SEM micrographs of microwave sintered-hot extruded pure Al and Al-X (X = SiC, Si3 N4 , and Al2 O3 ) composites.

**Figure 6.** FESEM images for (a) pure Al, (b) Al-SiC, (c) Al-Si3 N4 , and (d) Al-Al2 O3 composites.

The results revealed a fair uniform distribution of ceramic reinforcement particles in the aluminum matrix. It can be further noted that SEM images show two main phases: the grey matrix is the Al phase while the dispersed phase showing white spots represents the SiC, Si3 N4 , and Al2 O3 particles used as reinforcements. At some spaces, agglomeration of reinforcement particulates has been observed (**Figure 6(d)**) which is due to density differences of reinforcements and the aluminum matrix.

The mechanical properties of Al-MMCs are dependent on the nature and distribution of the reinforcement particles. Homogeneous and intragranular distribution is preferred to attain improved properties, and importantly, the hot extrusion process has led to the desirable distribution. Previous studies have reported that agglomeration of reinforcement particles in Al matrix has resulted in the degradation of mechanical properties, as reinforcement clustering along with voids act as pre-existing cracks, limiting the stress transfer from soft matrix to hard phase particles during the deformation process [36, 37].

In our developed Al-composites, the agglomeration of reinforcements is observed only in a few locations, confirming a uniform reinforcement distribution in the Al-X composites. This near-uniform distribution of reinforcement promotes even heating (by absorbing the microwave energy) through the compact during sintering and demonstrates the effectiveness of using powder metallurgy and microwave sintering for the synthesis of Al-based composites [38]. Hence, the SEM results show that microwave sintering followed by hot extrusion process has an appropriate potential for manufacturing the ceramic particle-reinforced metal matrix composites.

#### **4.3. EDS analysis of AMMCs**

The energy dispersive spectroscopy (EDS) technique was used to study the composition and elemental distribution of phases present in the Al-based composites.

fraction of the reinforced ceramic particles and the effect of hot extrusion, as a secondary processing, play an active role in improving the mechanical properties of the composites.

N4

, and (c) Al2

O3 .

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31

**Al Ka 1 O Ka 1**

**Al Ka 1 Si Ka 1 C Ka 1\_**

The microhardness is a very useful important property that reflects the strength of the material. Generally, several factors would affect the microhardness of the composites, such as particle shape, size, amount, distribution, density of reinforcement, and method of preparation [39].

composites exhibit a hardness of 82 ± 4, 101 ± 3, and 92 ± 5 Hv, respectively; these values are comparatively higher than the unreinforced aluminum. However, a remarkable enhancement

AMMCs indicates that the ceramic particles has a major contribution in the strengthening of Al matrix. This increase in the hardness is because of the contribution of the reduced crystallite

The presence of hard ceramic particles can enhance the microhardness of composites accord-

*<sup>m</sup>* + *Hr f*

N4 , and

O3

) composites is

N4

reinforced Al composite. The increase in the microhardness of

, and Al2

*<sup>r</sup>* (1)

N4

O3

, and Al-15 vol.% Al<sup>2</sup>

**Figure 8** shows the microhardness of the microwave sintered-extruded pure Al, SiC, Si3

) compared to 35 nm (SiC) in the composite.

reinforced composite. The hardness of the Al-X (X = SiC, Si3

higher than the pure aluminum. The Al-1.5 vol.% SiC, Al-1.5 vol.% Si<sup>3</sup>

N4

*4.4.1. Microhardness studies of AMMCs*

**Figure 7.** EDS maps of AMMCs reinforced with (a) SiC, (b) Si3

**(c)**

to 101 ± 3 is observed for Si3

ing to the rule of mixtures [40].

N4

*Hc* = *Hm f*

size of ~15 nm (Si3

Al2 O3

**(b)**

**(a)**

**Figure 7(a–c)** shows the EDS mapping analysis of microwave-hot extruded pure Al and Al-X (X = SiC, Si3 N4 , and Al2 O3 ) composites. The elemental distribution of phases such as aluminum (matrix) and ceramic particles (reinforcements) is clearly observable. Furthermore, the reinforcing elements are uniformly dispersed all over the aluminum matrix. This confirms the appropriate mixing of ceramic reinforcement particles with the aluminum matrix. **Figure 7(a)** represents the corresponding Al, Si, and composition maps of the Al-SiC composite. **Figure 7(b)** shows the corresponding Al, Si, and N composition maps of the Al-Si3 N4 composite. **Figure 7(c)** shows the EDS Al and O composition maps of the Al-Al2 O3 composite. Further, the elemental distribution map evidently reveals the uniform distribution of ceramic reinforcement particles in Al matrix and confirms the presence of aluminum, SiC, Si<sup>3</sup> N4 , and Al2 O3 phases in respective composition.

#### **4.4. Mechanical properties of AMMCs**

The amount and size of reinforcement particles, type of processing technique, and the matrix/ particle integrity greatly influence the mechanical properties of an Al-based composite. A strong matrix/particle interface integrity was obtained in this study. Therefore, the volume

**Figure 7.** EDS maps of AMMCs reinforced with (a) SiC, (b) Si3 N4 , and (c) Al2 O3 .

fraction of the reinforced ceramic particles and the effect of hot extrusion, as a secondary processing, play an active role in improving the mechanical properties of the composites.

#### *4.4.1. Microhardness studies of AMMCs*

The results revealed a fair uniform distribution of ceramic reinforcement particles in the aluminum matrix. It can be further noted that SEM images show two main phases: the grey matrix is the Al phase while the dispersed phase showing white spots represents the SiC,

ment particulates has been observed (**Figure 6(d)**) which is due to density differences of rein-

The mechanical properties of Al-MMCs are dependent on the nature and distribution of the reinforcement particles. Homogeneous and intragranular distribution is preferred to attain improved properties, and importantly, the hot extrusion process has led to the desirable distribution. Previous studies have reported that agglomeration of reinforcement particles in Al matrix has resulted in the degradation of mechanical properties, as reinforcement clustering along with voids act as pre-existing cracks, limiting the stress transfer from soft matrix to hard

In our developed Al-composites, the agglomeration of reinforcements is observed only in a few locations, confirming a uniform reinforcement distribution in the Al-X composites. This near-uniform distribution of reinforcement promotes even heating (by absorbing the microwave energy) through the compact during sintering and demonstrates the effectiveness of using powder metallurgy and microwave sintering for the synthesis of Al-based composites [38]. Hence, the SEM results show that microwave sintering followed by hot extrusion process has an appropriate potential for manufacturing the ceramic particle-reinforced metal matrix

The energy dispersive spectroscopy (EDS) technique was used to study the composition and

**Figure 7(a–c)** shows the EDS mapping analysis of microwave-hot extruded pure Al and

aluminum (matrix) and ceramic particles (reinforcements) is clearly observable. Furthermore, the reinforcing elements are uniformly dispersed all over the aluminum matrix. This confirms the appropriate mixing of ceramic reinforcement particles with the aluminum matrix. **Figure 7(a)** represents the corresponding Al, Si, and composition maps of the Al-SiC composite. **Figure 7(b)** shows the corresponding Al, Si, and N composition maps of the Al-Si3

Further, the elemental distribution map evidently reveals the uniform distribution of ceramic

The amount and size of reinforcement particles, type of processing technique, and the matrix/ particle integrity greatly influence the mechanical properties of an Al-based composite. A strong matrix/particle interface integrity was obtained in this study. Therefore, the volume

composite. **Figure 7(c)** shows the EDS Al and O composition maps of the Al-Al2

reinforcement particles in Al matrix and confirms the presence of aluminum, SiC, Si<sup>3</sup>

) composites. The elemental distribution of phases such as

N4

O3 composite.

N4 , and

elemental distribution of phases present in the Al-based composites.

O3

particles used as reinforcements. At some spaces, agglomeration of reinforce-

Si3 N4

, and Al2

30 Sintering of Functional Materials

composites.

Al-X (X = SiC, Si3

Al2 O3

**4.3. EDS analysis of AMMCs**

N4

, and Al2

phases in respective composition.

**4.4. Mechanical properties of AMMCs**

O3

forcements and the aluminum matrix.

phase particles during the deformation process [36, 37].

The microhardness is a very useful important property that reflects the strength of the material. Generally, several factors would affect the microhardness of the composites, such as particle shape, size, amount, distribution, density of reinforcement, and method of preparation [39].

**Figure 8** shows the microhardness of the microwave sintered-extruded pure Al, SiC, Si3 N4 , and Al2 O3 reinforced composite. The hardness of the Al-X (X = SiC, Si3 N4 , and Al2 O3 ) composites is higher than the pure aluminum. The Al-1.5 vol.% SiC, Al-1.5 vol.% Si<sup>3</sup> N4 , and Al-15 vol.% Al<sup>2</sup> O3 composites exhibit a hardness of 82 ± 4, 101 ± 3, and 92 ± 5 Hv, respectively; these values are comparatively higher than the unreinforced aluminum. However, a remarkable enhancement to 101 ± 3 is observed for Si3 N4 reinforced Al composite. The increase in the microhardness of AMMCs indicates that the ceramic particles has a major contribution in the strengthening of Al matrix. This increase in the hardness is because of the contribution of the reduced crystallite size of ~15 nm (Si3 N4 ) compared to 35 nm (SiC) in the composite.

The presence of hard ceramic particles can enhance the microhardness of composites according to the rule of mixtures [40].

$$H\_c = -H\_m f\_m + H\_r f\_r \tag{1}$$

**Figure 8.** Hardness of aluminum metal matrix composites.

where *Hc* represents hardness of the composite, *Hm* and *Hr* represent hardness of the matrix and the reinforcing particle, respectively, and *f <sup>m</sup>* and *f r* represent the volume fraction of the matrix and the reinforcing particle, respectively.

an increase of ~31 and ~115%, respectively, compared to pure Al. The addition of micron-sized

This significant improvement in compression strength properties of the extruded Al-X

effects of (i) uniform distribution of reinforcing particles in the matrix and (ii) enhanced dislocation density [39]. For a clearer comparison, we have noticed that the compressive properties

**CYS (MPa) UCS (MPa) Failure** 

Pure Al 5.15 ± 0.3 73 ± 5 70 ± 3 313 ± 5 7.17 105 ± 2 119 ± 4 13.6 ± 0.3

9.60 ± 0.6 81 ± 6 114 ± 7 392 ± 6 7.48 158 ± 9 178 ± 6 7.3 ± 0.9

16.34±0.4 94±2 142±6 412±3 8.07 165±5 191±5 8.2±0.4

24.56 ± 0.8 106 ± 9 136 ± 5 388 ± 8 6.19 139 ± 8 154 ± 6 7.2 ± 0.7

) composites exhibited higher compressive failure strain values when compared to

which is ~24 and ~106%, respectively, higher than that of pure Al. The Al-X (X = SiC, Si<sup>3</sup>

composite exhibited (UCS) ~136 MPa and (CYS) 338 MPa

Development of Metal Matrix Composites Using Microwave Sintering Technique

http://dx.doi.org/ 10.5772/68081

33

) composites compared to the pure Al can be ascribed to the coupled

N4

**Strain (%)**

, and Al2

O3

) composites are inter-

**(%)**

**TYS (MPa) UTS (MPa) Elongation** 

N4 ,

O3

**Figure 9.** Compressive stress-strain curves of aluminum metal matrix composites.

alumina particles, Al-15 vol.% Al<sup>2</sup>

, and Al2

O3

**Young's modulus (GPa)**

**Table 1.** Mechanical properties of aluminum metal matrix composites.

of the microwave sintered-hot extruded Al-X (X = SiC, Si3

estingly superior to that of conventional sintered AMMCs [43–47].

**Materials Nanoindentation data Compressive properties Tensile properties**

and Al2

(X = SiC, Si3

Al-1.5 vol.% SiC

Al-1.5 vol.% Si<sup>3</sup> N4

Al-15 vol.% Al2 O3

O3

that of pure Al (~7.1%).

N4

**Hardness (GPa)**

The dispersion of hard ceramic reinforcement in the soft aluminum matrix results in strengthening of the structure. Referring to Hall-Petch relationship, the mechanical properties of the metallic materials are affected by the grain size. The grain size of metal matrix composites is smaller than that of the aluminum matrix because of grain refinement of reinforced ceramic particles. The fine grains enhance the hardness of the resulting structure. In addition, the difference in thermal shrinkage between the aluminum matrix and the ceramic particles produces quench hardening effect [41]. The presence of hard ceramic particles also improves the mechanical properties due to dispersion hardening of soft aluminum matrix. In fact, the presence of hard particles impedes the motion dislocation and thus improves the mechanical properties [42].

#### *4.4.2. Compressive studies of AMMCs*

The true stress-strain curves of the microwave sintered-hot extruded pure Al and Al-X (X = SiC, Si3 N4 , and Al2 O3 ) composites under compression loading at room temperature are shown in **Figure 9**. The average compressive yield strength (CYS) and ultimate compressive strength (UCS) values of the extruded composites are listed in **Table 1**.

A significant improvement in the strength of Al-X composites are observed compared to pure aluminum. The compression strength of pure Al was increased by adding various ceramic reinforcement particles. The Al-1.5 vol.% SiC composite showed the compressive yield strength (0.2% CYS) and the ultimate compressive strength (UCS) of ~114 and ~392 MPa, respectively, the incremental increase is ~26 and ~72%, respectively, compared to pure Al. In the case of Al-1.5 vol.% Si<sup>3</sup> N4 composite, the (CYS) (UCS) were ~142 and ~412 MPa, respectively, showing

**Figure 9.** Compressive stress-strain curves of aluminum metal matrix composites.

where *Hc*

32 Sintering of Functional Materials

properties [42].

(X = SiC, Si3

Al-1.5 vol.% Si<sup>3</sup>

*4.4.2. Compressive studies of AMMCs*

, and Al2

N4

O3

strength (UCS) values of the extruded composites are listed in **Table 1**.

N4

represents hardness of the composite, *Hm* and *Hr*

*<sup>m</sup>* and *f r*

The dispersion of hard ceramic reinforcement in the soft aluminum matrix results in strengthening of the structure. Referring to Hall-Petch relationship, the mechanical properties of the metallic materials are affected by the grain size. The grain size of metal matrix composites is smaller than that of the aluminum matrix because of grain refinement of reinforced ceramic particles. The fine grains enhance the hardness of the resulting structure. In addition, the difference in thermal shrinkage between the aluminum matrix and the ceramic particles produces quench hardening effect [41]. The presence of hard ceramic particles also improves the mechanical properties due to dispersion hardening of soft aluminum matrix. In fact, the presence of hard particles impedes the motion dislocation and thus improves the mechanical

The true stress-strain curves of the microwave sintered-hot extruded pure Al and Al-X

shown in **Figure 9**. The average compressive yield strength (CYS) and ultimate compressive

A significant improvement in the strength of Al-X composites are observed compared to pure aluminum. The compression strength of pure Al was increased by adding various ceramic reinforcement particles. The Al-1.5 vol.% SiC composite showed the compressive yield strength (0.2% CYS) and the ultimate compressive strength (UCS) of ~114 and ~392 MPa, respectively, the incremental increase is ~26 and ~72%, respectively, compared to pure Al. In the case of

) composites under compression loading at room temperature are

composite, the (CYS) (UCS) were ~142 and ~412 MPa, respectively, showing

and the reinforcing particle, respectively, and *f*

**Figure 8.** Hardness of aluminum metal matrix composites.

matrix and the reinforcing particle, respectively.

represent hardness of the matrix

represent the volume fraction of the

an increase of ~31 and ~115%, respectively, compared to pure Al. The addition of micron-sized alumina particles, Al-15 vol.% Al<sup>2</sup> O3 composite exhibited (UCS) ~136 MPa and (CYS) 338 MPa which is ~24 and ~106%, respectively, higher than that of pure Al. The Al-X (X = SiC, Si<sup>3</sup> N4 , and Al2 O3 ) composites exhibited higher compressive failure strain values when compared to that of pure Al (~7.1%).

This significant improvement in compression strength properties of the extruded Al-X (X = SiC, Si3 N4 , and Al2 O3 ) composites compared to the pure Al can be ascribed to the coupled effects of (i) uniform distribution of reinforcing particles in the matrix and (ii) enhanced dislocation density [39]. For a clearer comparison, we have noticed that the compressive properties of the microwave sintered-hot extruded Al-X (X = SiC, Si3 N4 , and Al2 O3 ) composites are interestingly superior to that of conventional sintered AMMCs [43–47].


**Table 1.** Mechanical properties of aluminum metal matrix composites.

#### *4.4.3. Tensile studies of AMMCs*

The representative tensile stress-strain curves for microwave sintered-hot extruded pure Al and Al-X (X = SiC, Si3 N4 , and Al2 O3 ) composites at room temperature are shown in **Figure 10**. The variations in the tensile strength, yield strength, and ductility with reinforcement addition are listed in **Table 1**. It can be observed that all composites exhibited higher tensile strengths in comparison to that of pure Al. However, the elongation of the composites decreases as compared to pure Al. The calculated decrease is in elongation compared to microwave sintered-extruded pure Al is 46, 39, and 47% for Al-1.5 vol.% SiC, Al-1.5 vol.% Si<sup>3</sup> N4 , and Al-15 vol.% Al<sup>2</sup> O3 composites, respectively.

The efficient load transfer (*σ*load) between the ductile matrix and the hard ceramic reinforcement particles during tensile testing occurs, particularly when there is a good interfacial contact between the matrix and the reinforcement and is represented as following [48–50]:

*σload* = 0.5 *Vf σYM* (2)

The interaction between the dislocations and the reinforcement particles enhances the strength of the composite materials in agreement with the Orowan mechanism. Due to the existence of dispersed reinforcement particles in the matrix, dislocation loops are formed when disloca-

where *G* is the shear modulus of matrix, *b* is the Burgers vector, *λ* is the inter-particle spacing,

The difference in the CTE values of the reinforcement particles and the metal matrix produces geometrically necessary dislocations and thermally induced residual stresses. The thermal stresses at the particles and matrix interface make the plastic deformation more tough which, hence, enhances the level of hardness and flow stress. The mismatch strain effect due to the

difference between the CTE values of particles and that of the matrix is given by [38]:

\_\_ 3 *βGm b*

forcement, and Δ*T* is the difference between the test and process temperature.

particles and Al of the microwave sintered-extruded Al-X (X = SiC, Si3

where *b* is the strengthening coefficient, Δα is the difference between CTE of matrix and rein-

Some selected compression and tensile tested fractured surfaces were studied using SEM in order to understand the type of fracture and nature of the bonding between the reinforcing

The fracture morphology of microwave sintered-extruded pure Al and Al-X (X = SiC, Si3

) composites during compression test are shown in **Figure 11(a–d)**. The fracture surfaces are comparatively smooth and the formation of shear band can barely be seen in the fractured samples. The fractured compressive samples reveal a crack at 45° to the test axis. **Figure 11(a–d)** shows a typical shear mode fracture in pure Al and Al-X (composites reinforced with various ceramic particles. It approves that the compressive deformation of the Al composites is expressively indifferent. This is due to heterogeneous deformation and work hardening behavior [52]. In contrast, mixed fracture surface and the shear band formation is found in Al-Al2

√

\_\_\_\_\_\_\_\_\_ <sup>24</sup> *Vf* <sup>Δ</sup>*α*Δ*<sup>T</sup>* \_\_\_\_\_\_\_\_\_ (1 − *Vf* ) *b rp*

tions interact with the reinforcing particles. *σOrowan* can be calculated as [51]:

*<sup>σ</sup>Orowan* <sup>=</sup> \_\_\_\_\_\_ 0.13*Gb*

Δ *σCTE* = √

and *r* is the particle radius.

*4.4.4. Fractography of AMMCs*

Al2 O3 is the volume fraction of ceramic reinforcement particles and *σYM* is the matrix yield

Development of Metal Matrix Composites Using Microwave Sintering Technique

*<sup>λ</sup>* ln \_\_*<sup>r</sup>*

*<sup>b</sup>* (3)

http://dx.doi.org/ 10.5772/68081

35

N4 , Al2 O3 (4)

) composites.

N4 ,

O3

where *Vf*

stress.

Moreover, **Table 1** also shows that the tensile properties of the Al-1. 5 vol.% Si<sup>3</sup> N4 composites are comparable/superior to that of the SiC and Al2 O3 reinforced Al composites. This can be endorsed to the reduced size of the reinforcing particles employed [37]. Like compressive properties, the tensile properties of microwave sintered-extruded Al-X (X = SiC, Si3 N4 , and Al2 O3 ) composites are found superior to the conventional sintered AMMCs [31–35].

To understand the strengthening effects of ceramic reinforcement particles on the hardness, compression, and tensile properties of composites, such as UTS and YS, it is favorable to discuss the strengthening mechanism in detail. In the present study, strengthening occurs due to following mechanisms: (i) active load transfer from the matrix to the reinforcement, (ii) Orowan strengthening, and (iii) generation of internal thermal stresses because of the difference in the coefficient of thermal expansion (CTE) between the reinforcement particles and matrix phase.

**Figure 10.** Tensile stress-strain curves of aluminum metal matrix composites.

The efficient load transfer (*σ*load) between the ductile matrix and the hard ceramic reinforcement particles during tensile testing occurs, particularly when there is a good interfacial contact between the matrix and the reinforcement and is represented as following [48–50]:

$$
\sigma\_{load} = 0.5 \, V\_f \sigma\_{\text{YM}} \tag{2}
$$

where *Vf* is the volume fraction of ceramic reinforcement particles and *σYM* is the matrix yield stress.

The interaction between the dislocations and the reinforcement particles enhances the strength of the composite materials in agreement with the Orowan mechanism. Due to the existence of dispersed reinforcement particles in the matrix, dislocation loops are formed when dislocations interact with the reinforcing particles. *σOrowan* can be calculated as [51]:

$$
\sigma\_{\text{Oouun}} = \frac{0.13Gb}{\lambda} \ln \frac{r}{b} \tag{3}
$$

where *G* is the shear modulus of matrix, *b* is the Burgers vector, *λ* is the inter-particle spacing, and *r* is the particle radius.

The difference in the CTE values of the reinforcement particles and the metal matrix produces geometrically necessary dislocations and thermally induced residual stresses. The thermal stresses at the particles and matrix interface make the plastic deformation more tough which, hence, enhances the level of hardness and flow stress. The mismatch strain effect due to the difference between the CTE values of particles and that of the matrix is given by [38]:

$$
\Delta \sigma\_{\rm CT} = \sqrt{3} \,\beta G\_m \, b \sqrt{\frac{24 \, V\_f \Delta a \Delta T}{(1 - V\_f) \, b \, r\_p}} \tag{4}
$$

where *b* is the strengthening coefficient, Δα is the difference between CTE of matrix and reinforcement, and Δ*T* is the difference between the test and process temperature.

#### *4.4.4. Fractography of AMMCs*

*4.4.3. Tensile studies of AMMCs*

N4

composites, respectively.

are comparable/superior to that of the SiC and Al2

**Figure 10.** Tensile stress-strain curves of aluminum metal matrix composites.

, and Al2

O3

tered-extruded pure Al is 46, 39, and 47% for Al-1.5 vol.% SiC, Al-1.5 vol.% Si<sup>3</sup>

Moreover, **Table 1** also shows that the tensile properties of the Al-1. 5 vol.% Si<sup>3</sup>

and Al-X (X = SiC, Si3

34 Sintering of Functional Materials

O3

vol.% Al<sup>2</sup>

Al2 O3

matrix phase.

The representative tensile stress-strain curves for microwave sintered-hot extruded pure Al

The variations in the tensile strength, yield strength, and ductility with reinforcement addition are listed in **Table 1**. It can be observed that all composites exhibited higher tensile strengths in comparison to that of pure Al. However, the elongation of the composites decreases as compared to pure Al. The calculated decrease is in elongation compared to microwave sin-

endorsed to the reduced size of the reinforcing particles employed [37]. Like compressive

To understand the strengthening effects of ceramic reinforcement particles on the hardness, compression, and tensile properties of composites, such as UTS and YS, it is favorable to discuss the strengthening mechanism in detail. In the present study, strengthening occurs due to following mechanisms: (i) active load transfer from the matrix to the reinforcement, (ii) Orowan strengthening, and (iii) generation of internal thermal stresses because of the difference in the coefficient of thermal expansion (CTE) between the reinforcement particles and

properties, the tensile properties of microwave sintered-extruded Al-X (X = SiC, Si3

) composites are found superior to the conventional sintered AMMCs [31–35].

O3

) composites at room temperature are shown in **Figure 10**.

N4

N4

reinforced Al composites. This can be

, and Al-15

composites

N4 , and

> Some selected compression and tensile tested fractured surfaces were studied using SEM in order to understand the type of fracture and nature of the bonding between the reinforcing particles and Al of the microwave sintered-extruded Al-X (X = SiC, Si3 N4 , Al2 O3 ) composites.

> The fracture morphology of microwave sintered-extruded pure Al and Al-X (X = SiC, Si3 N4 , Al2 O3 ) composites during compression test are shown in **Figure 11(a–d)**. The fracture surfaces are comparatively smooth and the formation of shear band can barely be seen in the fractured samples. The fractured compressive samples reveal a crack at 45° to the test axis. **Figure 11(a–d)** shows a typical shear mode fracture in pure Al and Al-X (composites reinforced with various ceramic particles. It approves that the compressive deformation of the Al composites is expressively indifferent. This is due to heterogeneous deformation and work hardening behavior [52]. In contrast, mixed fracture surface and the shear band formation is found in Al-Al2 O3

**(a) (b)**

Development of Metal Matrix Composites Using Microwave Sintering Technique

http://dx.doi.org/ 10.5772/68081

37

**(c) (d)**

**Figure 13.** CTE of the aluminum metal matrix composites.

Al-Al2 O3 .

**Figure 12.** SEM micrographs of the tensile fracture surfaces of (a) pure aluminum, (b)Al-SiC, (c) Al-Si3

N4

, and (d)

**Figure 11.** SEM micrographs of the compression fracture surfaces of (a) pure aluminum, (b)Al-SiC, (c) Al-Si3 N4 , and (d) Al-Al2 O3 .

composite (**Figure 11(d)**). The plastic deformation in the composites was inhibited due to the dispersion of second phases. This led to the significant reduction in compressive failure strain in the composite (see **Table 1**).

Fracture morphology of pure Al and its composites during tensile testing are presented in **Figure 12**. The examination of fractured surfaces reveals the formation of similar ductile fracture in all composites. For Al-SiC composites, the fractured surfaces shows dimple-like fracture which can be related to the observed failure strain of more than 7% (see **Table 1**). The presence of SiC particles in the dimple cores and walls suggests that the fractured particles and agglomerates are potential stress concentration sites and susceptible to void formation. Al-Si3 N4 composites showed the strongest bonding as revealed by the good matrix/reinforcement performance. For Al-Al2 O3 composites, it can be seen that ductile failure occurs in the matrix, whereas brittle, cleavage-type failure is seen to be predominant in regions where Al2 O3 particles are present. Large number of dimples with tear ridges is also seen in the Al-Al2 O3 composite.

#### *4.4.5. Coefficient of thermal expansion of AMMCs*

The variation of CTE of microwave sintered-extruded pure Al and Al-X composites (X = SiC, Si3 N4 , Al2 O3 ) is shown in **Figure 13**. It can be observed that the CTE values decrease with reinforced ceramic particles. It is in accordance with the theory that the thermal expansion Development of Metal Matrix Composites Using Microwave Sintering Technique http://dx.doi.org/ 10.5772/68081 37

**Figure 12.** SEM micrographs of the tensile fracture surfaces of (a) pure aluminum, (b)Al-SiC, (c) Al-Si3 N4 , and (d) Al-Al2 O3 .

**Figure 13.** CTE of the aluminum metal matrix composites.

composite (**Figure 11(d)**). The plastic deformation in the composites was inhibited due to the dispersion of second phases. This led to the significant reduction in compressive failure strain

**Figure 11.** SEM micrographs of the compression fracture surfaces of (a) pure aluminum, (b)Al-SiC, (c) Al-Si3

Fracture morphology of pure Al and its composites during tensile testing are presented in **Figure 12**. The examination of fractured surfaces reveals the formation of similar ductile fracture in all composites. For Al-SiC composites, the fractured surfaces shows dimple-like fracture which can be related to the observed failure strain of more than 7% (see **Table 1**). The presence of SiC particles in the dimple cores and walls suggests that the fractured particles and agglomer-

ites showed the strongest bonding as revealed by the good matrix/reinforcement performance.

The variation of CTE of microwave sintered-extruded pure Al and Al-X composites (X = SiC,

reinforced ceramic particles. It is in accordance with the theory that the thermal expansion

composites, it can be seen that ductile failure occurs in the matrix, whereas brittle,

) is shown in **Figure 13**. It can be observed that the CTE values decrease with

N4

N4 , and (d)

particles are present.

O3

composite.

O3

compos-

ates are potential stress concentration sites and susceptible to void formation. Al-Si3

cleavage-type failure is seen to be predominant in regions where Al2

Large number of dimples with tear ridges is also seen in the Al-Al2

**(a) (b)**

36 Sintering of Functional Materials

**(c) (d)**

*4.4.5. Coefficient of thermal expansion of AMMCs*

in the composite (see **Table 1**).

For Al-Al2

Al-Al2 O3 .

Si3 N4 , Al2 O3

O3

of the composites is governed by the competing interactions of expansion of Al matrix and the constraint of ceramic particles through their interfaces [52]. The CTE of pure Al was measured to be 23.31 × 10−6/K which is in close agreement with the theoretical CTE of aluminum (24 × 10−6/K). The addition of nano-sized 1.5 vol.% SiC, nano-sized 1.5 vol.% Si<sup>3</sup> N4 , and micronsized 15 vol.% Al<sup>2</sup> O3 particles to Al reduced the CTE value to ~19.20 × 10−6/K, 19.43 × 10−6/K, and 19.66 × 10−6/K which are ~17.63, ~16.66, and ~15.65% reduction when compared to pure Al. This considerable decrease in CTE values may be due to the high thermal stability of SiC, Si3 N4 , and Al2 O3 reinforcement particles having theoretical CTE of 4.3 × 10−6/K [53], 3.3 × 10−6/K [54], 7.4 × 10−6/K [53], respectively. The linear decrease in CTE values with the addition of ceramic particles can be attributed to: (i) the lower CTE values of ceramic (SiC, Si<sup>3</sup> N4 , and Al2 O3 ) particles reinforcements as compared to that of the pure Al matrix; and (ii) uniform distribution of the ceramic reinforcements in the matrix.

• Coefficient of thermal expansion values decreases with the addition of ceramic-reinforced particles into Al matrix confirming high-dimensional stability of microwave sintered-ex-

This chapter was made possible by NPRP Grant 7-159-2-076 from Qatar National Research Fund (a member of the Qatar Foundation). Statements made herein are solely the responsibility

2 Department of Metallurgical and Materials Engineering, Suez University, Suez, Egypt

[1] Sutton WH. Microwave processing of ceramic materials. American Ceramic Society

[2] Sutton WH. Ceramic transactions microwaves: Theory and applications in materials

[3] Das S, Mukhopadhyay AK, Datta S, Basu D. Prospects of microwave processing: An

[4] Xie Z, Wang C, Fan X, Huang Y. Microwave processing and properties of Ce-Y-ZrO<sup>2</sup>

[5] Agrawal DK. Microwave processing of ceramics: A review. Current Opinion in Solid

[6] Penchal Reddy M, Madhuri W, Ramamanohar Reddy N, Siva Kumar KV, Murthy VRK, Ramakrishna Reddy R. Magnetic properties of Ni-Zn ferrites prepared by microwave

[7] Hao HS, Xu LH, Huang Y, Zhang XM, Xie ZP. Kinetics mechanism of microwave sintering in ceramic materials. Science in China Series E: Technological Sciences.

processing II. Key issues in microwave process technology. 1993;**36**:3-18

ceramics with 2.45 GHz irradiation. Materials Letters. 1999;**38**:190-196

) composites making them suitable for automotive and

http://dx.doi.org/ 10.5772/68081

39

Development of Metal Matrix Composites Using Microwave Sintering Technique

\* and Adel Mohamed Amer Mohamed2

truded Al-X (X = SiC, Si3

**Acknowledgements**

of the authors.

**References**

Bulletin. 1989;**68**:376-386

2009;**52**:2727-2731

**Author details**

Penchal Reddy Matli1

many other related applications.

N4 , Al2 O3

, Rana Abdul Shakoor1

1 Center for Advanced Materials, Qatar University, Doha, Qatar

overview. Bulletin of Materials Science. 2009;**32**:1-13

sintering method. Journal of Electroceramics. 2012;**28**:1-9

State Materials Science. 1998;**3**:480-486

\*Address all correspondence to: shakoor@qu.edu.qa

The compatible CTE of Al-X composites (X = SiC, Si3 N4 , Al2 O3 ) and high dimensional stability makes these microwave sintered-extruded composite very competitive for application in aerospace and automotive industry.

## **5. Conclusions**

Al-X (X = SiC, Si3 N4 , and Al2 O3 ) composites were successfully synthesized through microwave-assisted powder metallurgy route coupled with hot extrusion process. Various ceramic reinforcement particles were added into Al matrix, and their effect on structural, mechanical, and thermal properties has led to the following conclusions.


• Coefficient of thermal expansion values decreases with the addition of ceramic-reinforced particles into Al matrix confirming high-dimensional stability of microwave sintered-extruded Al-X (X = SiC, Si3 N4 , Al2 O3 ) composites making them suitable for automotive and many other related applications.

## **Acknowledgements**

of the composites is governed by the competing interactions of expansion of Al matrix and the constraint of ceramic particles through their interfaces [52]. The CTE of pure Al was measured to be 23.31 × 10−6/K which is in close agreement with the theoretical CTE of aluminum

and 19.66 × 10−6/K which are ~17.63, ~16.66, and ~15.65% reduction when compared to pure Al. This considerable decrease in CTE values may be due to the high thermal stability of SiC, Si3

7.4 × 10−6/K [53], respectively. The linear decrease in CTE values with the addition of ceramic

ticles reinforcements as compared to that of the pure Al matrix; and (ii) uniform distribution

ity makes these microwave sintered-extruded composite very competitive for application in

wave-assisted powder metallurgy route coupled with hot extrusion process. Various ceramic reinforcement particles were added into Al matrix, and their effect on structural, mechanical,

N4

coupled with extrusion process has an appropriate potential to process high-performance

• Homogeneous reinforcement particles distribution was found in microwave sintered-ex-

• A comparison of mechanical properties (hardness, strength) indicates that microwave sin-

microwave sintered-extruded pure Al. The improvement in mechanical properties can be attributed to (i) active load transfer from the matrix to the reinforcement, (ii) Orowan strengthening, and (iii) generation of internal thermal stresses because of the difference in the coefficient of thermal expansion (CTE) between the reinforcement particles and matrix phase.

, and Al2

O3 .

reinforcement particles having theoretical CTE of 4.3 × 10−6/K [53], 3.3 × 10−6/K [54],

N4 , Al2 O3

) composites were successfully synthesized through micro-

) composites indicate that the main components

) composites. This shows that the microwave sintering

) composites have superior properties compared to

) composites have low ductility compared to pure

particles to Al reduced the CTE value to ~19.20 × 10−6/K, 19.43 × 10−6/K,

N4

N4

) and high dimensional stabil-

, and Al2

, and micron-

O3 ) par-

N4 ,

N4 ,

(24 × 10−6/K). The addition of nano-sized 1.5 vol.% SiC, nano-sized 1.5 vol.% Si<sup>3</sup>

particles can be attributed to: (i) the lower CTE values of ceramic (SiC, Si<sup>3</sup>

sized 15 vol.% Al<sup>2</sup>

38 Sintering of Functional Materials

O3

**5. Conclusions**

Al-X (X = SiC, Si3

and Al2

O3

of the ceramic reinforcements in the matrix.

aerospace and automotive industry.

N4

• XRD patterns of Al-X (X = SiC, Si<sup>3</sup>

tered-extruded Al-X (X = SiC, Si3

• The produced Al-X (X = SiC, Si3

Al2 O3

truded Al-X (X = SiC, Si3

, and Al2

of the synthesized composites are Al, SiC, Si3

particle-reinforced metal matrix composites.

O3

and thermal properties has led to the following conclusions.

N4 , Al2 O3

N4 , Al2 O3

N4 , Al2 O3

N4 , Al2 O3

tured surfaces endorsing ductile mode of fracture.

Al due to low inherent ductility of ceramic particles used as reinforcement.

• The fractography results indicate that under compressive loading, the Al-X (X = SiC, Si3

) composites show the presence of shear bands which confirms the brittle mode of fracture. However, under tensile loading, the dimple formation was noticed on the frac-

The compatible CTE of Al-X composites (X = SiC, Si3

This chapter was made possible by NPRP Grant 7-159-2-076 from Qatar National Research Fund (a member of the Qatar Foundation). Statements made herein are solely the responsibility of the authors.

## **Author details**

Penchal Reddy Matli1 , Rana Abdul Shakoor1 \* and Adel Mohamed Amer Mohamed2

\*Address all correspondence to: shakoor@qu.edu.qa

1 Center for Advanced Materials, Qatar University, Doha, Qatar

2 Department of Metallurgical and Materials Engineering, Suez University, Suez, Egypt

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**Chapter 3**

**Provisional chapter**

) can be mixed with tricalcium phosphate

to the TCP‐40 wt% TiO<sup>2</sup>

= 27 MPa; H = 360 Hv; E = 51 GPa and G = 20 GPa)

, β‐TCP) based on biomaterial has attracted considerable

composites are similar to those of bone

and MgF<sup>2</sup>

for biomedi‐

composites

composites increase

, β‐TCP) to make bioceramic composites, which would combine the biocom‐

cal applications. The samples were characterized by different characterization techniques such as physicochemical and mechanical. The sintering of the TCP at various tempera‐ tures (1000, 1100, 1200 and 1300°C) with different percentages of titania (2.5, 5, 7.5, 10, 20,

with both the sintering temperature and the amount of titania. The highest values of the composites' (H = 270 Hv; E = 33.1 GPa and G = 15.7 GPa) were obtained with 40 wt% tita‐

at 1200°C for 1 h. The amelioration of these properties is due to the formation of a new compound and the liquid phase which helps to fill the pores in the microstructure. The

‐MgF<sup>2</sup>

**Keywords:** sintering, biomaterial, composites, mechanical properties, tricalcium

**Sintering of the Tricalcium Phosphate-Titania-**

**Sintering of the Tricalcium Phosphate-Titania-**

DOI: 10.5772/intechopen.68501

© 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.

interest for orthopedic and dental applications [1–13]. The tricalcium phosphate has been used clinically to repair bone defects for many years [6–10]. However, the major limitation of the β‐TCP uses as load bearing biomaterial is their poor fatigue resistance [6, 9]. Moreover, the mechanical properties of tricalcium phosphate are generally inadequate for many load‐carrying

**Magnesium Fluoride Composites**

**Magnesium Fluoride Composites**

Additional information is available at the end of the chapter

) and magnesium fluoride (MgF<sup>2</sup>

nia at 1200°C. Moreover, the addition of 4 wt% MgF<sup>2</sup>

(PO4 ) 2

leads to better mechanical properties (σr

obtained performances of the TCP‐TiO<sup>2</sup>

tissue and especially as enamel.

patibility of the β‐TCP and the high tribological properties of TiO<sup>2</sup>

30, 40 and 50 wt%) was studied. The performances of the TCP‐TiO<sup>2</sup>

Additional information is available at the end of the chapter

Ibticem Ayadi and Foued Ben Ayed

Ibticem Ayadi and Foued Ben Ayed

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

**Abstract**

(β‐Ca<sup>3</sup>

Titania (TiO2

phosphate

**1. Introduction**

Tricalcium phosphate (β‐Ca<sup>3</sup>

(PO4 ) 2

**Provisional chapter**

## **Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites Magnesium Fluoride Composites**

**Sintering of the Tricalcium Phosphate-Titania-**

DOI: 10.5772/intechopen.68501

Ibticem Ayadi and Foued Ben Ayed Additional information is available at the end of the chapter

Ibticem Ayadi and Foued Ben Ayed

Additional information is available at the end of the chapter

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

#### **Abstract**

Titania (TiO2 ) and magnesium fluoride (MgF<sup>2</sup> ) can be mixed with tricalcium phosphate (β‐Ca<sup>3</sup> (PO4 ) 2 , β‐TCP) to make bioceramic composites, which would combine the biocom‐ patibility of the β‐TCP and the high tribological properties of TiO<sup>2</sup> and MgF<sup>2</sup> for biomedi‐ cal applications. The samples were characterized by different characterization techniques such as physicochemical and mechanical. The sintering of the TCP at various tempera‐ tures (1000, 1100, 1200 and 1300°C) with different percentages of titania (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%) was studied. The performances of the TCP‐TiO<sup>2</sup> composites increase with both the sintering temperature and the amount of titania. The highest values of the composites' (H = 270 Hv; E = 33.1 GPa and G = 15.7 GPa) were obtained with 40 wt% tita‐ nia at 1200°C. Moreover, the addition of 4 wt% MgF<sup>2</sup> to the TCP‐40 wt% TiO<sup>2</sup> composites leads to better mechanical properties (σr = 27 MPa; H = 360 Hv; E = 51 GPa and G = 20 GPa) at 1200°C for 1 h. The amelioration of these properties is due to the formation of a new compound and the liquid phase which helps to fill the pores in the microstructure. The obtained performances of the TCP‐TiO<sup>2</sup> ‐MgF<sup>2</sup> composites are similar to those of bone tissue and especially as enamel.

**Keywords:** sintering, biomaterial, composites, mechanical properties, tricalcium phosphate

## **1. Introduction**

Tricalcium phosphate (β‐Ca<sup>3</sup> (PO4 ) 2 , β‐TCP) based on biomaterial has attracted considerable interest for orthopedic and dental applications [1–13]. The tricalcium phosphate has been used clinically to repair bone defects for many years [6–10]. However, the major limitation of the β‐TCP uses as load bearing biomaterial is their poor fatigue resistance [6, 9]. Moreover, the mechanical properties of tricalcium phosphate are generally inadequate for many load‐carrying

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. © 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.

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons

applications [3, 6, 9]. Hence, inert oxides such as alumina (Al<sup>2</sup> O3 ) or zirconia (ZrO<sup>2</sup> ) have been widely studied due to their bio‐inertness, excellent tribological properties, fracture toughness and strength [6, 9, 11–15]. Thus, the study conducted by Sakka et al. has recently concerned with the tricalcium phosphate/alumina system [11]. The researchers were interested in pro‐ ducing the Al<sup>2</sup> O3 ‐TCP composites with different percentages of β‐TCP (10, 20, 40 and 50 wt%). In fact, the best mechanical properties of this study reached 13.5 MPa with a composition of the alumina‐10 wt% tricalcium phosphate composite after the sintering at 1600°C [11]. Thereby, the next study directed by Sakka et al. proved that the incorporation of 5 wt% TiO<sup>2</sup> to the alu‐ mina‐10 wt% tricalcium phosphate composite matrix leads to the highest mechanical perfor‐ mances (74 MPa) at 1600°C [12]. However, the very low content of tricalcium phosphate (10 wt%) in the composites based on alumina, titania and tricalcium phosphate limits its use in the biomedical applications. Furthermore, the presence of the sizeable grains hinders the per‐ formances of these composites caused by the high temperature sintering [12]. Also, the study made by Sallemi et al. was interested to elaborate and to characterize the tricalcium phosphate‐ zirconia composites with different percentages of zirconia (25, 50 and 75 wt%) [14]. Thus, the ultimate values of the performances of the tricalcium phosphate were obtained with 50 wt% zirconia [14]. But, the experimental results in this study indicate the appearance of low mechanical properties of the tricalcium phosphate‐zirconia composites [13, 14]. Based on the results, the inverse allotropic transformation of zirconia is sufficient to cause the degradation of mechanical properties of these composites [13, 14]. Thus, the expansion in volume of the zirconia samples is responsible for the fragility of the tricalcium phosphate‐zirconia compos‐ ites [13, 14]. In conclusion, the addition of alumina or zirconia to the TCP matrix did not enhance the mechanical properties [11–14]. For this reason, we tested another inert oxide like titania (TiO2 ). Besides, titania has attracted much attention due to its excellent ability to chemi‐ cally bond with living hard tissue [16, 17]. This happens through the formation of a bone min‐ eral, such as tricalcium phosphate phase on the material surface that ultimately induces direct bonding with native bone tissue [18–20]. Thus, titania has been attracting attention as an implantable material because it is harmless to a living body, and with good mechanical prop‐ erties. Moreover, titania is an attractive material, applicable to various fields, such as biomedi‐ cal applications [16–20]. Furthermore, the titania is considered as having excellent biocompatibility, manifested in various biomedical applications [16–20]. So, the choice of the titania added to the tricalcium phosphate matrix was based on those considerations. Thereby, we would combine the biocompatibility of the TCP with the high mechanical and tribological properties of titania in order to elaborate a bioceramic composite. This combination of calcium phosphate/inert oxide system could give rise to more biomaterials in the physiological envi‐ ronment [21]. Few papers were interested on the study of the TCP‐TiO<sup>2</sup> composites using hydroxyapatite and titania as starting materials [22–24]. In fact, titania has been added to the tricalcium phosphate matrix in the order of enhancing the mechanical performances of the TCP and not degrading its biocompatibility. Thereby, this study focused on the sintering and the mechanical properties of the tricalcium phosphate with different percentages of TiO<sup>2</sup> (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%). In this study, the performances of the TCP‐TiO<sup>2</sup> composites increase with the augmentation of the sintering temperature and the amount of titania. Thus, the best mechanical properties resulting from this study were obtained with the TCP‐40 wt% TiO2 composites after the sintering process at 1200°C for 1 h. The optimum values of the

TCP‐40 wt% TiO<sup>2</sup>

ion (F−

tion of MgF<sup>2</sup>

tution of F−

in the HAp/ZrO<sup>2</sup>

for OH‐

of the TCP‐40 wt% TiO<sup>2</sup>

Brazilian test (σr

into TCP below 1400°C owing to the ion substitution F−

crystal structure [37]. Many studies show that, the addition of MgF<sup>2</sup>

cisely the enamel. Hence, MgF<sup>2</sup>

composites of mechanical strength, Vickers hardness, Young's and Shear

has been chosen as a suitable compound for doping calcium

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

composites by the inclusion

47

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

) are incorporated in

(5 and 10 vol%) added to

to the tetragonal zirconia

enhances

) and

in the hydroxyapatite (HAp)

modulus reached 33 MPa, 270 Hv, 33.1 GPa and 15.7 GPa, respectively. Thus, the aim of this

of several additives to reach mechanical property similar to those of the bone tissues and pre‐

phosphate [25]. Magnesium fluoride is an important material, which attracts attention thanks to its wide range of applications [26–32]. In the light of its technological properties, magnesium fluoride was studied in different area [26, 30, 33]. Among the alkaline‐earth fluorides, MgF<sup>2</sup> crystallizes in the rutile‐type structure [26, 27]. Thanks to the higher compressibility of MgF<sup>2</sup> and the relative sizes of its ions, the phase transformations of this material are found at lower

either the hydroxyapatite (HAp) or TCP crystal structure to replace the position of calcium, phosphorus or hydroxyl ion [25]. Fluoride is known to be important in suppressing dental car‐ ies [33, 34]. It stimulates the proliferation and the differentiation of bone cells [35]. The fluorine

) has been investigated as an essential element for bone and dental formation in the human body [35, 36]. Fluoride compound is known as an effective additive for reducing the phase decomposition of hydroxyapatite due to the crystal structure stabilization [37]. Magnesium is an abundant and essential cation in the human body, since it has significant effects on human metabolism [38]. Mg2+ promotes the dental caries formation caused by the high calcification process in the bone formation [37, 38]. Even though, the lack of Mg2+ ion results from the inhibition of bone growth, the degradation of bone structure and the decrease of osteoblast adhesion [37, 38]. Evis and Pinar Sun [34] demonstrate the β‐TCP structure stabi‐ lization and the hydroxyapatite formation by the heat treatment above 800°C. Hot pressed magnesium fluoride has been used as a dome material for many devices due to its excellent mechanical strength and thermal stability [39–41]. The study conducted by Kim et al. aimed to increase the mechanical properties and to inhibit the phase decomposition of HAp by the addi‐

composites [37]. As a result, the MgF<sup>2</sup>

for OH‐

composites sintered at 1200°C for different lengths of the sintering

), Vickers indentation and ultrasonic technique, the magic angle scanning

the HAp‐zirconia composites completely suppresses the decomposition of hydroxyapatite

and HAp mixtures strongly reduces the tendency of HAp decomposition due to partial substi‐

magnesium (Mg2+) ions. On the other hand, fluoride presents the much more stable ion in acidic environment (pH = 4). This property is very important in medical fields and especially in dentistry [44]. According to the previous studies, fluoride has a great influence on the physi‐ cal and biological properties of materials [45–47]. Additionally, fluoride is an essential trace element required for normal dental and skeletal development [47]. It has been shown that the presence of fluoride offers beneficial effects on increasing the quantity and quality of bone formation in the body [47, 48]. In the second time, we are interested in the examination of the effect of magnesium fluoride addition (1, 2.5, 3, 4, 4.5, 5, 6, 7.5 and 10 wt%) on the performances

time. Then, we will characterize the resulting composites by different techniques such as

the tricalcium phosphate‐titania composites performances by the inclusion of fluoride (F−

ions in the calcium phosphate structure [42, 43]. Therefore, MgF<sup>2</sup>

study is to ameliorate the performances of the TCP‐40 wt% TiO<sup>2</sup>

pressures than those in many oxides [32]. The minor ions (Mg2+ and F−

TCP‐40 wt% TiO<sup>2</sup> composites of mechanical strength, Vickers hardness, Young's and Shear modulus reached 33 MPa, 270 Hv, 33.1 GPa and 15.7 GPa, respectively. Thus, the aim of this study is to ameliorate the performances of the TCP‐40 wt% TiO<sup>2</sup> composites by the inclusion of several additives to reach mechanical property similar to those of the bone tissues and pre‐ cisely the enamel. Hence, MgF<sup>2</sup> has been chosen as a suitable compound for doping calcium phosphate [25]. Magnesium fluoride is an important material, which attracts attention thanks to its wide range of applications [26–32]. In the light of its technological properties, magnesium fluoride was studied in different area [26, 30, 33]. Among the alkaline‐earth fluorides, MgF<sup>2</sup> crystallizes in the rutile‐type structure [26, 27]. Thanks to the higher compressibility of MgF<sup>2</sup> and the relative sizes of its ions, the phase transformations of this material are found at lower pressures than those in many oxides [32]. The minor ions (Mg2+ and F− ) are incorporated in either the hydroxyapatite (HAp) or TCP crystal structure to replace the position of calcium, phosphorus or hydroxyl ion [25]. Fluoride is known to be important in suppressing dental car‐ ies [33, 34]. It stimulates the proliferation and the differentiation of bone cells [35]. The fluorine ion (F− ) has been investigated as an essential element for bone and dental formation in the human body [35, 36]. Fluoride compound is known as an effective additive for reducing the phase decomposition of hydroxyapatite due to the crystal structure stabilization [37]. Magnesium is an abundant and essential cation in the human body, since it has significant effects on human metabolism [38]. Mg2+ promotes the dental caries formation caused by the high calcification process in the bone formation [37, 38]. Even though, the lack of Mg2+ ion results from the inhibition of bone growth, the degradation of bone structure and the decrease of osteoblast adhesion [37, 38]. Evis and Pinar Sun [34] demonstrate the β‐TCP structure stabi‐ lization and the hydroxyapatite formation by the heat treatment above 800°C. Hot pressed magnesium fluoride has been used as a dome material for many devices due to its excellent mechanical strength and thermal stability [39–41]. The study conducted by Kim et al. aimed to increase the mechanical properties and to inhibit the phase decomposition of HAp by the addi‐ tion of MgF<sup>2</sup> in the HAp/ZrO<sup>2</sup> composites [37]. As a result, the MgF<sup>2</sup> (5 and 10 vol%) added to the HAp‐zirconia composites completely suppresses the decomposition of hydroxyapatite into TCP below 1400°C owing to the ion substitution F− for OH‐ in the hydroxyapatite (HAp) crystal structure [37]. Many studies show that, the addition of MgF<sup>2</sup> to the tetragonal zirconia and HAp mixtures strongly reduces the tendency of HAp decomposition due to partial substi‐ tution of F− for OH‐ ions in the calcium phosphate structure [42, 43]. Therefore, MgF<sup>2</sup> enhances the tricalcium phosphate‐titania composites performances by the inclusion of fluoride (F− ) and magnesium (Mg2+) ions. On the other hand, fluoride presents the much more stable ion in acidic environment (pH = 4). This property is very important in medical fields and especially in dentistry [44]. According to the previous studies, fluoride has a great influence on the physi‐ cal and biological properties of materials [45–47]. Additionally, fluoride is an essential trace element required for normal dental and skeletal development [47]. It has been shown that the presence of fluoride offers beneficial effects on increasing the quantity and quality of bone formation in the body [47, 48]. In the second time, we are interested in the examination of the effect of magnesium fluoride addition (1, 2.5, 3, 4, 4.5, 5, 6, 7.5 and 10 wt%) on the performances of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for different lengths of the sintering time. Then, we will characterize the resulting composites by different techniques such as Brazilian test (σr ), Vickers indentation and ultrasonic technique, the magic angle scanning

applications [3, 6, 9]. Hence, inert oxides such as alumina (Al<sup>2</sup>

ducing the Al<sup>2</sup>

46 Sintering of Functional Materials

titania (TiO2

TiO2

O3

O3

‐TCP composites with different percentages of β‐TCP (10, 20, 40 and 50 wt%).

). Besides, titania has attracted much attention due to its excellent ability to chemi‐

cally bond with living hard tissue [16, 17]. This happens through the formation of a bone min‐ eral, such as tricalcium phosphate phase on the material surface that ultimately induces direct bonding with native bone tissue [18–20]. Thus, titania has been attracting attention as an implantable material because it is harmless to a living body, and with good mechanical prop‐ erties. Moreover, titania is an attractive material, applicable to various fields, such as biomedi‐ cal applications [16–20]. Furthermore, the titania is considered as having excellent biocompatibility, manifested in various biomedical applications [16–20]. So, the choice of the titania added to the tricalcium phosphate matrix was based on those considerations. Thereby, we would combine the biocompatibility of the TCP with the high mechanical and tribological properties of titania in order to elaborate a bioceramic composite. This combination of calcium phosphate/inert oxide system could give rise to more biomaterials in the physiological envi‐

hydroxyapatite and titania as starting materials [22–24]. In fact, titania has been added to the tricalcium phosphate matrix in the order of enhancing the mechanical performances of the TCP and not degrading its biocompatibility. Thereby, this study focused on the sintering and the mechanical properties of the tricalcium phosphate with different percentages of TiO<sup>2</sup>

increase with the augmentation of the sintering temperature and the amount of titania. Thus, the best mechanical properties resulting from this study were obtained with the TCP‐40 wt%

composites after the sintering process at 1200°C for 1 h. The optimum values of the

ronment [21]. Few papers were interested on the study of the TCP‐TiO<sup>2</sup>

5, 7.5, 10, 20, 30, 40 and 50 wt%). In this study, the performances of the TCP‐TiO<sup>2</sup>

widely studied due to their bio‐inertness, excellent tribological properties, fracture toughness and strength [6, 9, 11–15]. Thus, the study conducted by Sakka et al. has recently concerned with the tricalcium phosphate/alumina system [11]. The researchers were interested in pro‐

In fact, the best mechanical properties of this study reached 13.5 MPa with a composition of the alumina‐10 wt% tricalcium phosphate composite after the sintering at 1600°C [11]. Thereby,

mina‐10 wt% tricalcium phosphate composite matrix leads to the highest mechanical perfor‐ mances (74 MPa) at 1600°C [12]. However, the very low content of tricalcium phosphate (10 wt%) in the composites based on alumina, titania and tricalcium phosphate limits its use in the biomedical applications. Furthermore, the presence of the sizeable grains hinders the per‐ formances of these composites caused by the high temperature sintering [12]. Also, the study made by Sallemi et al. was interested to elaborate and to characterize the tricalcium phosphate‐ zirconia composites with different percentages of zirconia (25, 50 and 75 wt%) [14]. Thus, the ultimate values of the performances of the tricalcium phosphate were obtained with 50 wt% zirconia [14]. But, the experimental results in this study indicate the appearance of low mechanical properties of the tricalcium phosphate‐zirconia composites [13, 14]. Based on the results, the inverse allotropic transformation of zirconia is sufficient to cause the degradation of mechanical properties of these composites [13, 14]. Thus, the expansion in volume of the zirconia samples is responsible for the fragility of the tricalcium phosphate‐zirconia compos‐ ites [13, 14]. In conclusion, the addition of alumina or zirconia to the TCP matrix did not enhance the mechanical properties [11–14]. For this reason, we tested another inert oxide like

the next study directed by Sakka et al. proved that the incorporation of 5 wt% TiO<sup>2</sup>

) or zirconia (ZrO<sup>2</sup>

) have been

to the alu‐

composites using

(2.5,

composites

nuclear magnetic resonance (31P), the scanning electron microscopy, the infrared spectroscopy and the X‐ray diffraction (XRD).

## **2. Materials and methods**

The β‐TCP powder resulted from a mixture of calcium carbonate (CaCO<sup>3</sup> : Fluka, purity ≥ 98.5%) and calcium phosphate dibasic anhydrous (CaHPO<sup>4</sup> : Fluka, purity ≥ 99%) after a heat treatment at 1000°C for 2 h according to the following reaction (Eq. (1))" [49]:

$$\text{2CaHPO}\_4\text{(s)} + \text{CaCO}\_3\text{(s)} \leftrightarrow \text{Ca}\_3\text{(PO}\_4\text{)} + \text{H}\_2\text{O} + \text{CO}\_2\text{(g)}\tag{1}$$

d = (3 . 07 A′ + 3 . 89 B′ + 3 . 15 C′

and MgF<sup>2</sup>

of β‐TCP, anatase‐TiO<sup>2</sup>

spectrometer.

with a gold layer.

 **(wt%) B′<sup>a</sup>**

A′, B′ and C′ are the weight rates of β‐TCP, TiO<sup>2</sup>

**Table 2.** The weight ratios of the different TCP‐TiO<sup>2</sup>

2010) [50].

**A′<sup>a</sup>**

a

bTheoretical density.

posite are illustrated in **Table 2**.

where A′, B′ and C′ are the weight ratios and 3.07, 3.89 and 3.15 are the theoretical densities

The powders were mixed with absolute ethanol in an agate mortar. After milling these pow‐ ders, the mixture was dried at 80°C for 24 h. After drying, the powder mixtures were molded in a cylinder with diameter 20 mm and thickness 4 mm and pressed under 150 MPa. The

The determination of the phase transformation in the microstructure of the elaborated composites was investigated by X‐ray diffraction (XRD) analysis. The identification of the components phases was done by means Seifert XRD 3000 TT diffractometer with CuKα radiation (λ = 1.54056 Å). The phase identification was operated resulting from the compar‐ ison between experimental XRD patterns and standards files compiled by the International Center for Diffraction Data (ICDD). The powders were then characterized by infrared spectrometric analysis with attenuated total reflection method (ATR) (Agilent Cary 630 Fourier Transform Infrared Spectrometer (FTIR)). The powders were submitted to the 31P magic angle scanning nuclear magnetic resonance (MAS‐NMR) on a Brucker 300 WB

The microstructure of the fractured surfaces of the sintered samples was investigated with scanning electron microscope (SEM) (JEOL JSM 5800LV) after enhancing their conductivity

The powder's size was measured through a Micromeritics Sedigraph 5000. The specific sur‐

 **(wt%) db**

face area (SSA) was determined by the BET method with azote (N<sup>2</sup>

 **(wt%) C′<sup>a</sup>**

59.50 39.50 1.00 3.3947 58.75 38.75 2.50 3.3897 58.50 38.50 3.00 3.3881 58.00 38.00 4.00 3.3848 57.75 37.75 4.50 3.3831 57.50 37.50 5.00 3.3815 57.00 37.00 6.00 3.3782 56.25 36.25 7.50 3.3732 55.00 35.00 10.00 3.3650

and MgF<sup>2</sup>

‐MgF<sup>2</sup>

, respectively.

composites.

specimens were heated and cooled at rates of 10°C min−1 and 20°C min−1, respectively.

) /100 (3)

49

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

) as the adsorbed gas (ASAP

, respectively. The calculated theoretical densities of all‐com‐

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

The initial powders used to obtain the TCP‐TiO<sup>2</sup> ‐MgF<sup>2</sup> composites were synthesized trical‐ cium phosphate (β‐TCP), magnesium fluoride MgF<sup>2</sup> (Sigma Aldrich, purity > 98%) and titania TiO2 (Fluka, purity > 98%). Firstly, Titania was introduced in the β‐TCP matrix at different contents (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%) followed by homogeneous mixing in a mortar. The heat treatment of the compacted disks was carried out in a vertical programmable muffle furnace (Pyrox 2408) at various temperatures (1000, 1100, 1200 and 1300°C) for 1 h. The theo‐ retical density (d) was determined using the following Eq. (2):

$$\mathbf{d} = (3..07\mathbf{A} + 3..89\mathbf{B} \,)/100\tag{2}$$

where A and B are the weight ratios and 3.07 and 3.89 are the theoretical densities of β‐TCP and anatase‐TiO<sup>2</sup> , respectively. The calculated theoretical densities of each composite are illustrated in **Table 1**.

Different amounts of MgF<sup>2</sup> (1, 2.5, 3, 4, 4.5, 5, 6, 7.5 and 10 wt%) were added to the TCP‐40 wt% TiO2 composites followed by homogeneous mixing in a mortar. The heat treatment of the compacted specimens was carried out in a vertical programmable muffle furnace (Pyrox 2408) at 1200°C for different sintering times. The theoretical density (d) was determined using the following Eq. (3):


a A and B are the weight rates of TiO<sup>2</sup> and β‐TCP, respectively.

**Table 1.** The weight ratios of different TCP‐TiO<sup>2</sup> composites.

b Theoretical density.

$$\mathbf{d} = \langle \mathbf{3} \ .07\,\mathrm{A}^{\prime} + \mathbf{3} \ .89\,\mathrm{B}^{\prime} + \mathbf{3} \ .15\,\mathrm{C}^{\prime} \rangle / 100 \tag{3}$$

where A′, B′ and C′ are the weight ratios and 3.07, 3.89 and 3.15 are the theoretical densities of β‐TCP, anatase‐TiO<sup>2</sup> and MgF<sup>2</sup> , respectively. The calculated theoretical densities of all‐com‐ posite are illustrated in **Table 2**.

nuclear magnetic resonance (31P), the scanning electron microscopy, the infrared spectroscopy

(s)↔C a3

 d = (3 . 07A + 3 . 89B ) /100 (2) where A and B are the weight ratios and 3.07 and 3.89 are the theoretical densities of β‐TCP

 composites followed by homogeneous mixing in a mortar. The heat treatment of the compacted specimens was carried out in a vertical programmable muffle furnace (Pyrox 2408) at 1200°C for different sintering times. The theoretical density (d) was determined using the

 **(wt%) db**

 (Fluka, purity > 98%). Firstly, Titania was introduced in the β‐TCP matrix at different contents (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%) followed by homogeneous mixing in a mortar. The heat treatment of the compacted disks was carried out in a vertical programmable muffle furnace (Pyrox 2408) at various temperatures (1000, 1100, 1200 and 1300°C) for 1 h. The theo‐

(P O4

‐MgF<sup>2</sup>

, respectively. The calculated theoretical densities of each composite are

(1, 2.5, 3, 4, 4.5, 5, 6, 7.5 and 10 wt%) were added to the TCP‐40 wt%

) +H<sup>2</sup>

O +CO2

: Fluka, purity ≥

: Fluka, purity ≥ 99%) after a heat

composites were synthesized trical‐

(Sigma Aldrich, purity > 98%) and titania

(g) (1)

The β‐TCP powder resulted from a mixture of calcium carbonate (CaCO<sup>3</sup>

treatment at 1000°C for 2 h according to the following reaction (Eq. (1))" [49]:

(s) +CaCO3

98.5%) and calcium phosphate dibasic anhydrous (CaHPO<sup>4</sup>

retical density (d) was determined using the following Eq. (2):

2.50 97.50 3.0905 5.00 95.00 3.1110 7.50 92.50 3.1315 10.00 90.00 3.1520 20.00 80.00 3.2340 30.00 70.00 3.3160 40.00 60.00 3.3980 50.00 50.00 3.4800

and β‐TCP, respectively.

composites.

The initial powders used to obtain the TCP‐TiO<sup>2</sup>

cium phosphate (β‐TCP), magnesium fluoride MgF<sup>2</sup>

and the X‐ray diffraction (XRD).

48 Sintering of Functional Materials

**2. Materials and methods**

2CaHPO4

TiO2

TiO2

**Aa**

a

b

and anatase‐TiO<sup>2</sup>

following Eq. (3):

illustrated in **Table 1**.

Different amounts of MgF<sup>2</sup>

 **(wt%) Ba**

A and B are the weight rates of TiO<sup>2</sup>

**Table 1.** The weight ratios of different TCP‐TiO<sup>2</sup>

Theoretical density.

The powders were mixed with absolute ethanol in an agate mortar. After milling these pow‐ ders, the mixture was dried at 80°C for 24 h. After drying, the powder mixtures were molded in a cylinder with diameter 20 mm and thickness 4 mm and pressed under 150 MPa. The specimens were heated and cooled at rates of 10°C min−1 and 20°C min−1, respectively.

The determination of the phase transformation in the microstructure of the elaborated composites was investigated by X‐ray diffraction (XRD) analysis. The identification of the components phases was done by means Seifert XRD 3000 TT diffractometer with CuKα radiation (λ = 1.54056 Å). The phase identification was operated resulting from the compar‐ ison between experimental XRD patterns and standards files compiled by the International Center for Diffraction Data (ICDD). The powders were then characterized by infrared spectrometric analysis with attenuated total reflection method (ATR) (Agilent Cary 630 Fourier Transform Infrared Spectrometer (FTIR)). The powders were submitted to the 31P magic angle scanning nuclear magnetic resonance (MAS‐NMR) on a Brucker 300 WB spectrometer.

The microstructure of the fractured surfaces of the sintered samples was investigated with scanning electron microscope (SEM) (JEOL JSM 5800LV) after enhancing their conductivity with a gold layer.

The powder's size was measured through a Micromeritics Sedigraph 5000. The specific sur‐ face area (SSA) was determined by the BET method with azote (N<sup>2</sup> ) as the adsorbed gas (ASAP 2010) [50].


a A′, B′ and C′ are the weight rates of β‐TCP, TiO<sup>2</sup> and MgF<sup>2</sup> , respectively. bTheoretical density.

**Table 2.** The weight ratios of the different TCP‐TiO<sup>2</sup> ‐MgF<sup>2</sup> composites. The main particle size (DBET) was calculated by assuming that the primary particles are spheri‐ cal (Eq. (4)) [51]:

$$\mathbf{D}\_{\text{net}} = \frac{6}{\text{Sp}} \tag{4}$$

at 1470°C is relative to the second transformation of TCP (α–α′) (**Figure 1a**). These results are mentioned in previous works from literature [12, 56]. The Differential thermal analysis (DTA)

of two endothermic peaks (**Figure 1c**). The first peak at 1470°C is attributed to the second transformation of the β‐TCP (α–α′) while the new peak at 1440°C is probably relative to the

**/g) DBET (μm) D50**

TiO2 12.00 0.11 0.20 3.89 β‐TCP 0.80 2.80 6.00 3.07 MgF<sup>2</sup> 37.00 0.05 0.10 3.15

liquid phase formation. Furthermore, the melting points of both β‐TCP and TiO<sup>2</sup>

and it is not relative to the melting point of the initial powders. The results are similar to that

and (c) TCP‐10 wt% TiO<sup>2</sup>

composites.

and 1855°C, respectively [22]. Thus, the liquid phase is formed between TiO<sup>2</sup>

found in the previous studies [12, 22]. They show that in the TiO<sup>2</sup>

**Figure 1.** DTA curves of: (a) β‐TCP, (b) anatase‐TiO<sup>2</sup>

shows that no evolution can be obtained with the sintering temperature

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

**a**

 **(μm) db**

composites reveals the appearance

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

51

are 1756

and β‐TCP,

/β‐TCP system, there is an

curve of the TiO<sup>2</sup>

a

b

Mean diameter.

Theoretical density.

**Compounds SSA (m2**

(**Figure 1b**). The thermogram of the TCP‐10 wt% TiO<sup>2</sup>

**Table 3.** Characteristics of the different powders used in the study.

where ρ is the theoretical density and S is the surface specific area.

The densification of the sintered specimens was evaluated by the measurements of the speci‐ men dimensions. The relative error of the densification value was about 1%.

Thermal expansion‐shrinkage of the compact powder was measured with a dilatometer (Model Setaram TMA 92 dilatometer). About 20 mg of powder was heated with a heating rate of 10°C min−1 along with alumina as a reference material and a low stream of argon gas. The heating and cooling rates were 10 and 20°C min−1, respectively.

The Brazilian test was used to measure the mechanical strength of the samples [52, 53]. The rupture strength (σr ) (or mechanical strength) was calculated based on the maximum applied load recorded (P) and sample's dimensions (D being the diameter and t being the thickness) using (Eq. (5)):

$$
\sigma\_v = \frac{2 \times \text{P}}{\pi \times \text{D} \times t} \tag{5}
$$

The sintered samples were examined by Vickers test indentation using loading values of 49 or 98 N applied for 15 s. This test was investigated after polishment of the samples surfaces between 1 and 3 µm with diamond paste. Vickers hardness value (H) was determined by Eq. (6) [54]:

$$\mathbf{H} = \mathbf{1} \, . 854 \mathbf{(P/d^2)}\tag{6}$$

where H is Vickers hardness, 'd' is the indent diagonal length and P is the indentation load.

The characterization of the samples using ultrasound technique promotes the determina‐ tion of both Young's and Shear moduli value [55]. The measurement was done by a high‐fre‐ quency generator Model 5077PR (Olympus). Young's modulus and the Shear modulus were calculated from the point of the longitudinal and the transversal ultrasonic velocities [55].

## **3. Results**

#### **3.1. Characterization of the starting powders**

**Table 3** presents the particles size distribution data (measured by granulometric repartition), the SSA and the calculated average grain sizes (DBET) of the β‐TCP, TiO<sup>2</sup> and MgF<sup>2</sup> powders. The averages of the grain's size of the samples obtained through the SSA (DBET) and the granu‐ lometric repartition (D50) are different. This difference may be due to the presence of agglom‐ erates in the initial powders (**Table 3**).

The curve of the differential thermal analysis of the β‐TCP shows the presence of two endo‐ thermic peaks: the first large band located between 1230 and 1270°C is attributed to the first allotropic transformation of the tricalcium phosphate (β–α) and the second peak registered at 1470°C is relative to the second transformation of TCP (α–α′) (**Figure 1a**). These results are mentioned in previous works from literature [12, 56]. The Differential thermal analysis (DTA) curve of the TiO<sup>2</sup> shows that no evolution can be obtained with the sintering temperature (**Figure 1b**). The thermogram of the TCP‐10 wt% TiO<sup>2</sup> composites reveals the appearance of two endothermic peaks (**Figure 1c**). The first peak at 1470°C is attributed to the second transformation of the β‐TCP (α–α′) while the new peak at 1440°C is probably relative to the


b Theoretical density.

The main particle size (DBET) was calculated by assuming that the primary particles are spheri‐

The densification of the sintered specimens was evaluated by the measurements of the speci‐

Thermal expansion‐shrinkage of the compact powder was measured with a dilatometer (Model Setaram TMA 92 dilatometer). About 20 mg of powder was heated with a heating rate of 10°C min−1 along with alumina as a reference material and a low stream of argon gas. The

The Brazilian test was used to measure the mechanical strength of the samples [52, 53]. The

load recorded (P) and sample's dimensions (D being the diameter and t being the thickness)

The sintered samples were examined by Vickers test indentation using loading values of 49 or 98 N applied for 15 s. This test was investigated after polishment of the samples surfaces between 1 and 3 µm with diamond paste. Vickers hardness value (H) was determined by

 H = 1 . 854(P/d2 ) (6) where H is Vickers hardness, 'd' is the indent diagonal length and P is the indentation load. The characterization of the samples using ultrasound technique promotes the determina‐ tion of both Young's and Shear moduli value [55]. The measurement was done by a high‐fre‐ quency generator Model 5077PR (Olympus). Young's modulus and the Shear modulus were calculated from the point of the longitudinal and the transversal ultrasonic velocities [55].

**Table 3** presents the particles size distribution data (measured by granulometric repartition),

The averages of the grain's size of the samples obtained through the SSA (DBET) and the granu‐ lometric repartition (D50) are different. This difference may be due to the presence of agglom‐

The curve of the differential thermal analysis of the β‐TCP shows the presence of two endo‐ thermic peaks: the first large band located between 1230 and 1270°C is attributed to the first allotropic transformation of the tricalcium phosphate (β–α) and the second peak registered

the SSA and the calculated average grain sizes (DBET) of the β‐TCP, TiO<sup>2</sup>

) (or mechanical strength) was calculated based on the maximum applied

Sρ (4)

<sup>π</sup> <sup>×</sup> <sup>D</sup> <sup>×</sup> <sup>t</sup> (5)

and MgF<sup>2</sup>

powders.

cal (Eq. (4)) [51]:

50 Sintering of Functional Materials

rupture strength (σr

using (Eq. (5)):

Eq. (6) [54]:

**3. Results**

DBET <sup>=</sup> \_\_\_6

where ρ is the theoretical density and S is the surface specific area.

heating and cooling rates were 10 and 20°C min−1, respectively.

σr <sup>=</sup> \_\_\_\_\_\_\_ <sup>2</sup> <sup>×</sup> <sup>P</sup>

**3.1. Characterization of the starting powders**

erates in the initial powders (**Table 3**).

men dimensions. The relative error of the densification value was about 1%.

**Table 3.** Characteristics of the different powders used in the study.

**Figure 1.** DTA curves of: (a) β‐TCP, (b) anatase‐TiO<sup>2</sup> and (c) TCP‐10 wt% TiO<sup>2</sup> composites.

liquid phase formation. Furthermore, the melting points of both β‐TCP and TiO<sup>2</sup> are 1756 and 1855°C, respectively [22]. Thus, the liquid phase is formed between TiO<sup>2</sup> and β‐TCP, and it is not relative to the melting point of the initial powders. The results are similar to that found in the previous studies [12, 22]. They show that in the TiO<sup>2</sup> /β‐TCP system, there is an

**Figure 2.** Phase equilibrium diagram of the TiO<sup>2</sup> ‐Ca<sup>3</sup> (PO4 )2 system [22].

eutectic with a composition of 63 wt% TCP‐37 wt% TiO<sup>2</sup> at 1380°C [22]. **Figure 2** confirms the presence of the liquid phase which is showed in the binary diagram of the Ca<sup>3</sup> (PO4 )2 / TiO2 [22].

**3.2. Effect of the addition of titania on the densification and the mechanical properties of** 

, (c) MgF<sup>2</sup>

The evolution of the densification of the tricalcium phosphate was studied with the addi‐ tion of titania between 1000 and 1300°C. **Figure 4** shows the typical relationship between temperature and density. The density of the β‐TCP sintered with different percentages of the

and the sintering temperature (**Figure 5**). At 1200°C, the rupture strength of the TCP‐40 wt%

The evolution of Vickers hardness with different percentages of titania at various tempera‐ tures was shown in **Figure 6**. Vickers hardness reached its optimum value (270 Hv) at 1200°C

(**Figure 5h**). The amelioration of these performances of the TCP‐40 wt% TiO<sup>2</sup>

composites reached its maximum value (33 MPa) (**Figure 5g**). Above 1200°C, the rupture

 (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%) increases with the sintering temperature (**Figure 4**). At 1200°C, the optimum value of the densification (89%) was obtained with 40 wt% TiO<sup>2</sup> (**Figure 4g**). Above 1200°C, the performances of the composites decrease abruptly (**Figure 4**). **Figure 5** shows the influence of the titania additive (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%) at various sintering temperatures (1000, 1100, 1200 and 1300°C) on the rupture strength of TCP.

composites improves with both the content of TiO<sup>2</sup>

in the tricalcium phosphate matrix

composites.

and β‐TCP.

composites

composites was hindered abruptly (**Figure 5**). The discrepancy of

and (d) TCP‐10 wt% TiO<sup>2</sup>

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

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

53

. Then, Vickers hardness value decreases with the sintering temperature.

**the tricalcium phosphate**

strength of the TCP‐TiO<sup>2</sup>

with 40 wt% TiO<sup>2</sup>

The mechanical strength of the TCP‐TiO<sup>2</sup>

**Figure 3.** Linear shrinkage curves of: (a) β‐TCP, (b) TiO<sup>2</sup>

results appears especially after the addition of 50 wt% TiO<sup>2</sup>

could be according to the formation of a liquid phase between TiO<sup>2</sup>

TiO2

TiO2

**Figure 3** shows the dilatometric measurements of different powders used in this study (β‐TCP, TiO<sup>2</sup> , TCP‐10 wt% TiO<sup>2</sup> composites and MgF<sup>2</sup> ). The sintering temperature of the initial powder began at about 1000, 900 and 800°C for the β‐TCP, the TiO<sup>2</sup> and the MgF<sup>2</sup> , respectively (**Figure 3a**–**c**). The peak at 1230°C is attributed to the first allotropic transfor‐ mation of the tricalcium phosphate (**Figure 3a**). The shrinkage curve of the titania powder reveals one peak which is relative to the phase transformation from anatase to rutile at 1090°C (**Figure 3b**). This result was well‐confirmed by literature [57]. In fact, they showed that the commercial anatase‐titania powder has been transformed into the rutile structure at about 1000°C [57], which confirms our results. The curve of the pure MgF<sup>2</sup> indicates that the shrinkage was manifested at 800°C until 1150°C (**Figure 3c**). In fact, the rate of maximum densification corresponds to the inflection point, which is obtained at 890°C (**Figure 3c**). **Figure 3d** presents the shrinkage curve of the TCP‐10 wt% TiO<sup>2</sup> composites. In fact, no evolution was reported with the TCP‐10 wt% TiO<sup>2</sup> composites (**Figure 3d**). As a result, the content of 10 wt% TiO<sup>2</sup> stabilizes the TCP structure and prevents the inverse allotropic transformation of TCP from the α phase to the β phase. The stabilization was well‐explained by the exchange of Ca2+ and Ti4+ ions derived from β‐TCP and TiO<sup>2</sup> .

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites http://dx.doi.org/10.5772/intechopen.68501 53

**Figure 3.** Linear shrinkage curves of: (a) β‐TCP, (b) TiO<sup>2</sup> , (c) MgF<sup>2</sup> and (d) TCP‐10 wt% TiO<sup>2</sup> composites.

eutectic with a composition of 63 wt% TCP‐37 wt% TiO<sup>2</sup>

, TCP‐10 wt% TiO<sup>2</sup>

**Figure 2.** Phase equilibrium diagram of the TiO<sup>2</sup>

a result, the content of 10 wt% TiO<sup>2</sup>

TiO2

[22].

52 Sintering of Functional Materials

(β‐TCP, TiO<sup>2</sup>

the presence of the liquid phase which is showed in the binary diagram of the Ca<sup>3</sup>

**Figure 3** shows the dilatometric measurements of different powders used in this study

system [22].

respectively (**Figure 3a**–**c**). The peak at 1230°C is attributed to the first allotropic transfor‐ mation of the tricalcium phosphate (**Figure 3a**). The shrinkage curve of the titania powder reveals one peak which is relative to the phase transformation from anatase to rutile at 1090°C (**Figure 3b**). This result was well‐confirmed by literature [57]. In fact, they showed that the commercial anatase‐titania powder has been transformed into the rutile structure

that the shrinkage was manifested at 800°C until 1150°C (**Figure 3c**). In fact, the rate of maximum densification corresponds to the inflection point, which is obtained at 890°C

allotropic transformation of TCP from the α phase to the β phase. The stabilization was

composites and MgF<sup>2</sup>

initial powder began at about 1000, 900 and 800°C for the β‐TCP, the TiO<sup>2</sup>

‐Ca<sup>3</sup> (PO4 )2

at about 1000°C [57], which confirms our results. The curve of the pure MgF<sup>2</sup>

(**Figure 3c**). **Figure 3d** presents the shrinkage curve of the TCP‐10 wt% TiO<sup>2</sup>

well‐explained by the exchange of Ca2+ and Ti4+ ions derived from β‐TCP and TiO<sup>2</sup>

In fact, no evolution was reported with the TCP‐10 wt% TiO<sup>2</sup>

at 1380°C [22]. **Figure 2** confirms

). The sintering temperature of the

(PO4 )2 /

indicates

composites.

.

composites (**Figure 3d**). As

stabilizes the TCP structure and prevents the inverse

,

and the MgF<sup>2</sup>

#### **3.2. Effect of the addition of titania on the densification and the mechanical properties of the tricalcium phosphate**

The evolution of the densification of the tricalcium phosphate was studied with the addi‐ tion of titania between 1000 and 1300°C. **Figure 4** shows the typical relationship between temperature and density. The density of the β‐TCP sintered with different percentages of the TiO2 (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%) increases with the sintering temperature (**Figure 4**). At 1200°C, the optimum value of the densification (89%) was obtained with 40 wt% TiO<sup>2</sup> (**Figure 4g**). Above 1200°C, the performances of the composites decrease abruptly (**Figure 4**).

**Figure 5** shows the influence of the titania additive (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%) at various sintering temperatures (1000, 1100, 1200 and 1300°C) on the rupture strength of TCP. The mechanical strength of the TCP‐TiO<sup>2</sup> composites improves with both the content of TiO<sup>2</sup> and the sintering temperature (**Figure 5**). At 1200°C, the rupture strength of the TCP‐40 wt% TiO2 composites reached its maximum value (33 MPa) (**Figure 5g**). Above 1200°C, the rupture strength of the TCP‐TiO<sup>2</sup> composites was hindered abruptly (**Figure 5**). The discrepancy of results appears especially after the addition of 50 wt% TiO<sup>2</sup> in the tricalcium phosphate matrix (**Figure 5h**). The amelioration of these performances of the TCP‐40 wt% TiO<sup>2</sup> composites could be according to the formation of a liquid phase between TiO<sup>2</sup> and β‐TCP.

The evolution of Vickers hardness with different percentages of titania at various tempera‐ tures was shown in **Figure 6**. Vickers hardness reached its optimum value (270 Hv) at 1200°C with 40 wt% TiO<sup>2</sup> . Then, Vickers hardness value decreases with the sintering temperature.

**Figure 4.** Relative density versus temperature of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup> : (a) 2.5 wt%, (b) 5 wt%, (c) 7.5 wt%, (d) 10 wt%, (e) 20 wt%, (f) 30 wt%, (g) 40 wt% and (h) 50 wt%.

**Figures 7** and **8** present the evolution of elastics moduli (E and G) of the β‐TCP sintered with different contents of TiO<sup>2</sup> (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%) at various temperatures (1000, 1100, 1200 and 1300°C). The optimum values of both Young's modulus and Shear modulus reached 33.1 and 15.7 GPa, respectively, with 40 wt% TiO<sup>2</sup> . Beyond 40 wt% TiO<sup>2</sup> , the properties of the composites were hold up with the increase of the sintering temperature (**Figures 7** and **8**).

MgF<sup>2</sup>

full densification.

with 10 wt% MgF<sup>2</sup>

4 wt% of MgF<sup>2</sup>

, then the density of the composites decreases abruptly (**Figure 9**). The optimum value

**Figure 5.** Mechanical strength versus temperature of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup>

(a) 2.5 wt%, (b) 5 wt%, (c) 7.5 wt%, (d) 10 wt%, (e) 20 wt%, (f) 30 wt%, (g) 40 wt% and (h) 50 wt%.

rupture strength reached its maximum value (27 MPa) after the addition of 4 wt% MgF<sup>2</sup>

The influence of the magnesium fluoride addition on Vickers hardness was studied at 1200°C

(**Figure 11**). The addition of MgF<sup>2</sup>

to the composites led to a maximum of Vickers hardness (360 Hv) after the

. This result is obviously close to the

(1, 2.5, 3, 4, 4.5, 5, 6, 7.5 and 10 wt%). Thus, the

composites (**Figure 11**). Thus, the incorporation of

, the rupture strength decreases sharply (**Figure 10**). A remark‐

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composites sintered at 1200°C

:

55

composites was obtained

outstandingly enhanced

of the densification was about 94% with 4 wt% MgF<sup>2</sup>

for 1 h with different percentages of MgF<sup>2</sup>

(**Figure 10**).

(**Figure 10**). Above 4 wt% MgF<sup>2</sup>

with the same percentages of MgF<sup>2</sup>

Vickers hardness of the TCP‐40 wt% TiO<sup>2</sup>

**Figure 10** shows the rupture strength of the TCP‐40 wt% TiO<sup>2</sup>

able amelioration of the rupture strength of the TCP‐40 wt% TiO<sup>2</sup>

#### **3.3. Effect of the magnesium fluoride addition on the densification and the mechanical properties of the tricalcium phosphate‐titania composites**

The effect of the magnesium fluoride addition on the performances of the TCP‐40 wt% TiO2 composites has been assessed by the measurement of the density and the rupture strength. The densification behavior of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h with different amounts of MgF<sup>2</sup> (1, 2.5, 3, 4, 4.5, 5, 6, 7.5 and 10 wt%) is reported in **Figure 9**. These results show a significant improvement of the relative density of the TCP‐ TiO2 composites as a function of the magnesium fluoride added (**Figure 9**). The density of the TCP‐TiO<sup>2</sup> composites increases from 3 wt% MgF<sup>2</sup> and remains constant until 6 wt% Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites http://dx.doi.org/10.5772/intechopen.68501 55

**Figure 5.** Mechanical strength versus temperature of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup> : (a) 2.5 wt%, (b) 5 wt%, (c) 7.5 wt%, (d) 10 wt%, (e) 20 wt%, (f) 30 wt%, (g) 40 wt% and (h) 50 wt%.

**Figures 7** and **8** present the evolution of elastics moduli (E and G) of the β‐TCP sintered with

**Figure 4.** Relative density versus temperature of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup>

(a) 2.5 wt%, (b) 5 wt%, (c) 7.5 wt%, (d) 10 wt%, (e) 20 wt%, (f) 30 wt%, (g) 40 wt% and (h) 50 wt%.

1100, 1200 and 1300°C). The optimum values of both Young's modulus and Shear modulus

of the composites were hold up with the increase of the sintering temperature (**Figures 7** and **8**).

The effect of the magnesium fluoride addition on the performances of the TCP‐40 wt%

**Figure 9**. These results show a significant improvement of the relative density of the TCP‐

composites has been assessed by the measurement of the density and the rupture

composites as a function of the magnesium fluoride added (**Figure 9**). The density of

**3.3. Effect of the magnesium fluoride addition on the densification and the mechanical** 

(2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%) at various temperatures (1000,

. Beyond 40 wt% TiO<sup>2</sup>

(1, 2.5, 3, 4, 4.5, 5, 6, 7.5 and 10 wt%) is reported in

, the properties

:

composites sintered at 1200°C

and remains constant until 6 wt%

different contents of TiO<sup>2</sup>

54 Sintering of Functional Materials

TiO2

TiO2

the TCP‐TiO<sup>2</sup>

reached 33.1 and 15.7 GPa, respectively, with 40 wt% TiO<sup>2</sup>

**properties of the tricalcium phosphate‐titania composites**

strength. The densification behavior of the TCP‐40 wt% TiO<sup>2</sup>

composites increases from 3 wt% MgF<sup>2</sup>

for 1 h with different amounts of MgF<sup>2</sup>

MgF<sup>2</sup> , then the density of the composites decreases abruptly (**Figure 9**). The optimum value of the densification was about 94% with 4 wt% MgF<sup>2</sup> . This result is obviously close to the full densification.

**Figure 10** shows the rupture strength of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h with different percentages of MgF<sup>2</sup> (1, 2.5, 3, 4, 4.5, 5, 6, 7.5 and 10 wt%). Thus, the rupture strength reached its maximum value (27 MPa) after the addition of 4 wt% MgF<sup>2</sup> (**Figure 10**). Above 4 wt% MgF<sup>2</sup> , the rupture strength decreases sharply (**Figure 10**). A remark‐ able amelioration of the rupture strength of the TCP‐40 wt% TiO<sup>2</sup> composites was obtained with 10 wt% MgF<sup>2</sup> (**Figure 10**).

The influence of the magnesium fluoride addition on Vickers hardness was studied at 1200°C with the same percentages of MgF<sup>2</sup> (**Figure 11**). The addition of MgF<sup>2</sup> outstandingly enhanced Vickers hardness of the TCP‐40 wt% TiO<sup>2</sup> composites (**Figure 11**). Thus, the incorporation of 4 wt% of MgF<sup>2</sup> to the composites led to a maximum of Vickers hardness (360 Hv) after the

**Figure 6.** Vickers hardness versus temperature of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup> at: (a) 1000°C; (b) 1100°C; (c) 1200°C and (d) 1300°C.

sintering process at 1200°C for 1 h. Beyond 4 wt% MgF<sup>2</sup> , Vickers hardness of the composites was hindered (**Figure 11**). However, Vickers hardness of the TCP‐40 wt% TiO<sup>2</sup> composites increases slowly with 10 wt% MgF<sup>2</sup> (**Figure 11**).

composites was enhanced after the sintering process at 1200°C for 1 h and reached 27 MPa then, decreases abruptly (**Figure 14**). This amelioration was associated to the important densification

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posites with the length of the sintering time. Vickers hardness increases and reaches its opti‐ mum value (360 Hv) for 1 h (**Figure 15**). Then, a remarkable diminution of Vickers hardness

ites sintered for different length of the sintering time at 1200°C is pictured in **Figure 16A** and **B**, respectively. The sintering time has no effect on the elastic moduli (**Figure 16A** and **B**). Young's modulus reached 51 GPa as maximum value while Shear modulus reached 20 GPa as

After the sintering process, the samples were investigated by different characterization tech‐

 **composites**

**‐MgF2**

‐4 wt% MgF<sup>2</sup>

at: (a) 1000°C; (b) 1100°C, (c)

‐4 wt% MgF<sup>2</sup>

com‐

compos‐

**Figure 15** shows the evolution of Vickers hardness of the TCP‐38 wt% TiO<sup>2</sup>

**Figure 7.** Young's modulus of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup>

The evolution of the elastic moduli (E and G) of the TCP‐38 wt% TiO<sup>2</sup>

of the samples was registered above 1 h (**Figure 15**).

niques such as: XRD, IR, 31P MAS‐NMR and SEM.

an optimum value (**Figure 16A** and **B**).

**3.4. Characterization of the TCP‐TiO2**

(**Figures 13** and **14**).

1200°C and (d) 1300°C.

The evolution of the elastic modulus (E and G) of the TCP‐40 wt% TiO<sup>2</sup> composites sintered with different contents of MgF<sup>2</sup> is shown in **Figure 12A** and **B**, respectively. The elastic modu‐ lus of the TCP‐40 wt% TiO<sup>2</sup> composites increase with the addition of MgF<sup>2</sup> (**Figure 12A** and **B**). Thus, the optimum values of Young's modulus and the Shear modulus of the TCP‐40 wt% TiO2 composites are obtained by adding 4 wt% MgF<sup>2</sup> and these values reach 51 and 20 GPa, respectively. The performances of the samples were hindered with the increase of the percent‐ ages of MgF<sup>2</sup> in the TCP‐40 wt% TiO<sup>2</sup> composites (**Figure 12A** and **B**).

**Figure 13** illustrates a typical relation between the different lengths of the sintering time and the densification of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites after the sintering tempera‐ ture at 1200°C. The maximum value of densification (94%) was registered after the sintering process at 1200°C for 1 h (**Figure 13**). The densification's value decreases with the sintering time (**Figure 13**).

**Figure 14** depicts the rupture strength evolution of the samples after the sintering process at 1200°C for different length of the sintering time. The mechanical strength of the TCP‐TiO<sup>2</sup> ‐MgF<sup>2</sup>

**Figure 7.** Young's modulus of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup> at: (a) 1000°C; (b) 1100°C, (c) 1200°C and (d) 1300°C.

composites was enhanced after the sintering process at 1200°C for 1 h and reached 27 MPa then, decreases abruptly (**Figure 14**). This amelioration was associated to the important densification (**Figures 13** and **14**).

**Figure 15** shows the evolution of Vickers hardness of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> com‐ posites with the length of the sintering time. Vickers hardness increases and reaches its opti‐ mum value (360 Hv) for 1 h (**Figure 15**). Then, a remarkable diminution of Vickers hardness of the samples was registered above 1 h (**Figure 15**).

The evolution of the elastic moduli (E and G) of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> compos‐ ites sintered for different length of the sintering time at 1200°C is pictured in **Figure 16A** and **B**, respectively. The sintering time has no effect on the elastic moduli (**Figure 16A** and **B**). Young's modulus reached 51 GPa as maximum value while Shear modulus reached 20 GPa as an optimum value (**Figure 16A** and **B**).

#### **3.4. Characterization of the TCP‐TiO2 ‐MgF2 composites**

sintering process at 1200°C for 1 h. Beyond 4 wt% MgF<sup>2</sup>

composites are obtained by adding 4 wt% MgF<sup>2</sup>

in the TCP‐40 wt% TiO<sup>2</sup>

the densification of the TCP‐38 wt% TiO<sup>2</sup>

increases slowly with 10 wt% MgF<sup>2</sup>

1000°C; (b) 1100°C; (c) 1200°C and (d) 1300°C.

56 Sintering of Functional Materials

with different contents of MgF<sup>2</sup>

lus of the TCP‐40 wt% TiO<sup>2</sup>

TiO2

ages of MgF<sup>2</sup>

time (**Figure 13**).

was hindered (**Figure 11**). However, Vickers hardness of the TCP‐40 wt% TiO<sup>2</sup>

The evolution of the elastic modulus (E and G) of the TCP‐40 wt% TiO<sup>2</sup>

(**Figure 11**).

**Figure 6.** Vickers hardness versus temperature of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup>

, Vickers hardness of the composites

and these values reach 51 and 20 GPa,

composites after the sintering tempera‐

is shown in **Figure 12A** and **B**, respectively. The elastic modu‐

composites increase with the addition of MgF<sup>2</sup>

composites (**Figure 12A** and **B**).

**B**). Thus, the optimum values of Young's modulus and the Shear modulus of the TCP‐40 wt%

respectively. The performances of the samples were hindered with the increase of the percent‐

**Figure 13** illustrates a typical relation between the different lengths of the sintering time and

‐4 wt% MgF<sup>2</sup>

ture at 1200°C. The maximum value of densification (94%) was registered after the sintering process at 1200°C for 1 h (**Figure 13**). The densification's value decreases with the sintering

**Figure 14** depicts the rupture strength evolution of the samples after the sintering process at 1200°C for different length of the sintering time. The mechanical strength of the TCP‐TiO<sup>2</sup>

composites

at: (a)

composites sintered

(**Figure 12A** and

‐MgF<sup>2</sup>

After the sintering process, the samples were investigated by different characterization tech‐ niques such as: XRD, IR, 31P MAS‐NMR and SEM.

**Figure 8.** Shear modulus of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup> at: (a) 1000°C, (b) 1100°C, (c) 1200°C and (d) 1300°C.

**Figure 10.** Rupture strength of the TCP‐40 wt% TiO<sup>2</sup>

**Figure 11.** Vickers hardness of the TCP‐40 wt% TiO<sup>2</sup>

of MgF<sup>2</sup> .

of MgF<sup>2</sup> . composites sintered at 1200°C for 1 h with different percentages

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composites sintered at 1200°C for 1 h with different percentages

**Figure 9.** Relative density of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h with different percentages of MgF<sup>2</sup> .

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites http://dx.doi.org/10.5772/intechopen.68501 59

**Figure 10.** Rupture strength of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h with different percentages of MgF<sup>2</sup> .

**Figure 8.** Shear modulus of the β‐TCP sintered for 1 h with different percentages of TiO<sup>2</sup>

1200°C and (d) 1300°C.

58 Sintering of Functional Materials

**Figure 9.** Relative density of the TCP‐40 wt% TiO<sup>2</sup>

MgF<sup>2</sup> . at: (a) 1000°C, (b) 1100°C, (c)

composites sintered at 1200°C for 1 h with different percentages of

**Figure 11.** Vickers hardness of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h with different percentages of MgF<sup>2</sup> .

**Figure 13.** Relative density of the TCP‐38 wt% TiO<sup>2</sup>

**Figure 14.** Rupture strength of the TCP‐38 wt% TiO<sup>2</sup>

the sintering time.

the sintering time.

‐4 wt% MgF<sup>2</sup>

‐4 wt% MgF<sup>2</sup>

composites sintered at 1200°C for different length of

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composites sintered at 1200°C for different length of

**Figure 12.** Elastic modulus of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h with different percentages of MgF<sup>2</sup> : (A) Young's modulus and (B) the Shear modulus.

**Figure 17** shows the XRD patterns of the samples sintered at 1200°C for 1 h without and with different contents of MgF<sup>2</sup> (1, 2.5, 4, 4.5, 5, 7.5 and 10 wt%). The diffraction pattern of the TCP‐40 wt% TiO<sup>2</sup> composites sintered without MgF<sup>2</sup> displays the α‐TCP phase traces pres‐ ence (ICDD data file no. 43‐1484), the calcium titanate (CaTiO<sup>3</sup> ) (ICDD data file no. 89‐8033), the β‐TCP phase (ICDD data file no. 70‐2065) and the rutile phase of titania (ICDD data file no. 65‐1119) (**Figure 17a**). The introduction of 1 wt% MgF<sup>2</sup> to the TCP‐40 wt% TiO<sup>2</sup> led to the presence of traces of both calcium titanate (CaTiO<sup>3</sup> ) and fluorapatite (FAp) (ICDD data file no. 76‐0558) beside the β‐TCP phase and the rutile phase of titania in majority peaks (**Figure 17b**). Further increasing the percentage of MgF<sup>2</sup> to the TCP‐TiO<sup>2</sup> composites, we Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites http://dx.doi.org/10.5772/intechopen.68501 61

**Figure 13.** Relative density of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites sintered at 1200°C for different length of the sintering time.

**Figure 17** shows the XRD patterns of the samples sintered at 1200°C for 1 h without and with

the β‐TCP phase (ICDD data file no. 70‐2065) and the rutile phase of titania (ICDD data file

file no. 76‐0558) beside the β‐TCP phase and the rutile phase of titania in majority peaks

composites sintered without MgF<sup>2</sup>

ence (ICDD data file no. 43‐1484), the calcium titanate (CaTiO<sup>3</sup>

no. 65‐1119) (**Figure 17a**). The introduction of 1 wt% MgF<sup>2</sup>

the presence of traces of both calcium titanate (CaTiO<sup>3</sup>

(**Figure 17b**). Further increasing the percentage of MgF<sup>2</sup>

(1, 2.5, 4, 4.5, 5, 7.5 and 10 wt%). The diffraction pattern of the

displays the α‐TCP phase traces pres‐

to the TCP‐40 wt% TiO<sup>2</sup>

to the TCP‐TiO<sup>2</sup>

composites sintered at 1200°C for 1 h with different percentages of

) and fluorapatite (FAp) (ICDD data

) (ICDD data file no. 89‐8033),

led to

composites, we

different contents of MgF<sup>2</sup>

**Figure 12.** Elastic modulus of the TCP‐40 wt% TiO<sup>2</sup>

: (A) Young's modulus and (B) the Shear modulus.

TCP‐40 wt% TiO<sup>2</sup>

60 Sintering of Functional Materials

MgF<sup>2</sup>

**Figure 14.** Rupture strength of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites sintered at 1200°C for different length of the sintering time.

**Figure 15.** Vickers hardness of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites sintered at 1200°C for different length of the sintering time.

obtain more of fluorapatite (**Figure 17b**–**h**). Nevertheless, the β‐TCP phase is disappeared and the rutile phase quite occupied the majority of peaks (**Figure 17c**–**h**). The appearance of a new phase relative to Mg<sup>2</sup> (PO4 )F (ICDD data file no. 42‐0582) was detected after the addi‐ tion of 2.5, 4 and 4.5 wt% MgF<sup>2</sup> (**Figure 17c**–**e**). At higher contents of MgF<sup>2</sup> (after 4.5 wt%), the Mg<sup>2</sup> (PO4 )F phase is completely disappeared and the intensity of fluorapatite increased (**Figure 17f**–**h**).

**Figure 18** displays the FTIR spectra of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h without and with different percentages of MgF<sup>2</sup> (1, 2.5, 4, 4.5, 5, 7.5 and 10 wt%). Most bands characterize the phosphate group of calcium phosphate (**Figure 18**). Therefore, the bands at 954 and 965 cm−1 are assigned to the symmetric stretching of the PO4 3 ions while the bands at 1033 and 1093 cm−1 are relative to the asymmetric stretching of the PO4 3 ions (**Figure 18**).

The 31P MAS‐NMR spectra of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h without and with different amounts of MgF<sup>2</sup> (1, 2.5, 4, 4.5, 5, 7.5 and 10 wt%) are illustrated in **Figure 19**. This analysis indicates the presence of the tetrahedral environment, which is assigned to the resonance characteristic of the phosphate group (**Figure 19**). Moreover, the 31P MAS‐NMR spectrum of the TCP‐40 wt% TiO<sup>2</sup> composites sintered without MgF<sup>2</sup> shows many tetrahedral environments: one broad peak centered at 4 ppm, a shoulder at 0.47 ppm and an intense peak at −0.62 ppm (**Figure 19a**). The NMR spectrum of the TCP‐40 wt% TiO<sup>2</sup> ‐1 wt% MgF<sup>2</sup> composites presents an allure modification (**Figure 19b**). One intense peak located at 3 ppm is probably attributed to the phosphor of the phosphate groups of the FAp (**Figure 19c**–**h**). The second peak at 0.6 ppm is probably relative to the phosphor of the phosphate groups of the

Mg<sup>2</sup> (PO4

wt% MgF<sup>2</sup>

composites (**Figure 19f**–**h**).

sintering time: (A) Young's modulus and (B) the Shear modulus.

**Figure 16.** Elastic modulus of the TCP‐38 wt% TiO<sup>2</sup>

shows the micrographs of the TCP‐40 wt% TiO<sup>2</sup>

)F (**Figure 19b**–**e**). The intensity of the first peak increases with the amount of MgF<sup>2</sup>

‐4.5

composites sintered at 1200°C for 1 h

composites sintered at 1200°C for different length of the

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while the intensity of the second peak decreases and disappears with the TCP‐40 wt% TiO<sup>2</sup>

‐4 wt% MgF<sup>2</sup>

The microstructure of the sintered samples was observed using the SEM analysis. **Figure 20**

obtain more of fluorapatite (**Figure 17b**–**h**). Nevertheless, the β‐TCP phase is disappeared and the rutile phase quite occupied the majority of peaks (**Figure 17c**–**h**). The appearance of

‐4 wt% MgF<sup>2</sup>

characterize the phosphate group of calcium phosphate (**Figure 18**). Therefore, the bands at

in **Figure 19**. This analysis indicates the presence of the tetrahedral environment, which is assigned to the resonance characteristic of the phosphate group (**Figure 19**). Moreover, the 31P

tetrahedral environments: one broad peak centered at 4 ppm, a shoulder at 0.47 ppm and an

 composites presents an allure modification (**Figure 19b**). One intense peak located at 3 ppm is probably attributed to the phosphor of the phosphate groups of the FAp (**Figure 19c**–**h**). The second peak at 0.6 ppm is probably relative to the phosphor of the phosphate groups of the

intense peak at −0.62 ppm (**Figure 19a**). The NMR spectrum of the TCP‐40 wt% TiO<sup>2</sup>

)F (ICDD data file no. 42‐0582) was detected after the addi‐

(after 4.5 wt%),

composites sintered at 1200°C for 1

ions while the bands at

shows many

‐1 wt%

ions (**Figure 18**).

(1, 2.5, 4, 4.5, 5, 7.5 and 10 wt%). Most bands

composites sintered at 1200°C for different length of

composites sintered at 1200°C for 1 h

3

3

(1, 2.5, 4, 4.5, 5, 7.5 and 10 wt%) are illustrated

composites sintered without MgF<sup>2</sup>

(**Figure 17c**–**e**). At higher contents of MgF<sup>2</sup>

)F phase is completely disappeared and the intensity of fluorapatite increased

a new phase relative to Mg<sup>2</sup>

(PO4

(**Figure 17f**–**h**).

the sintering time.

62 Sintering of Functional Materials

the Mg<sup>2</sup>

MgF<sup>2</sup>

tion of 2.5, 4 and 4.5 wt% MgF<sup>2</sup>

**Figure 15.** Vickers hardness of the TCP‐38 wt% TiO<sup>2</sup>

(PO4

**Figure 18** displays the FTIR spectra of the TCP‐40 wt% TiO<sup>2</sup>

954 and 965 cm−1 are assigned to the symmetric stretching of the PO4

1033 and 1093 cm−1 are relative to the asymmetric stretching of the PO4

h without and with different percentages of MgF<sup>2</sup>

The 31P MAS‐NMR spectra of the TCP‐40 wt% TiO<sup>2</sup>

without and with different amounts of MgF<sup>2</sup>

MAS‐NMR spectrum of the TCP‐40 wt% TiO<sup>2</sup>

**Figure 16.** Elastic modulus of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites sintered at 1200°C for different length of the sintering time: (A) Young's modulus and (B) the Shear modulus.

Mg<sup>2</sup> (PO4 )F (**Figure 19b**–**e**). The intensity of the first peak increases with the amount of MgF<sup>2</sup> while the intensity of the second peak decreases and disappears with the TCP‐40 wt% TiO<sup>2</sup> ‐4.5 wt% MgF<sup>2</sup> composites (**Figure 19f**–**h**).

The microstructure of the sintered samples was observed using the SEM analysis. **Figure 20** shows the micrographs of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h

**Figure 17.** XRD patterns of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h without and with different percentages of MgF<sup>2</sup> : (a) 0 wt%, (b) 1 wt%, (c) 2.5 wt%, (d) 4 wt%, (e) 4.5 wt%, (f) 5 wt%, (g) 7.5 wt% and (h) 10 wt%. (o: α‐TCP, ♦: β‐TCP, +: CaTiO<sup>3</sup> , \*: rutile, ♣: fluorapatite, ●: Mg<sup>2</sup> (PO4 )F).

without and with different percentages of MgF<sup>2</sup>

**Figure 19.** 31P MAS‐NMR spectra of the TCP‐40 wt% TiO<sup>2</sup>

composites sintered with 2.5 wt% MgF<sup>2</sup>

composites sintered without MgF<sup>2</sup>

the formation of a new lamella form (**Figure 20b**). At higher amount of MgF<sup>2</sup>

sification (94%) was registered after the addition of 4 wt% MgF<sup>2</sup>

‐7.5 wt% MgF<sup>2</sup>

form (**Figure 20e**–**f**). Moreover, the micrographs of the TCP‐35 wt% TiO<sup>2</sup>

Moreover, the microstructure of the TCP‐38 wt% TiO<sup>2</sup>

Both the addition of 7.5 and 10 wt% of MgF<sup>2</sup>

grains of the TCP and the titania (**Figure 20a**). The microstructure of the TCP‐40 wt% TiO<sup>2</sup> composites shows a liquid phase, a continuous phase relative to the β‐TCP phases and a small‐sized grain relative to the titania (**Figure 20a**). The morphology of the TCP=40 wt%

This fact is explained by the intergranular porosity presence beside the grain growth and

this new phase still presented in the microstructure (**Figure 20c**–**d**). An important den‐

composites (**Figure 20c**–**d**). This result can be explicated by the reduction in the porosity.

continuous phases with dense contacts between the grains of the samples (**Figure 20c**–**d**).

and the properties of the samples were hindered by the formation in the bubbles form and the grain growth in the microstructure (**Figure 20e**–**h**). Thus, the microstructure of

depict the grains in hexagonal form formation beside the bubbles form (**Figure 20g**–**h**).

of the TCP‐40 wt% TiO<sup>2</sup>

different percentages of MgF<sup>2</sup>

the TCP‐36.25 wt% TiO<sup>2</sup>

TiO2

(h) 10 wt%.

(2.5, 4, 7.5 and 10 wt%). The micrograph

composites sintered at 1200°C for 1 h without and with

: (a) 0 wt%, (b) 1 wt%, (c) 2.5 wt%, (d) 4 wt%, (e) 4.5 wt%, (f) 5 wt%, (g) 7.5 wt% and

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

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

65

was completely transformed (**Figure 20b**).

‐4 wt% MgF<sup>2</sup>

composites sintered at 1200°C presents the bubbles

led to the disappearance of the lamella form

depicts the coalescence of the

to the TCP‐40 wt% TiO<sup>2</sup>

composites shows

‐10 wt% MgF<sup>2</sup>

(4 wt%),

**Figure 18.** Infrared spectra of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h without and with different percentages of MgF<sup>2</sup> : (a) 0 wt%, (b) 1 wt%, (c) 2.5 wt%, (d) 4 wt%, (e) 4.5 wt%, (f) 5 wt%, (g) 7.5 wt% and (h) 10 wt%.

**Figure 19.** 31P MAS‐NMR spectra of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h without and with different percentages of MgF<sup>2</sup> : (a) 0 wt%, (b) 1 wt%, (c) 2.5 wt%, (d) 4 wt%, (e) 4.5 wt%, (f) 5 wt%, (g) 7.5 wt% and (h) 10 wt%.

**Figure 17.** XRD patterns of the TCP‐40 wt% TiO<sup>2</sup>

**Figure 18.** Infrared spectra of the TCP‐40 wt% TiO<sup>2</sup>

percentages of MgF<sup>2</sup>

, \*: rutile, ♣: fluorapatite, ●: Mg<sup>2</sup>

percentages of MgF<sup>2</sup>

(o: α‐TCP, ♦: β‐TCP, +: CaTiO<sup>3</sup>

64 Sintering of Functional Materials

composites sintered at 1200°C for 1 h without and with different

composites sintered at 1200°C for 1 h without and with different

: (a) 0 wt%, (b) 1 wt%, (c) 2.5 wt%, (d) 4 wt%, (e) 4.5 wt%, (f) 5 wt%, (g) 7.5 wt% and (h) 10 wt%.

: (a) 0 wt%, (b) 1 wt%, (c) 2.5 wt%, (d) 4 wt%, (e) 4.5 wt%, (f) 5 wt%, (g) 7.5 wt% and (h) 10 wt%.

(PO4 )F).

> without and with different percentages of MgF<sup>2</sup> (2.5, 4, 7.5 and 10 wt%). The micrograph of the TCP‐40 wt% TiO<sup>2</sup> composites sintered without MgF<sup>2</sup> depicts the coalescence of the grains of the TCP and the titania (**Figure 20a**). The microstructure of the TCP‐40 wt% TiO<sup>2</sup> composites shows a liquid phase, a continuous phase relative to the β‐TCP phases and a small‐sized grain relative to the titania (**Figure 20a**). The morphology of the TCP=40 wt% TiO2 composites sintered with 2.5 wt% MgF<sup>2</sup> was completely transformed (**Figure 20b**). This fact is explained by the intergranular porosity presence beside the grain growth and the formation of a new lamella form (**Figure 20b**). At higher amount of MgF<sup>2</sup> (4 wt%), this new phase still presented in the microstructure (**Figure 20c**–**d**). An important den‐ sification (94%) was registered after the addition of 4 wt% MgF<sup>2</sup> to the TCP‐40 wt% TiO<sup>2</sup> composites (**Figure 20c**–**d**). This result can be explicated by the reduction in the porosity. Moreover, the microstructure of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites shows continuous phases with dense contacts between the grains of the samples (**Figure 20c**–**d**). Both the addition of 7.5 and 10 wt% of MgF<sup>2</sup> led to the disappearance of the lamella form and the properties of the samples were hindered by the formation in the bubbles form and the grain growth in the microstructure (**Figure 20e**–**h**). Thus, the microstructure of the TCP‐36.25 wt% TiO<sup>2</sup> ‐7.5 wt% MgF<sup>2</sup> composites sintered at 1200°C presents the bubbles form (**Figure 20e**–**f**). Moreover, the micrographs of the TCP‐35 wt% TiO<sup>2</sup> ‐10 wt% MgF<sup>2</sup> depict the grains in hexagonal form formation beside the bubbles form (**Figure 20g**–**h**).

of the TCP while some researchers observed a minimal or very slow resorption [6, 59, 60]. Consequently, the addition of the inert oxide can enhance the mechanical performances of the TCP in order to combine the high mechanical and tricological properties of titania with the resorbability of TCP. The study of the tricalcium phosphate/titania system interested on improving the densification and the mechanical properties of the TCP by the addition

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

. The properties of the TCP‐TiO<sup>2</sup>

addition of a reinforcement agent like magnesium fluoride (MgF<sup>2</sup>

on the mechanical properties of the TCP‐40 wt% TiO<sup>2</sup>

ison between the previous work derived from the TCP‐40 wt% TiO<sup>2</sup>

‐4 wt% MgF<sup>2</sup>

(the TCP‐38 wt% TiO<sup>2</sup>

. The optimum values (ρ = 94%, σr

accordance with the amount of titania and the sintered temperature. Thus, the optimum val‐ ues of these composites were acquired after the sintering at 1200°C with 40 wt% of titania. The objective of this study aimed to ameliorate the performances of these composites by the

a dopant in both apatite and β‐TCP structures, has been the subject of specific interest due to its important role in biological processes after implantation [61, 62]. Magnesium plays an essential role in the formation of dental caries and bone deposition. Magnesium ions promote bone mineralization and control the growth of the calcium phosphate crystals under in vivo conditions [62, 63]. Magnesium indirectly influences bone resorption and directly stimulates osteoblast proliferation with an effect comparable to that of insulin (a known growth factor for osteoblast) [62–64]. Thus, magnesium containing calcium phosphates is more appropri‐ ate for biomaterials and their applications [62]. Magnesium ions play important roles in the formation and transformation of biologically related calcium phosphates [65, 66]. MgF<sup>2</sup>

is known as the most promising dental caries‐promoting element due to a high initial cal‐ cification process in the bone formation [37, 38]. Also, fluoride is known to be important in suppressing dental caries [33, 34]. It stimulates the proliferation and the differentiation of

and dental formation in the human body [35, 36]. The effect of adding different percentages

for different lengths of the sintering time was investigated. The mechanical performances were studied by Vickers indentation, Brazilian test and ultrasound techniques to determine

of the bone tissues and precisely the enamel [6, 8, 67–70]. Thus, the enamel is mainly made of the mineral in the calcium phosphate [69, 71]. It is the hardest tissue in the human body because it contains almost no water [69]. For this reason, it is good to get a similar behavior as enamel. In fact, Chun et al. studied the mechanical properties of both enamel and dentin [69]. Therefore, if we compare the value of the strain and stress, enamel tend to fractures earlier than dentin, so it is more brittle than dentine. Nevertheless, if we compare from Vickers hard‐ ness value, the enamel is harder than dentine. In fact, dentin was characterized by a higher force resistance while enamel was characterized by a higher wear resistance [69]. The compar‐

conclusion described between enamel and dentin. However, as far as hardness is considered,

‐4 wt% MgF<sup>2</sup>

composites are considered harder than the TCP‐40 wt%

) has been investigated as an essential element for bone

composites were increased in

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

). Recently, magnesium, as

composites sintered at 1200°C

composites increase as in terms of

= 27 MPa, E = 51 GPa, G = 20 GPa

composites and the pres‐

composites) leads to the same

) and the elastic moduli (E and G), respec‐

. These values are relatively similar to those

composites are attained after the sintering process

ion

67

of different amount of TiO<sup>2</sup>

bone cells [35]. The fluorine ion (F−

Vickers hardness (H), the mechanical strength (σr

at 1200°C for 1 h with the addition of 4 wt% MgF<sup>2</sup>

and H = 360 Hv) of the TCP‐40 wt% TiO<sup>2</sup>

tively. The mechanical properties of the TCP‐40 wt% TiO<sup>2</sup>

of MgF<sup>2</sup>

the amounts of MgF<sup>2</sup>

ent work, with adding MgF<sup>2</sup>

the TCP‐38 wt% TiO<sup>2</sup>

**Figure 20.** SEM micrographs of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for 1 h without and with different percentages of MgF<sup>2</sup> : (a) 0 wt% (×10,000), (b) 2.5 wt% (×5000), (c) 4 wt% (×20,000), (d) 4 wt% (×5000), (e) 7.5 wt% (×20,000), (f) 7.5 wt% (×5000), (g) 10 wt% (×20,000) and (h) 10 wt% (×5000).

### **4. Discussion**

In the light of its astonishing properties in terms of biocompatibility, in vivo bioactivity, osteoconductivity and bio‐resorbability, tricalcium phosphate has been emerged as one of the most imperative biomaterials [6, 58, 59]. Many authors reported rather fast degradation of the TCP while some researchers observed a minimal or very slow resorption [6, 59, 60]. Consequently, the addition of the inert oxide can enhance the mechanical performances of the TCP in order to combine the high mechanical and tricological properties of titania with the resorbability of TCP. The study of the tricalcium phosphate/titania system interested on improving the densification and the mechanical properties of the TCP by the addition of different amount of TiO<sup>2</sup> . The properties of the TCP‐TiO<sup>2</sup> composites were increased in accordance with the amount of titania and the sintered temperature. Thus, the optimum val‐ ues of these composites were acquired after the sintering at 1200°C with 40 wt% of titania. The objective of this study aimed to ameliorate the performances of these composites by the addition of a reinforcement agent like magnesium fluoride (MgF<sup>2</sup> ). Recently, magnesium, as a dopant in both apatite and β‐TCP structures, has been the subject of specific interest due to its important role in biological processes after implantation [61, 62]. Magnesium plays an essential role in the formation of dental caries and bone deposition. Magnesium ions promote bone mineralization and control the growth of the calcium phosphate crystals under in vivo conditions [62, 63]. Magnesium indirectly influences bone resorption and directly stimulates osteoblast proliferation with an effect comparable to that of insulin (a known growth factor for osteoblast) [62–64]. Thus, magnesium containing calcium phosphates is more appropri‐ ate for biomaterials and their applications [62]. Magnesium ions play important roles in the formation and transformation of biologically related calcium phosphates [65, 66]. MgF<sup>2</sup> ion is known as the most promising dental caries‐promoting element due to a high initial cal‐ cification process in the bone formation [37, 38]. Also, fluoride is known to be important in suppressing dental caries [33, 34]. It stimulates the proliferation and the differentiation of bone cells [35]. The fluorine ion (F− ) has been investigated as an essential element for bone and dental formation in the human body [35, 36]. The effect of adding different percentages of MgF<sup>2</sup> on the mechanical properties of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C for different lengths of the sintering time was investigated. The mechanical performances were studied by Vickers indentation, Brazilian test and ultrasound techniques to determine Vickers hardness (H), the mechanical strength (σr ) and the elastic moduli (E and G), respec‐ tively. The mechanical properties of the TCP‐40 wt% TiO<sup>2</sup> composites increase as in terms of the amounts of MgF<sup>2</sup> . The optimum values (ρ = 94%, σr = 27 MPa, E = 51 GPa, G = 20 GPa and H = 360 Hv) of the TCP‐40 wt% TiO<sup>2</sup> composites are attained after the sintering process at 1200°C for 1 h with the addition of 4 wt% MgF<sup>2</sup> . These values are relatively similar to those of the bone tissues and precisely the enamel [6, 8, 67–70]. Thus, the enamel is mainly made of the mineral in the calcium phosphate [69, 71]. It is the hardest tissue in the human body because it contains almost no water [69]. For this reason, it is good to get a similar behavior as enamel. In fact, Chun et al. studied the mechanical properties of both enamel and dentin [69].

Therefore, if we compare the value of the strain and stress, enamel tend to fractures earlier than dentin, so it is more brittle than dentine. Nevertheless, if we compare from Vickers hard‐ ness value, the enamel is harder than dentine. In fact, dentin was characterized by a higher force resistance while enamel was characterized by a higher wear resistance [69]. The compar‐ ison between the previous work derived from the TCP‐40 wt% TiO<sup>2</sup> composites and the pres‐ ent work, with adding MgF<sup>2</sup> (the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites) leads to the same conclusion described between enamel and dentin. However, as far as hardness is considered, the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites are considered harder than the TCP‐40 wt%

**4. Discussion**

percentages of MgF<sup>2</sup>

66 Sintering of Functional Materials

**Figure 20.** SEM micrographs of the TCP‐40 wt% TiO<sup>2</sup>

(f) 7.5 wt% (×5000), (g) 10 wt% (×20,000) and (h) 10 wt% (×5000).

In the light of its astonishing properties in terms of biocompatibility, in vivo bioactivity, osteoconductivity and bio‐resorbability, tricalcium phosphate has been emerged as one of the most imperative biomaterials [6, 58, 59]. Many authors reported rather fast degradation

: (a) 0 wt% (×10,000), (b) 2.5 wt% (×5000), (c) 4 wt% (×20,000), (d) 4 wt% (×5000), (e) 7.5 wt% (×20,000),

composites sintered at 1200°C for 1 h without and with different

TiO2 composites. Such similarity between the present work and the enamel behaviors allows us to study the performances of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites in biomedical applications and specifically as enamel.

The values of the elastic modulus (E and G) of the TCP‐38 wt% TiO<sup>2</sup>

values with the elastic modulus of the TCP‐40 wt% TiO<sup>2</sup>

The evolution of the mechanical properties (E, G, H and σr

MgF<sup>2</sup>

MgF<sup>2</sup>

wt% TiO<sup>2</sup>

(Ca5 (PO4 )3

[Mg<sup>2</sup> (PO4

wt% MgF<sup>2</sup>

wt% TiO<sup>2</sup>

wt% MgF<sup>2</sup>

the Mg<sup>2</sup>

phase (Mg<sup>2</sup>

to the TCP‐TiO<sup>2</sup>

percentages of MgF<sup>2</sup>

(PO4

amounts of MgF<sup>2</sup>

‐4 wt% MgF<sup>2</sup>

FAp with more quantities in the TCP‐TiO<sup>2</sup>

proves the presence of traces of CaTiO<sup>3</sup>

The XRD spectra of the TCP‐TiO2‐MgF<sup>2</sup>

F) and Mg<sup>2</sup>

(PO4

F and the fluorapatite (Ca5

2Ca3

)F and Ca5

of the TCP‐38 wt% TiO<sup>2</sup>

‐4 wt% MgF<sup>2</sup>

(PO4 )3 F].

The diffractograms of the TCP‐40 wt% TiO<sup>2</sup>

(PO4

compounds. The interaction between TCP and MgF<sup>2</sup>

(PO4 )3

‐4 wt% MgF<sup>2</sup>

, the mechanical properties of the TCP‐TiO<sup>2</sup>

titania (**Figure 17**). The spectrum of the TCP‐40 wt% TiO<sup>2</sup>

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

‐MgF<sup>2</sup>

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 MgF<sup>2</sup>

(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% TiO<sup>2</sup>

composites structure enhances their performances with 4 wt%MgF<sup>2</sup>

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

‐MgF<sup>2</sup>

)2(s) + 2MgF2(s) + TiO2(s) + H<sup>2</sup>

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% TiO<sup>2</sup>

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 Mg<sup>2</sup>

)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% MgF<sup>2</sup>

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

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

composites sintered without MgF<sup>2</sup>

) of the TCP‐38 wt% TiO<sup>2</sup>

composites. Consequently, the MgF<sup>2</sup>

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 Mg<sup>2</sup>

O ↔ Ca<sup>5</sup>

composites. These results confirm those obtained with

(PO4

composites sintered without MgF<sup>2</sup>

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% MgF<sup>2</sup>

.

and H) of the TCP‐38

The densification discrepancy of the samples prepared by a mixture of the TCP‐40 wt% TiO<sup>2</sup> composites with adding 4 wt% MgF<sup>2</sup> (<sup>ρ</sup> <sup>=</sup> 94%) and without MgF<sup>2</sup> (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‐TiO<sup>2</sup> ‐MgF<sup>2</sup> 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% TiO<sup>2</sup> composites was obtained after the addition of 4 wt% MgF<sup>2</sup> (H = 360 Hv). If we compare this result with that of the TCP‐40 wt% TiO<sup>2</sup> composites without adding the magnesium fluoride, we notice an increase from 270 to 360 Hv. So, adding 4 wt% MgF<sup>2</sup> enhances Vickers hardness of the TCP‐40 wt% TiO<sup>2</sup> 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% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> 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% TiO<sup>2</sup> composites sintered without MgF<sup>2</sup> , we notice that the elastic corresponding modulus (E and G) of the composites sintered with 4 wt% MgF<sup>2</sup> 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

(<sup>ρ</sup> <sup>=</sup> 94%) and without MgF<sup>2</sup>

The densification discrepancy of the samples prepared by a mixture of the TCP‐40 wt% TiO<sup>2</sup>

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% MgF<sup>2</sup>

‐MgF<sup>2</sup>

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

enhances Vickers hardness of the TCP‐40 wt% TiO<sup>2</sup>

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% TiO<sup>2</sup>

ing helps to fill the pore in the microstructure of the TCP‐TiO<sup>2</sup>

optimum value of Vickers hardness of the TCP‐40 wt% TiO<sup>2</sup>

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

applications and specifically as enamel.

composites with adding 4 wt% MgF<sup>2</sup>

the addition of 4 wt% MgF<sup>2</sup>

to 360 Hv. So, adding 4 wt% MgF<sup>2</sup>

wt% TiO<sup>2</sup>

Hv) (**Table 4**).

β‐TCP‐40 wt% TiO<sup>2</sup> composites

‐4 wt% MgF<sup>2</sup> 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% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> 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% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites sintered for 300 min or the TCP‐35 wt% TiO<sup>2</sup> ‐10 wt% MgF<sup>2</sup> composites sintered for 60 min are improved which is probably due to the formation of the FAp with more quantities in the TCP‐TiO<sup>2</sup> ‐MgF<sup>2</sup> composites. Consequently, the MgF<sup>2</sup> added to the TCP‐TiO<sup>2</sup> composites structure enhances their performances with 4 wt%MgF<sup>2</sup> .

The diffractograms of the TCP‐40 wt% TiO<sup>2</sup> composites sintered at 1200°C with different percentages of MgF<sup>2</sup> indicate that the majority of peaks are relative to the rutile phase of titania (**Figure 17**). The spectrum of the TCP‐40 wt% TiO<sup>2</sup> composites sintered without MgF<sup>2</sup> proves the presence of traces of CaTiO<sup>3</sup> and α‐TCP besides the rutile phase and the β‐TCP. The XRD spectra of the TCP‐TiO2‐MgF<sup>2</sup> composites depicts the formation of both fluorapatite (Ca5 (PO4 )3 F) and Mg<sup>2</sup> (PO4 )F which is in agreement with the 31P MAS‐NMR analysis. Thus, the Mg<sup>2</sup> (PO4 )F was obtained only for 2.5, 4 and 4.5 wt% of MgF<sup>2</sup> and disappears for higher amounts of MgF<sup>2</sup> (**Figure 17**).This disappearance is attributed to the incorporation of this phase (Mg<sup>2</sup> (PO4 )F) into the calcium phosphate structure by a solid state reaction through the compounds. The interaction between TCP and MgF<sup>2</sup> promotes the formation of the Mg<sup>2</sup> (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 [Mg<sup>2</sup> (PO4 )F and Ca5 (PO4 )3 F].

SEM analyses prove that excellent performances are attained with the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> 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% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites. These results confirm those obtained with the XRD analysis. Thus, the new lamella form detected in the microstructure of the TCP‐38 wt% TiO<sup>2</sup> ‐4 wt% MgF<sup>2</sup> composites is probably accorded to the Mg<sup>2</sup> (PO4 )F derived from the solid reaction between MgF2 and TCP according to the previous reaction (Eq. (7)). Above 4 wt% MgF<sup>2</sup> , the mechanical properties of the TCP‐TiO<sup>2</sup> ‐MgF<sup>2</sup> 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% TiO<sup>2</sup> ‐10 wt% MgF<sup>2</sup> 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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## **5. Conclusion**

The effect of the MgF<sup>2</sup> addition on the performances of the TCP‐40 wt% TiO<sup>2</sup> composites was studied during the sintering process at 1200°C for different length of the sintering time. The addition of 4 wt% MgF<sup>2</sup> to the TCP‐40 wt% TiO<sup>2</sup> 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% TiO<sup>2</sup> composites was attained with 4 wt% MgF<sup>2</sup> (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 Mg<sup>2</sup> (PO4 )F and the hexag‐ onal grains relative to FAp. Above 4 wt% MgF<sup>2</sup> , 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: benayedfoued@yahoo.fr

Laboratory of Industrial Chemistry, National Engineering School, Sfax University, Sfax, Tunisia

## **References**


[3] Bouslama N, Ben Ayed F, Bouaziz J. Sintering and mechanical properties of tricalcium phosphate‐fluorapatite composite. Ceramics International. 2009;**35**:1909‐1917. DOI: 10.1016/j.ceramint.2008.10.030

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% TiO<sup>2</sup>

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% TiO<sup>2</sup>

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‐

Laboratory of Industrial Chemistry, National Engineering School, Sfax University, Sfax, Tunisia

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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% MgF<sup>2</sup>

to the TCP‐40 wt% TiO<sup>2</sup>

structure indicates the presence of new lamella form relative to the Mg<sup>2</sup>

graph of the TCP‐35 wt% TiO<sup>2</sup>

70 Sintering of Functional Materials

**5. Conclusion**

The effect of the MgF<sup>2</sup>

addition of 4 wt% MgF<sup>2</sup>

tion‐precipitation phenomena.

Ibticem Ayadi and Foued Ben Ayed\*

**Author details**

**References**

composites was attained with 4 wt% MgF<sup>2</sup>

onal grains relative to FAp. Above 4 wt% MgF<sup>2</sup>

\*Address all correspondence to: benayedfoued@yahoo.fr


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[42] Evis Z, Usta M, Kutbay I. Improvement in sinterability and phase stability of hydroxy‐ apatite and partially stabilized zirconia composites. Journal of the European Ceramic Society. 2009;**29**:621‐628. DOI: http://dx.doi.org/10.1016/j.jeurceramsoc.2008.07.020 DOI:

[43] Pardun K, Treccani L, Volkmann E, Streckbein P, Heiss C, Gerlach JW, Maendl S, Rezwan K. Magnesium containing mixed coatings on zirconia for dental implants: Mechanical characterization and in vitro behavior. Journal of Biomaterials Applications.

[44] Chen X, Chen X, Brauer DS, Wilson RM, Hill RG, Karpukhina N. Novel alkali free bio‐ active fluorapatite glass ceramics. Journal of Non‐Crystalline Solids. 2014;**402**:172‐177.

[45] Kanazawa T, Umegaki T, Yamashita K, Monma H, Hiramatsu T. Effects of additives on sintering and some properties of calcium phosphates with various Ca/P Ratios. Journal

[46] Hench LL. Bioceramics: From concept to clinic. Journal of the American Ceramic Society.

[47] Azami M, Jalilifiroozinezhad S, Mozafari M, Rabiee M. Synthesis and solubility of cal‐ cium fluoride/hydroxy‐fluorapatite nanocrystals for dental applications. Ceramics International. 2011;**37**:2007‐2014. DOI: http://dx.doi.org/10.1016/j.ceramint.2011.02.025

[48] Carsten J, Marco B. 19F NMR spectroscopy of glass ceramics containing fluorapatites.

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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

[67] Ravaglioli A, Krajewski A. Handbook of Bioceramics: Materials, Properties and Appli cations. London, England: Chapman & Hall; 1992. 422 p. DOI: 10.1007/97 8‐94‐011‐2 336‐5

[68] Shi D, editor. Handbook of Biomaterials and Tissue Engineering. 1st ed. New York: Springer‐Verlag Berlin Heidelberg; 2004. 200 p. DOI: 10.1007/978‐3‐662‐06104‐6

[69] Chun KJ, Choi HH, Lee JY. Comparison of mechanical property and role between enamel and dentin in the human teeth. Journal of Dental Biomechanics. 2014;**5**:1‐6. DOI:

[70] Biswas N, Dey A, Kundu S, Chakraborty H, Mukhopadhyay AK. Mechanical proper‐ ties of enamel nanocomposite. ISRN Biomaterials Hindawi. 2013;**2013**:1‐15. DOI: http://

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

[72] Hosokawa M, Nogi K, Naito M, Yokoyama T, editors. Handbook of Nanoparticle Technology. 1st ed. Linacre House, Amsterdam, The Netherlands: Elsevier BV; 2007. 644

[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

Pharmaceutical Sciences. 2008;**97**(1):421‐432. DOI: 10.1002/jps.21131

10.1177/1758736014520809 DOI:10.1177%2F1758736014520809#pmc\_ext

Materials Science. 1981;**16**(9):2551‐2556. DOI: 10.1007/BF01113595

1870604

76 Sintering of Functional Materials

dx.doi.org/10.5402/2013/253761

p. ISBN: 9780080558028

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, Al<sup>2</sup>

iii. Rare-earths-based magnetic materials, for example, Nd<sup>9</sup>

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

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

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

)<sup>2</sup>

⋅6H

O2

)<sup>3</sup>

⋅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

)<sup>3</sup>

⋅9H

O2

)<sup>3</sup>

⋅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)**. The striking difference arises from the

MSS 0.19 0.21 0.23 0.24 0.43 1.07 1.23 2.65 4.98 SSS 0.13 0.15 0.19 0.23 0.30 0.39 1.05 2.08 5.35

MSS 3.48 3.62 3.7 3.98 4.02 4.2 4.64 4.73 4.56 SSS 4.23 4.49 4.62 4.73 4.78 4.81 4.91 4.93 4.88

MSS 23.9 29.7 97.8 503.0 522.8 865.3 908.3 949.7 1076.0 SSS 23.4 29.5 68.3 424.2 523.0 572.0 605.3 774.7 930.7

MSS 3.4 7.7 12.5 5.0 4.1 1.3 1.2 0.5 0.4 SSS 3.0 9.7 11.5 6.7 3.8 3.5 1.6 0.9 0.3

**Sintering temperature (°C) 600 700 800 900 1000 1100 1200 1300 1400**

**Table 2.** 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/cm<sup>3</sup>

Saturation induction, B<sup>s</sup> (Gauss)

Coercivity, H<sup>c</sup> (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, B<sup>s</sup> , of nickel zinc ferrite could range from 1000 to 3000 G [20, 41]. The experimental values of B<sup>s</sup> 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 B<sup>s</sup> 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 (B<sup>s</sup>

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

(M<sup>s</sup>

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 B<sup>s</sup>

**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 B<sup>s</sup> (M<sup>s</sup> ) 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 (B<sup>s</sup> , 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, B<sup>s</sup>

mental values of B<sup>s</sup>

90 Sintering of Functional Materials

23.4 to 930.7 G for SSS. The various ranges of B<sup>s</sup>

**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 H<sup>c</sup> 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 H<sup>c</sup> 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. H<sup>c</sup> 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 H<sup>c</sup> 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 H<sup>c</sup> [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 H<sup>c</sup> . 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 H<sup>c</sup> 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*<sup>i</sup> 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 H<sup>c</sup>

92 Sintering of Functional Materials

drop of H<sup>c</sup>

 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].

<sup>R</sup> = (1/µ<sup>i</sup>

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 H<sup>c</sup>

as compared to MSS where it drops at 0.23 µm. H<sup>c</sup>

few microns. An inversely proportional trend of H<sup>c</sup>

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

frequency fR and the initial permeability *μ*<sup>i</sup>

f

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

ened grains for Ni0.3Zn0.7Fe2

movement, thus decreasing the H<sup>c</sup>

a maximum H<sup>c</sup>

**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, Fe<sup>2</sup> 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

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

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, Fe<sup>2</sup>

3Y<sup>2</sup>

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 → 2Y<sup>3</sup>

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

**Grain size, (±0.01 μm)**

**Saturation induction, Bs (Gauss)**

**Saturation magnetization, Ms (emu/cm3**

**)**

**Coercivity, Hc (Oe)**

**Single-sample sintering Multi-sample sintering**

 0.16 2.1 1.7 2.6 0.20 16.9 2.2 0.1 0.17 2.3 2.6 3.1 0.21 24.9 2.7 0.7 0.18 16.4 3.9 6.7 0.25 35.8 4.7 1.3 0.28 20.6 4.8 10.9 0.26 49.6 5.2 3.9 0.33 120.7 5.4 15.5 0.28 128.6 5.5 15.8 0.60 173.2 5.7 18.5 0.58 185.7 6.2 19.3 1.14 223.7 10.9 12.4 0.80 244.5 12.8 15.2 1.68 378.9 21.2 7.4 1.25 463.1 23.3 8.8 2.71 570.4 26.3 4.3 3.09 714.6 29.1 2.9

**Saturation magnetization,** 

> **(emu/cm3 )**

**Ms**

Fe5

O3

(Alfa Aesar, 99.945%) and

O12 (2)

**garnet**

96 Sintering of Functional Materials

Y2 O3

**Sintering temperature, (T, °C)**

**Figure 10.** 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 Ismail<sup>1</sup>

**References**

\*, Idza Riati Ibrahim<sup>1</sup>

\*Address all correspondence to: ismayadi@upm.edu.my

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 Y<sup>3</sup>

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 Nazlan<sup>2</sup>

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-

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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 Ismail<sup>1</sup> \*, Idza Riati Ibrahim<sup>1</sup> and Rodziah Nazlan<sup>2</sup>

\*Address all correspondence to: ismayadi@upm.edu.my

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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**Chapter 5**

**Provisional chapter**

**Sintering of Whiteware Body Depending on Different**

The sintering of whiteware (porcelain) body can be affected by using fluxing agents or binders. The chapter describes the sintering process of porcelain body in case of different fluxing agent (different feldspar rocks, bone ash, zeolite) and binder (kaolin vs. calcium aluminate cement) utilization in the porcelain raw material mixture. Sintering process is presented according to thermodilatometrical curves and sintering tempera-

**Keywords:** whitewares, feldspar rocks, zeolite, bone ash, kaolin, calcium aluminate cement, sintering temperature, water absorption, mineralogical composition

Whiteware is a traditional ceramic material used to make pottery and porcelain. Traditional raw material mixture for whiteware (porcelain) production covers kaolin or/and kaolin clay, quartz and feldspar rock at a composition about 50:25:25 wt.%. Typical properties of porcelain body are low porosity (below 0.3%), high mechanical strength (bending strength over 40 MPa, Young Modulus over 60 GPa), firing temperature about 1300°C and high whiteness

Feldspar rocks are used in the fine ceramic industry as a fluxing agent to form a glassy phase for accelerating of sintering process. Feldspar rocks are a mixture of pure feldspars, quartz and mica especially from the mineralogical point of view. Pure feldspars are divided into potassium feldspars (orthoclase, microcline), sodium feldspars (albite) and calcium feldspars (anorthite). Solid solutions between K‐feldspar and albite are called alkali feldspars, and solid solutions between albite and anorthite are plagioclase feldspars.

**Sintering of Whiteware Body Depending on Different** 

DOI: 10.5772/68082

© 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.

**Fluxing Agents and Binders**

**Fluxing Agents and Binders**

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

Radomir Sokolar

Radomir Sokolar

http://dx.doi.org/10.5772/68082

**Abstract**

tures especially.

**1. Introduction**

and translucency [1–3].

**Provisional chapter**

## **Sintering of Whiteware Body Depending on Different Fluxing Agents and Binders Fluxing Agents and Binders**

**Sintering of Whiteware Body Depending on Different** 

DOI: 10.5772/68082

Radomir Sokolar Additional information is available at the end of the chapter

Radomir Sokolar

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/68082

#### **Abstract**

The sintering of whiteware (porcelain) body can be affected by using fluxing agents or binders. The chapter describes the sintering process of porcelain body in case of different fluxing agent (different feldspar rocks, bone ash, zeolite) and binder (kaolin vs. calcium aluminate cement) utilization in the porcelain raw material mixture. Sintering process is presented according to thermodilatometrical curves and sintering temperatures especially.

**Keywords:** whitewares, feldspar rocks, zeolite, bone ash, kaolin, calcium aluminate cement, sintering temperature, water absorption, mineralogical composition

## **1. Introduction**

Whiteware is a traditional ceramic material used to make pottery and porcelain. Traditional raw material mixture for whiteware (porcelain) production covers kaolin or/and kaolin clay, quartz and feldspar rock at a composition about 50:25:25 wt.%. Typical properties of porcelain body are low porosity (below 0.3%), high mechanical strength (bending strength over 40 MPa, Young Modulus over 60 GPa), firing temperature about 1300°C and high whiteness and translucency [1–3].

Feldspar rocks are used in the fine ceramic industry as a fluxing agent to form a glassy phase for accelerating of sintering process. Feldspar rocks are a mixture of pure feldspars, quartz and mica especially from the mineralogical point of view. Pure feldspars are divided into potassium feldspars (orthoclase, microcline), sodium feldspars (albite) and calcium feldspars (anorthite). Solid solutions between K‐feldspar and albite are called alkali feldspars, and solid solutions between albite and anorthite are plagioclase feldspars.

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons

The plagioclase series follows according to percentage of anorthite in parentheses [4]. Feldspar rocks are usually used as a source of alkali oxides (Na2 O, K2 O) and alumina (Al2 O3 ) for the preparation of glazes [5]. Suitable choice of feldspar rock can significantly affect the properties of the ceramic body [6], firing temperature and soaking time [10]. The densification of green body, cleanability and the stain resistance of polished sintered ceramic tiles is influenced by particle size distribution of used feldspar rocks [7]. Feldspar rocks may be successfully replaced by LCD waste glass [8]. Wollastonite is very suitable material for acceleration of sintering process in porcelain body. Only 1 wt.% addition of wollastonite is able to decrease firing temperature (about 25°C) in the mixture with kaolin, quartz and potassium feldspar rock [9].

bodies. The addition of an epiclastic rock (20 wt.%) brought significant advantages (better grind ability, lower firing temperature with improved mechanical strength and lower porosity) and disadvantages (increasing of slip viscosity, worse powder compressibility, higher

Sintering of Whiteware Body Depending on Different Fluxing Agents and Binders

Anorthite type of whiteware body on the basis of raw materials mixture of feldspar rock, quartz and calcium aluminate cement (CAC) was developed at the firing temperature of 1300°C. Calcium aluminate cement (substitution of traditional kaolin or quartz) increases the strength of green body (**Figure 1**) and lowers the density due to formation of anorthite in all the fired bodies. An optimal ratio between quartz and feldspar rock for optimal sintering of

Whiteware body based on anorthite was developed from the mixture of ball clay, alumina, quartz, wollastonite and magnesia mixture. Sintered whiteware body (1220°C) has approximately two times higher modulus of rupture (110 MPa) than traditional porcelain body based on mullite due to lower content of glassy phase (only 30% for anorthitic whiteware body) [26]. Deflocculation of raw materials mixture based on calcium aluminate cement for the production of whiteware body with low porosity is necessary [27]. Carboxylic acids [28], polyethylene glycol, polyacrylate derivatives and aqueous solutions of sodium carboxylate [29] for optimalization of rheological properties of aluminous cement pastes were

Direct sintering is very effective method how to decrease the energy consumption during the firing of porcelain. Direct sintering reduced total processing time by ~50% and also lowered

For the description of sintering process, the thermodilatometrical analysis and sintering temperature are used primarily. Sintering temperature is defined as temperature when the fired

**Figure 1.** Variation of flexural strength of green body with hydration time. F, feldspar; Q, quartz; A, calcium aluminate

O3

) [23, 24].

http://dx.doi.org/10.5772/68082

107

firing shrinkage, and a darker color of the body due to high amounts of Fe2

the body was found (**Figure 2**) [25].

the sintering temperature from 1200 to 1175°C [30].

body has water absorption exactly 2%.

tested.

cement (CAC) [25].

Bone ash is fluxing agent for artistic porcelain especially known as bone china. The amount of bone ash in the raw material mixture of bone china is about 50% [11]. Bone ash (cattle bones calcined at around 1000°C) consists predominantly of hydroxyapatite. The reactions of bone ash in porcelain body were studied in detail in Refs. [12, 13]. Bone ash—fluxing agent for bone porcelains (bone china)—is usually produced by the calcination of bovine bones at the temperature of 1100°C. The melting point of bone ash is about 1670°C [14]. The mineralogical composition of bone ash consists of tricalcium phosphate in the form of hydroxyapatite Ca5 (OH)(PO<sup>4</sup> ) 3 .

Very useful fluxing agent for sintered ceramic body production is zeolite, which is able to accelerate the sintering process very intensively. Zeolite is a natural mineral with exceptional physical properties that follow from its specific crystal structure. The latter consists of a 3D lattice of silicate tetrahedrons (SiO<sup>4</sup> )4− mutually connected by oxygen atoms, with part of silicon atoms replaced by aluminum atoms (AlO<sup>4</sup> )5−. Zeolite has a wide range of applications in agriculture, breeding, civil engineering, protection of environment, wastewater purification, and in various industrial sectors. In civil engineering, it began to be used as a partial replacement of cement in the production of concrete [15–19]. Different Italian low‐cost natural zeolitic rocks as a substitute of feldspar rocks in porcelain raw materials mixture were investigated. Zeolitic rocks increased the slip viscosity during wet grinding with a coarser grain size distribution. The technological properties (strength, porosity, resistance etc.) of zeolite‐based porcelain bodies are similar to current traditional porcelain bodies made in the system kaolin—feldspar rock—quartz [20]. The aim of the study [21] was to investigate the effect of natural zeolite addition on the sintering kinetics. Clinoptilolite, which is a type of natural zeolite, was added partially or fully in replacement of quartz at selected electro‐porcelain composition. It was found that the sintering activation energy decreased with increasing zeolite addition. Replacement of quartz with zeolite decreases activation energy for the start of sintering process in electro‐porcelain body—firing temperature (about 50–100°C) and soaking time were reduced. In the study [22], the effect of natural zeolite addition on the electrical properties of porcelain bodies was investigated. The resistivity of samples increased at 50°C temperature after zeolite addition, while it was decreasing after zeolite addition at higher temperatures. The resistivity of samples depends on sintering temperature. Low‐cost naturally occurring mixtures of feldspar and zeolite occurring in epiclastic rocks were promising substitutes for conventional quartz‐feldspathic fluxes in ceramic bodies. Different epiclastic outcrops, with a different zeolite‐to‐feldspar ratio, were tested in porcelain stoneware bodies. The addition of an epiclastic rock (20 wt.%) brought significant advantages (better grind ability, lower firing temperature with improved mechanical strength and lower porosity) and disadvantages (increasing of slip viscosity, worse powder compressibility, higher firing shrinkage, and a darker color of the body due to high amounts of Fe2 O3 ) [23, 24].

The plagioclase series follows according to percentage of anorthite in parentheses [4].

for the preparation of glazes [5]. Suitable choice of feldspar rock can significantly affect the properties of the ceramic body [6], firing temperature and soaking time [10]. The densification of green body, cleanability and the stain resistance of polished sintered ceramic tiles is influenced by particle size distribution of used feldspar rocks [7]. Feldspar rocks may be successfully replaced by LCD waste glass [8]. Wollastonite is very suitable material for acceleration of sintering process in porcelain body. Only 1 wt.% addition of wollastonite is able to decrease firing temperature (about 25°C) in the mixture with kaolin,

Bone ash is fluxing agent for artistic porcelain especially known as bone china. The amount of bone ash in the raw material mixture of bone china is about 50% [11]. Bone ash (cattle bones calcined at around 1000°C) consists predominantly of hydroxyapatite. The reactions of bone ash in porcelain body were studied in detail in Refs. [12, 13]. Bone ash—fluxing agent for bone porcelains (bone china)—is usually produced by the calcination of bovine bones at the temperature of 1100°C. The melting point of bone ash is about 1670°C [14]. The mineralogical composition of bone ash consists of tricalcium phosphate in the form of hydroxyapatite

Very useful fluxing agent for sintered ceramic body production is zeolite, which is able to accelerate the sintering process very intensively. Zeolite is a natural mineral with exceptional physical properties that follow from its specific crystal structure. The latter consists of

in agriculture, breeding, civil engineering, protection of environment, wastewater purification, and in various industrial sectors. In civil engineering, it began to be used as a partial replacement of cement in the production of concrete [15–19]. Different Italian low‐cost natural zeolitic rocks as a substitute of feldspar rocks in porcelain raw materials mixture were investigated. Zeolitic rocks increased the slip viscosity during wet grinding with a coarser grain size distribution. The technological properties (strength, porosity, resistance etc.) of zeolite‐based porcelain bodies are similar to current traditional porcelain bodies made in the system kaolin—feldspar rock—quartz [20]. The aim of the study [21] was to investigate the effect of natural zeolite addition on the sintering kinetics. Clinoptilolite, which is a type of natural zeolite, was added partially or fully in replacement of quartz at selected electro‐porcelain composition. It was found that the sintering activation energy decreased with increasing zeolite addition. Replacement of quartz with zeolite decreases activation energy for the start of sintering process in electro‐porcelain body—firing temperature (about 50–100°C) and soaking time were reduced. In the study [22], the effect of natural zeolite addition on the electrical properties of porcelain bodies was investigated. The resistivity of samples increased at 50°C temperature after zeolite addition, while it was decreasing after zeolite addition at higher temperatures. The resistivity of samples depends on sintering temperature. Low‐cost naturally occurring mixtures of feldspar and zeolite occurring in epiclastic rocks were promising substitutes for conventional quartz‐feldspathic fluxes in ceramic bodies. Different epiclastic outcrops, with a different zeolite‐to‐feldspar ratio, were tested in porcelain stoneware

O, K2

)4− mutually connected by oxygen atoms, with part of

)5−. Zeolite has a wide range of applications

O) and alumina (Al2

O3 )

Feldspar rocks are usually used as a source of alkali oxides (Na2

quartz and potassium feldspar rock [9].

a 3D lattice of silicate tetrahedrons (SiO<sup>4</sup>

silicon atoms replaced by aluminum atoms (AlO<sup>4</sup>

Ca5

(OH)(PO<sup>4</sup>

106 Sintering of Functional Materials

) 3 . Anorthite type of whiteware body on the basis of raw materials mixture of feldspar rock, quartz and calcium aluminate cement (CAC) was developed at the firing temperature of 1300°C. Calcium aluminate cement (substitution of traditional kaolin or quartz) increases the strength of green body (**Figure 1**) and lowers the density due to formation of anorthite in all the fired bodies. An optimal ratio between quartz and feldspar rock for optimal sintering of the body was found (**Figure 2**) [25].

Whiteware body based on anorthite was developed from the mixture of ball clay, alumina, quartz, wollastonite and magnesia mixture. Sintered whiteware body (1220°C) has approximately two times higher modulus of rupture (110 MPa) than traditional porcelain body based on mullite due to lower content of glassy phase (only 30% for anorthitic whiteware body) [26]. Deflocculation of raw materials mixture based on calcium aluminate cement for the production of whiteware body with low porosity is necessary [27]. Carboxylic acids [28], polyethylene glycol, polyacrylate derivatives and aqueous solutions of sodium carboxylate [29] for optimalization of rheological properties of aluminous cement pastes were tested.

Direct sintering is very effective method how to decrease the energy consumption during the firing of porcelain. Direct sintering reduced total processing time by ~50% and also lowered the sintering temperature from 1200 to 1175°C [30].

For the description of sintering process, the thermodilatometrical analysis and sintering temperature are used primarily. Sintering temperature is defined as temperature when the fired body has water absorption exactly 2%.

**Figure 1.** Variation of flexural strength of green body with hydration time. F, feldspar; Q, quartz; A, calcium aluminate cement (CAC) [25].

**Figure 2.** Water adsorption with various compositions in fired bodies with 20% of CAC [25].

## **2. Sintering of whiteware body depending on fluxing agent (feldspar rocks, bone ash, zeolite)**

temperature (**Figure 3**). The most intensive sintering activity of the pure feldspar rock body shows potassium‐sodium feldspar rock F‐KNa—dry pressed test samples have the lowest water absorption, the highest bulk density and modulus of rupture in all firing temperatures in the range of firing at temperatures 1120–1210°C. Sodium‐calcium feldspar rock F‐NaCa begins sintering at much higher firing temperatures. Sintering temperature (**Figure 3**) of tested alkali feldspar rocks F‐KNa and F‐K is significantly lower than oligoclase type of feld-

**Table 1.** Chemical composition of used feldspar rocks and zeolite in weight% (LOI = loss of ignition) and the equivalent

**F‐KNa F‐K F‐NaCa Zeolite**

Sintering of Whiteware Body Depending on Different Fluxing Agents and Binders

http://dx.doi.org/10.5772/68082

109

SiO2 79.76 70.96 66.67 68.20

O<sup>3</sup> 12.37 16.10 20.11 12.40

O<sup>3</sup> 0.42 0.10 0.26 1.40 TiO<sup>2</sup> 0.05 0.04 0.04 – CaO 0.48 0.30 4.23 3.30 MgO 0.10 0.06 0.07 1.00

O 3.35 10.36 0.83 2.80

O 2.67 1.90 7.13 1.00 LOI 0.80 0.20 0.74 – *d*(0.5) [μm] 20.8 18.4 16.6 20.0

The mixtures of feldspar rocks with kaolin (40 wt.%)—samples FK‐KNa, FK‐K, FK‐NaCa totally change (increase) the sintering temperatures (**Table 2**) of alkali feldspar rocks F‐K and F‐KNa. The most intensive fluxing agent in case of pure feldspar rock body (F‐KNa) exhibits the lowest sintering activity in the mixture with kaolin with the highest sintering temperature. This fact is confirmed according to thermodilatometrical curves (**Figure 4**). Conversely, the mixture with kaolin decreases the sintering temperature of oligoclase F‐NaCa with the highest content of pure feldspars. The sintering temperature of F‐NaCa mixture with kaolin is

The difference between the sintering of pure feldspar rocks (F‐KNa, F‐K, F‐NaCa) and mixtures of feldspar rocks with kaolin (FK‐KNa, FK‐K, FK‐NaCa) is evident from the thermodilatometrical curves (**Figure 4**). The highest content of quartz and muscovite in feldspar rock F‐KNa caused high expansion of the body during firing in the range of 200–900°C in comparison with other tested samples based on feldspar rocks F‐NaCa and F‐K. Dry pressed body based on pure feldspar rock F‐KNa shows the best sinterability of all compared feldspar rocks with maximal firing shrinkage (about 5%—**Figure 4**). Very significant is quartz transformation at the temperature 573°C on cooling part of thermodilatometrical curves (**Figure 4**) depending on quartz content (**Table 1**) in individual tested feldspar rocks. The quartz transformation

is most visible for F‐KNa feldspar rock with maximal (55%) content of quartz.

spar rock F‐NaCa.

mean spherical diameter *d*(0.5).

Al2

Fe<sup>2</sup>

K2

Na2

lower (about 20°C) than pure feldspar rock F‐NaCa.

Sintering and melting of feldspar rocks depend on many aspects, such as the fineness of milling (granulometry), the rate of heating and finally the content of alkali oxides, because it directly creates the melting effect. Very useful is to compare the sintering activity of different typical feldspar rocks with different content of pure K‐feldspar, Na‐feldspar and Ca‐feldspar using for the industrial production of whitewares. The comparison is performed for pure feldspar rocks and for mixtures of feldspar rocks with kaolin. For the comparison, next feldspar rocks were used:


The chemical composition of compared feldspar rocks (**Table 1**) reflects their mineralogical composition and volume of different types of pure feldspars (microcline, albite, anorthite). Granulometry of industrially milled feldspar rocks (the equivalent mean spherical diameter of particles *d*(0.5) in **Table 1**) is very similar and does not affect the presented results.

Sintering activity of dry pressed test samples based on tested pure feldspar rocks (**Table 1**) was determined according to dependence of water absorption (EN ISO 10545) on the firing


**Table 1.** Chemical composition of used feldspar rocks and zeolite in weight% (LOI = loss of ignition) and the equivalent mean spherical diameter *d*(0.5).

**2. Sintering of whiteware body depending on fluxing agent (feldspar** 

**Figure 2.** Water adsorption with various compositions in fired bodies with 20% of CAC [25].

Sintering and melting of feldspar rocks depend on many aspects, such as the fineness of milling (granulometry), the rate of heating and finally the content of alkali oxides, because it directly creates the melting effect. Very useful is to compare the sintering activity of different typical feldspar rocks with different content of pure K‐feldspar, Na‐feldspar and Ca‐feldspar using for the industrial production of whitewares. The comparison is performed for pure feldspar rocks and for mixtures of feldspar rocks with kaolin. For the comparison, next feldspar

• Sodium‐potassium feldspar rock F‐KNa with mineralogical composition: K‐feldspar (microcline) 20.0%, Na‐feldspar (albite) 22.6%, Ca‐feldspar (anorthite) 2.4% and quartz 55.0%. • Potassium feldspar rock F‐K with mineralogical composition: K‐feldspar (microcline) 57.2%, Na‐feldspar (albite) 16.0%, Ca‐feldspar (anorthite) 1.5%, quartz 21.3% and mica

• Sodium‐calcium feldspar rock F‐NaCa with mineralogical composition: Na‐feldspar (albite)

The chemical composition of compared feldspar rocks (**Table 1**) reflects their mineralogical composition and volume of different types of pure feldspars (microcline, albite, anorthite). Granulometry of industrially milled feldspar rocks (the equivalent mean spherical diameter

Sintering activity of dry pressed test samples based on tested pure feldspar rocks (**Table 1**) was determined according to dependence of water absorption (EN ISO 10545) on the firing

60.3%, Ca‐feldspar (anorthite) 21.0%, quartz 13.8% and mica (muscovite) 4.9%.

of particles *d*(0.5) in **Table 1**) is very similar and does not affect the presented results.

**rocks, bone ash, zeolite)**

108 Sintering of Functional Materials

rocks were used:

(muscovite) 4.0%.

temperature (**Figure 3**). The most intensive sintering activity of the pure feldspar rock body shows potassium‐sodium feldspar rock F‐KNa—dry pressed test samples have the lowest water absorption, the highest bulk density and modulus of rupture in all firing temperatures in the range of firing at temperatures 1120–1210°C. Sodium‐calcium feldspar rock F‐NaCa begins sintering at much higher firing temperatures. Sintering temperature (**Figure 3**) of tested alkali feldspar rocks F‐KNa and F‐K is significantly lower than oligoclase type of feldspar rock F‐NaCa.

The mixtures of feldspar rocks with kaolin (40 wt.%)—samples FK‐KNa, FK‐K, FK‐NaCa totally change (increase) the sintering temperatures (**Table 2**) of alkali feldspar rocks F‐K and F‐KNa. The most intensive fluxing agent in case of pure feldspar rock body (F‐KNa) exhibits the lowest sintering activity in the mixture with kaolin with the highest sintering temperature. This fact is confirmed according to thermodilatometrical curves (**Figure 4**). Conversely, the mixture with kaolin decreases the sintering temperature of oligoclase F‐NaCa with the highest content of pure feldspars. The sintering temperature of F‐NaCa mixture with kaolin is lower (about 20°C) than pure feldspar rock F‐NaCa.

The difference between the sintering of pure feldspar rocks (F‐KNa, F‐K, F‐NaCa) and mixtures of feldspar rocks with kaolin (FK‐KNa, FK‐K, FK‐NaCa) is evident from the thermodilatometrical curves (**Figure 4**). The highest content of quartz and muscovite in feldspar rock F‐KNa caused high expansion of the body during firing in the range of 200–900°C in comparison with other tested samples based on feldspar rocks F‐NaCa and F‐K. Dry pressed body based on pure feldspar rock F‐KNa shows the best sinterability of all compared feldspar rocks with maximal firing shrinkage (about 5%—**Figure 4**). Very significant is quartz transformation at the temperature 573°C on cooling part of thermodilatometrical curves (**Figure 4**) depending on quartz content (**Table 1**) in individual tested feldspar rocks. The quartz transformation is most visible for F‐KNa feldspar rock with maximal (55%) content of quartz.

**Figure 3.** Water absorption *E* depending on the firing temperature. Determination of sintering temperature (*E* = 2%).


**Figure 4.** Thermodilatometric curves of pure feldspar rocks (F‐K, F‐KNa, F‐NaCa) and the mixtures of feldspar rocks

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**Figure 5.** XRD patterns of sintered feldspar rocks at sintering temperature: M, microcline; Al, albite; Q, quartz; A, anorthite.

with kaolin (FK‐K, FK‐KNa, FK‐NaCa) (10°C/min without soaking time on the maximal temperature).

**Table 2.** Sintering temperatures of tested samples based on different feldspar rocks and mixtures of kaolin (60%) with feldspar rocks (40%) or bone ash (FK‐B).

The feldspar rock F‐KNa with the lowest sintering temperature based on microcline and albite is typical by the quickest disappearing of feldspars during the sintering. Sintering temperature (1190°C) means the existence of only quartz and amorphous glassy phase without any feldspars (**Figure 5**). Quartz, amorphous glassy phase, and microcline are represented in the body F‐K after the firing at sintering temperature (1205°C). It is not possible to find an explanation of this fact in granulometry parameters of used feldspar rocks, which influence sintering and melting of feldspars very much, but in the equilibrium phase diagrams (**Figure 6**). Mixed sodium‐potassium feldspar rock generated low melting eutectic melts, which accelerate the sintering and melting process of feldspars. It is surprising that leucite generating during the potassium feldspars melting according to theoretical assumptions [4] is not detected even in sintered body F‐KNa or sintered body F‐K, both based on the potassium feldspar microcline. After the firing of F‐NaCa sample at sintering temperature (1275°C), the body contains anorthite (calcium feldspar) with high theoretical melting temperature of about 1550°C [4] and albite (**Figure 5**).

The more intensive fluxing agent than feldspar rocks for the sintering process in the system kaolin‐fluxing agent is bone ash (**Figure 4**)—the mixture containing bone ash FK‐B (**Table 2**) shows sintering temperature 1200°C. That is about 50°C lower compared with the most inten ‐ sive feldspar rock‐based mixture (FK‐K) with potassium feldspar rock F‐K containing 75% of pure microcline (**Table 2**). After the exceeding, the temperature 1200°C is visible intensive Sintering of Whiteware Body Depending on Different Fluxing Agents and Binders http://dx.doi.org/10.5772/68082 111

**Figure 4.** Thermodilatometric curves of pure feldspar rocks (F‐K, F‐KNa, F‐NaCa) and the mixtures of feldspar rocks with kaolin (FK‐K, FK‐KNa, FK‐NaCa) (10°C/min without soaking time on the maximal temperature).

The feldspar rock F‐KNa with the lowest sintering temperature based on microcline and albite is typical by the quickest disappearing of feldspars during the sintering. Sintering temperature (1190°C) means the existence of only quartz and amorphous glassy phase without any feldspars (**Figure 5**). Quartz, amorphous glassy phase, and microcline are represented in the body F‐K after the firing at sintering temperature (1205°C). It is not possible to find an explanation of this fact in granulometry parameters of used feldspar rocks, which influence sintering and melting of feldspars very much, but in the equilibrium phase diagrams (**Figure 6**). Mixed sodium‐potassium feldspar rock generated low melting eutectic melts, which accelerate the sintering and melting process of feldspars. It is surprising that leucite generating during the potassium feldspars melting according to theoretical assumptions [4] is not detected even in sintered body F‐KNa or sintered body F‐K, both based on the potassium feldspar microcline. After the firing of F‐NaCa sample at sintering temperature (1275°C), the body contains anorthite (calcium feldspar) with high theoretical melting temperature of about 1550°C

**Table 2.** Sintering temperatures of tested samples based on different feldspar rocks and mixtures of kaolin (60%) with

FK‐B 1200

**Figure 3.** Water absorption *E* depending on the firing temperature. Determination of sintering temperature (*E* = 2%).

**Mixture Sintering temperature (°C) Mixture Sintering temperature (°C)**

F‐K 1205 FK‐K 1250 (+50) F‐NaCa 1275 FK‐NaCa 1255 (−20) F‐KNa 1190 FK‐KNa 1285 (+95)

The more intensive fluxing agent than feldspar rocks for the sintering process in the system kaolin‐fluxing agent is bone ash (**Figure 4**)—the mixture containing bone ash FK‐B (**Table 2**) shows sintering temperature 1200°C. That is about 50°C lower compared with the most inten ‐ sive feldspar rock‐based mixture (FK‐K) with potassium feldspar rock F‐K containing 75% of pure microcline (**Table 2**). After the exceeding, the temperature 1200°C is visible intensive

[4] and albite (**Figure 5**).

feldspar rocks (40%) or bone ash (FK‐B).

110 Sintering of Functional Materials

**Figure 5.** XRD patterns of sintered feldspar rocks at sintering temperature: M, microcline; Al, albite; Q, quartz; A, anorthite.

samples according to fluxing agent utilization (zeolite vs. feldspar rock F‐KNa) are shown in **Figure 7**. During the firing, there is evident (**Figure 7**) that zeolite (in mixture Z) is more intensive fluxing agent compared to feldspar rock F‐KNa (mixture F) for the creation of sintered body with low porosity. High firing shrinkage is typical for the sintering—the raw materials mixture not C, but Z with zeolite content starts intensive shrinking from temperature of about 900°C. Compared mixture F based on traditional ceramic fluxin agent ‐ potassium feldspar

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The quartz transformation at 573°C (change of the fired body volume) is visible on cooling part of thermodilatometric curve of the mixture F (**Figure 7**) due to high portion of quartz in the mixture F based on the mixture kaolin‐quartz‐feldspar. This phenomenon is not presented on the cooling curve of the sintered body Z based on zeolite—the raw material mixture not contains quartz, which is advantageous for lower relative expansion (coefficient of linear ther-

Sintering temperature of tested samples based on zeolite is 1180°C, which is about 100°C lower than for the sample based on standard flux feldspar rock F‐K (mixture F). From the picture (**Figure 7**), there is evident different coefficient of linear thermal expansion α (in the

The sintered body based on zeolite shows the lower coefficient of thermal expansion α compared with feldspar sample due to the formation of anorthite in the sample Z and absence of quartz.

**Figure 7.** Thermodilatometric analysis of compared samples with different fluxing agent (zeolite vs. feldspar rock F‐KNa)

rock F‐K ‐ starts the sintering process at a higher temperature (about 1100 °C).

temperature range of 30–500°C) of both compared sintered bodies (Z vs. F):

mal expansion) of sintered body Z (**Figure 7**).

• feldspar + quartz body: α30–500°C (F) = 70 × 10−7 K−1

• zeolite body: α30–500°C (Z) = 48 × 10−7 K−1

during the firing (5°C/min without soaking time).

**Figure 6.** Phase diagram NaAlSi3 O8 —KalSi3 O8 —CaAl2 Si2 O8 and theoretical melting temperature of used feldspar rocks (1: F‐KNa, 2: F‐K, 3: F‐NaCa).

bloating of the bone ash bodies, which is typical by creating of secondary porosity and increasing in water absorption (**Figure 4**).

Different mineralogical composition between feldspar rocks and bone ash‐based porcelain sintered bodies is possible to document according to XRD analyses. Traditional porcelain with high content of feldspar rocks in the raw materials mixture contains mullite and quartz as main mineralogical phases. Mineralogical composition of porcelain body based on bone ash is totally different—typical is high content of β‐tricalcium phosphate and anorthite. Bone ash in bone porcelain bodies decomposes into β‐tricalcium phosphate Ca3 (PO<sup>4</sup> ) 2 , lime CaO and water at around 775°C according to Eq. (1) [14]:

$$\text{Ca}\_{10}(\text{PO}\_4)\_6(\text{OH})\_2 \rightarrow 3\cdot \text{β-Ca}\_3(\text{PO}\_4)\_2 + \text{CaO} + \text{H}\_2\text{O} \tag{1}$$

Lime reacts with metakaolin from clay relicts to form of anorthite [CaAl2 Si2 O8 ] according to Eq. (2) [14]:

$$\text{Al}\_2\text{O}\_3 \cdot 2\,\text{SiO}\_2 + \text{CaO} \rightarrow \text{CaAl}\_2\text{Si}\_2\text{O}\_8\tag{2}$$

Eutectic composition in the ternary system of bone china (Ca3 (PO<sup>4</sup> )2 —CaAl2 Si2 O8 —SiO2 ) is about 11% tricalcium phosphate, 51% anorthite and 38% silica with a melting temperature of 1290 ± 5°C [14].

Zeolite rock was investigated as a fluxing agent for sintered ceramic body and its effect in the sintering process. The thermodilatometric heating and cooling curves dL/L0 of two different samples according to fluxing agent utilization (zeolite vs. feldspar rock F‐KNa) are shown in **Figure 7**. During the firing, there is evident (**Figure 7**) that zeolite (in mixture Z) is more intensive fluxing agent compared to feldspar rock F‐KNa (mixture F) for the creation of sintered body with low porosity. High firing shrinkage is typical for the sintering—the raw materials mixture not C, but Z with zeolite content starts intensive shrinking from temperature of about 900°C. Compared mixture F based on traditional ceramic fluxin agent ‐ potassium feldspar rock F‐K ‐ starts the sintering process at a higher temperature (about 1100 °C).

The quartz transformation at 573°C (change of the fired body volume) is visible on cooling part of thermodilatometric curve of the mixture F (**Figure 7**) due to high portion of quartz in the mixture F based on the mixture kaolin‐quartz‐feldspar. This phenomenon is not presented on the cooling curve of the sintered body Z based on zeolite—the raw material mixture not contains quartz, which is advantageous for lower relative expansion (coefficient of linear thermal expansion) of sintered body Z (**Figure 7**).

Sintering temperature of tested samples based on zeolite is 1180°C, which is about 100°C lower than for the sample based on standard flux feldspar rock F‐K (mixture F). From the picture (**Figure 7**), there is evident different coefficient of linear thermal expansion α (in the temperature range of 30–500°C) of both compared sintered bodies (Z vs. F):


bloating of the bone ash bodies, which is typical by creating of secondary porosity and

Different mineralogical composition between feldspar rocks and bone ash‐based porcelain sintered bodies is possible to document according to XRD analyses. Traditional porcelain with high content of feldspar rocks in the raw materials mixture contains mullite and quartz as main mineralogical phases. Mineralogical composition of porcelain body based on bone ash is totally different—typical is high content of β‐tricalcium phosphate and anorthite. Bone ash in bone

Ca10 (PO<sup>4</sup> )<sup>6</sup> (OH )<sup>2</sup> → 3 β‐Ca3 (PO<sup>4</sup> )<sup>2</sup> + CaO + H<sup>2</sup> O (1)

Al2 O<sup>3</sup> ⋅ 2 SiO2 + CaO → CaAl2 Si2 O8 (2)

about 11% tricalcium phosphate, 51% anorthite and 38% silica with a melting temperature of

Zeolite rock was investigated as a fluxing agent for sintered ceramic body and its effect in the sintering process. The thermodilatometric heating and cooling curves dL/L0 of two different

(PO<sup>4</sup> ) 2

and theoretical melting temperature of used feldspar rocks

(PO<sup>4</sup> )2 —CaAl2

, lime CaO and water at

] according to

) is

Si2 O8

> Si2 O8 —SiO2

increasing in water absorption (**Figure 4**).

O8 —KalSi3 O8 —CaAl2 Si2 O8

**Figure 6.** Phase diagram NaAlSi3

(1: F‐KNa, 2: F‐K, 3: F‐NaCa).

112 Sintering of Functional Materials

around 775°C according to Eq. (1) [14]:

Eq. (2) [14]:

1290 ± 5°C [14].

porcelain bodies decomposes into β‐tricalcium phosphate Ca3

Eutectic composition in the ternary system of bone china (Ca3

Lime reacts with metakaolin from clay relicts to form of anorthite [CaAl2

The sintered body based on zeolite shows the lower coefficient of thermal expansion α compared with feldspar sample due to the formation of anorthite in the sample Z and absence of quartz.

**Figure 7.** Thermodilatometric analysis of compared samples with different fluxing agent (zeolite vs. feldspar rock F‐KNa) during the firing (5°C/min without soaking time).

Important technical property of anorthite is its low coefficient of linear thermal expansion of 48.2 × 10−6 K−1[31] (mullite 60 × 10−7 K−1 [32]). The mineralogical composition of both bodies after firing in both cases is characterized by the existence of mullite and glass phase. The sintered body (fired at 1200°C—mixture Z or 1300°C—mixture F, respectively) based on feldspar and quartz (mixture F) also contains quartz, and the body made from zeolite contains anorthite and cristobalite.

**Mixture Content (%‐mass)**

temperatures 30–500°C.

**Table 3.** Composition of raw material mixtures (test samples).

K 25% kaolin + 50% F‐KNa + 25% quartz sand + 0.35% sodium hexametaphosphate (deflocculant) CAC 25% CAC + 50% F‐KNa + 25% quartz sand + 0.35% sodium hexametaphosphate (deflocculant)

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**Figure 9.** Thermodilatometric analysis of kaolin (K) and calcium aluminate cement (CAC)‐based bodies during the firing (1280°C, 3°C/min without soaking time). Determination of the coefficient of linear thermal expansion in the range of

**Figure 10.** XRD of fired bodies based on different binder kaolin or CAC (M, mullite; Q, quartz; A, anorthite).

Sintered body based on zeolite (mixture Z in **Figure 8**) as a fluxing agent not creates white body, which is typical for sintered body of the mixture F based on F‐KNa feldspar rock (**Figure 8**). This situation corresponds to chemical composition of natural zeolite with higher content of Fe<sup>2</sup> O3 (**Table 1**).

**Figure 8.** Color of sintered bodies with water absorption below 2%.

## **3. The effect of calcium aluminate cement as a binder for the sintering of whiteware bodies**

The sintering process of whiteware (porcelain) body is affected by the used binder—we can use traditional plastic material (kaolin) or calcium aluminate cement (CAC) according to latest research [25, 33]. Comparison of the properties of both types (**Table 3**) of porcelains made by pressing from dry granulate is documented.

The difference in sintering process of two whiteware bodies with different binder kaolin vs. CAC (**Table 3**) is documented according to thermodilatometric curves (**Figure 9**). The sintering activity of both compared mixtures is very different when the firing temperature exceeds 1200°C—the system based on CAC (mixture CAC) is more able to sinter—we can observe higher firing shrinkage.

Significant decrease of the coefficient of linear thermal expansion in the temperature range 30–500°C is evident (**Figure 9**) when calcium aluminate cement CAC is used as binder compared with kaolin based body. The explanation of this fact we can find in the formation of anorthite in the CAC‐based sample (**Figure 10**). The fired body based on kaolin also contains mullite and quartz as a main mineralogical phases. Anorthite exhibits lower coefficient of linear thermal expansion of 48.2 × 10−6 K−1 [31] than mullite 60 × 10−6 K−1 [32].


**Table 3.** Composition of raw material mixtures (test samples).

Important technical property of anorthite is its low coefficient of linear thermal expansion of 48.2 × 10−6 K−1[31] (mullite 60 × 10−7 K−1 [32]). The mineralogical composition of both bodies after firing in both cases is characterized by the existence of mullite and glass phase. The sintered body (fired at 1200°C—mixture Z or 1300°C—mixture F, respectively) based on feldspar and quartz (mixture F)

Sintered body based on zeolite (mixture Z in **Figure 8**) as a fluxing agent not creates white body, which is typical for sintered body of the mixture F based on F‐KNa feldspar rock (**Figure 8**). This situation corresponds to chemical composition of natural zeolite with higher

**3. The effect of calcium aluminate cement as a binder for the sintering of** 

The sintering process of whiteware (porcelain) body is affected by the used binder—we can use traditional plastic material (kaolin) or calcium aluminate cement (CAC) according to latest research [25, 33]. Comparison of the properties of both types (**Table 3**) of porcelains made

The difference in sintering process of two whiteware bodies with different binder kaolin vs. CAC (**Table 3**) is documented according to thermodilatometric curves (**Figure 9**). The sintering activity of both compared mixtures is very different when the firing temperature exceeds 1200°C—the system based on CAC (mixture CAC) is more able to sinter—we can observe

Significant decrease of the coefficient of linear thermal expansion in the temperature range 30–500°C is evident (**Figure 9**) when calcium aluminate cement CAC is used as binder compared with kaolin based body. The explanation of this fact we can find in the formation of anorthite in the CAC‐based sample (**Figure 10**). The fired body based on kaolin also contains mullite and quartz as a main mineralogical phases. Anorthite exhibits lower coefficient of linear thermal

also contains quartz, and the body made from zeolite contains anorthite and cristobalite.

content of Fe<sup>2</sup>

114 Sintering of Functional Materials

O3

**whiteware bodies**

higher firing shrinkage.

by pressing from dry granulate is documented.

**Figure 8.** Color of sintered bodies with water absorption below 2%.

expansion of 48.2 × 10−6 K−1 [31] than mullite 60 × 10−6 K−1 [32].

(**Table 1**).

**Figure 9.** Thermodilatometric analysis of kaolin (K) and calcium aluminate cement (CAC)‐based bodies during the firing (1280°C, 3°C/min without soaking time). Determination of the coefficient of linear thermal expansion in the range of temperatures 30–500°C.

**Figure 10.** XRD of fired bodies based on different binder kaolin or CAC (M, mullite; Q, quartz; A, anorthite).


accelerated when calcium aluminate cement is used as a binder instead of kaolin—the bodies can be fired at lower temperatures. Calcium aluminate cement significantly changes mineralogical composition of fired body—anorthite is the main mineralogical phase, mullite is typical phase for standard porcelain bodies made in the system of kaolin‐quartz‐feldspar rock.

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The results were achieved under the project no. LO1408 "AdMaS UP—Advanced Materials, Structures and Technologies", supported by the Ministry of Education, Youth and Sports under the "National Sustainability Programme I" (chapter 2). The results were achieved under the project—the Czech Science Foundation, research project no. P104/13/23051S "Anorthite

Faculty of Civil Engineering, Brno University of Technology, Brno, Czech Republic

ical behavior. Journal of American Ceramic Society. 1998;**81**:3‐20

modulus of porcelain. Industrial Ceramics, 2008;**28**(2):153‐154

ing porcelain bodies. Thermochimica Acta, 2003;**406**(1‐2):199‐206

sanitary ware production. Ceramics International. 2015;**41**(5):7097‐7102

[4] Barth TFW. Feldspars. 1st ed. Bath: John Wiley & Sons; 1969

Ceramic Society, 2012;**32**(10):2095‐2102

[1] Carty WM, Senapati U. Porcelain—raw materials, processing, phase evolution, and mechan-

[2] Rado P. An Introduction to the Technology of Pottery. Oxford: Pergamon Press; 1988

[3] Stubna I, Slavikova J, Vozar L. Relationship between mechanical strength and Young's

[5] Norton, FH. Fine Ceramics Technology and Applications. 1st ed. Malabar: R.E Krieger;

[6] Das S. Kr, Dana K. Differences in densification behaviour of K‐ and Na‐feldspar‐contain-

[7] Alves HJ, Melchiades FG, Boschi AO. Effect of feldspar particle size on the porous microstructure and stain resistance of polished porcelain tiles. Journal of the European

[8] Kim K, Kim K, Hwang J. LCD waste glass as a substitute for feldspar in the porcelain

porcelain body on the basis of aluminous cement" (chapter 3).

Address all correspondence to: sokolar.r@fce.vutbr.cz

**Acknowledgements**

**Author details**

Radomir Sokolar

**References**

1970

**Table 4.** Physicomechanical properties of fired bodies K and CAC depending on firing temperature: Ev—water absor‐ ption, MOR—modulus of rupture, B—bulk density.

The CAC mixture shows significantly higher sintering activity according to measured parameters of porosity—the prepared samples of a mixture CAC have a lower water absorption Ev and higher bulk density B (according to EN ISO 10545) than mixtures based on kaolin (K) after firing at the same temperature (**Table 4**). Higher modulus of rupture MOR of fired bodies is achieved for anorthitic type of body (CAC) compared with mullite whiteware body (K) when MOR values for samples with similar porosity are compared (K‐1280°C and CAC‐1250°C in **Table 4**). Similar results are published in Ref. [25].

## **4. Conclusions**

For sintering and melting of pure feldspar rocks, not just the total content of feldspar components is important, but also the ratio between potassium, sodium and calcium feldspars. At the appropriate ratio, low melting eutectics can be expected to rise, with a melting temperature substantially lower than the theoretical melting temperature of pure feldspars. The presence of calcium feldspar significantly reduces sintering ability and melting of feldspar rocks. Totally different results we can expect for the mixtures of feldspar rocks with the plastic part of whiteware raw materials mixture—kaolin. The reactions between feldspar rocks and kaolin (Al2 O3 , SiO2 ) during the sintering process are the cause of low melting eutectics, which accelerate sintering.

Natural zeolite is very intensive fluxing agent for ceramic technology. Using zeolite we can reduce the sintering temperature of the body of about 100°C, compared with traditional ceramic fluxing agent—potassium‐sodium feldspar rock F‐KNa. The sintered body (with water absorption below 2%) based on zeolite has lower coefficient of linear thermal expansion. The presence of zeolite in raw materials mixture significantly changes mineralogical composition of fired whiteware body—mullite, anorthite and cristobalite are the main mineralogical phases instead of mullite and quartz, which are typical for a standard whiteware bodies made from raw material mixtures based on kaolin, quartz and feldspar. The limiting factor for the use of natural zeolite as a flux for whiteware is its coloring effect.

Calcium aluminate cement CAC with high content of Al2 O3 (70%) in the raw materials mixture for whiteware production is suitable alternative to kaolin—higher strength of green and fired body, more intensive whiteness of body after firing and lower coefficient of linear thermal expansion is possible to expect using CAC. The sintering activity of the whiteware body is accelerated when calcium aluminate cement is used as a binder instead of kaolin—the bodies can be fired at lower temperatures. Calcium aluminate cement significantly changes mineralogical composition of fired body—anorthite is the main mineralogical phase, mullite is typical phase for standard porcelain bodies made in the system of kaolin‐quartz‐feldspar rock.

## **Acknowledgements**

The results were achieved under the project no. LO1408 "AdMaS UP—Advanced Materials, Structures and Technologies", supported by the Ministry of Education, Youth and Sports under the "National Sustainability Programme I" (chapter 2). The results were achieved under the project—the Czech Science Foundation, research project no. P104/13/23051S "Anorthite porcelain body on the basis of aluminous cement" (chapter 3).

## **Author details**

The CAC mixture shows significantly higher sintering activity according to measured parameters of porosity—the prepared samples of a mixture CAC have a lower water absorption Ev and higher bulk density B (according to EN ISO 10545) than mixtures based on kaolin (K) after firing at the same temperature (**Table 4**). Higher modulus of rupture MOR of fired bodies is achieved for anorthitic type of body (CAC) compared with mullite whiteware body (K) when MOR values for samples with similar porosity are compared (K‐1280°C and

**Table 4.** Physicomechanical properties of fired bodies K and CAC depending on firing temperature: Ev—water absor‐

**Ev (%) MOR (MPa) B (kg m−3] Ev (%) MOR (MPa) B (kg m−3]**

For sintering and melting of pure feldspar rocks, not just the total content of feldspar components is important, but also the ratio between potassium, sodium and calcium feldspars. At the appropriate ratio, low melting eutectics can be expected to rise, with a melting temperature substantially lower than the theoretical melting temperature of pure feldspars. The presence of calcium feldspar significantly reduces sintering ability and melting of feldspar rocks. Totally different results we can expect for the mixtures of feldspar rocks with the plastic part of whiteware raw materials mixture—kaolin. The reactions between feldspar rocks and

Natural zeolite is very intensive fluxing agent for ceramic technology. Using zeolite we can reduce the sintering temperature of the body of about 100°C, compared with traditional ceramic fluxing agent—potassium‐sodium feldspar rock F‐KNa. The sintered body (with water absorption below 2%) based on zeolite has lower coefficient of linear thermal expansion. The presence of zeolite in raw materials mixture significantly changes mineralogical composition of fired whiteware body—mullite, anorthite and cristobalite are the main mineralogical phases instead of mullite and quartz, which are typical for a standard whiteware bodies made from raw material mixtures based on kaolin, quartz and feldspar. The limiting

for whiteware production is suitable alternative to kaolin—higher strength of green and fired body, more intensive whiteness of body after firing and lower coefficient of linear thermal expansion is possible to expect using CAC. The sintering activity of the whiteware body is

factor for the use of natural zeolite as a flux for whiteware is its coloring effect.

Calcium aluminate cement CAC with high content of Al2

) during the sintering process are the cause of low melting eutectics, which

O3

(70%) in the raw materials mixture

CAC‐1250°C in **Table 4**). Similar results are published in Ref. [25].

**1250 1280**

CAC 1.3 58.9 2360 Melting of test samples

K 6.4 29.4 2180 1.9 38.5 2280

**4. Conclusions**

**Sample Firing temperature (°C)**

116 Sintering of Functional Materials

ption, MOR—modulus of rupture, B—bulk density.

kaolin (Al2

O3 , SiO2

accelerate sintering.

Radomir Sokolar

Address all correspondence to: sokolar.r@fce.vutbr.cz

Faculty of Civil Engineering, Brno University of Technology, Brno, Czech Republic

## **References**


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[26] Taskiran MU, Demirkol N, Capoglu A. A new porcelainised stoneware material based on anorthite. Journal of the European Ceramic Society. 2005;**25**(2005):293‐300

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[30] Lerdprom W, Chinnam RK, Jayaseelan DD, Lee WE. Porcelain production by direct sin-

[31] Potuzak M, Solvang M, Dingwell D. Temperature independent thermal expansivities of calcium aluminosilicates melts between 1150 and 1973 K in the system anorthite–wollostanite–gehlenite (An–Wo–Geh): a density model. Geochimica et Cosmochimica Acta.

[32] Camerucci MA, Urretavizcaya G, Castro MS, Cavalieri AL. Electrical properties and thermal expansion of cordierite and cordierite‐mullite materials. Journal of the European

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118 Sintering of Functional Materials


**Section 2**

**Energy Assisted Sintering**

**Energy Assisted Sintering**

**Chapter 6**

**Provisional chapter**

**Sintering and Reactive Sintering by Spark Plasma**

**Sintering and Reactive Sintering by Spark Plasma** 

DOI: 10.5772/intechopen.68871

A wide variety of technological applications, especially in electronics, requires high‐density nanostructured solids, consolidated by sintering from nanoparticles. A new sintering tech‐ nique known as spark plasma sintering (SPS) appears as the only method to reach high densities while preserving the final grain size within the nanometric range, with the added advantage of carrying out the process at significantly lower temperatures and shorter times as compared with the classical processes. Recent studies have revealed that in many cases, SPS can also accomplish the solid‐state reaction to achieve the desired compound, leading to reactive SPS (RSPS). In this chapter, a review of RSPS is presented, focusing particularly

**Keywords:** nanostructured solids, ceramic materials, reactive spark plasma sintering,

For the past two decades, the synthesis and applications of magnetic nanoparticles (MNPs) have gained immense interest in a wide range of technologies, especially in the biomedical field [1–4]. These applications are based on the novel magnetic properties associated with the nanoscale [5]. In the electronics field, nanostructured materials point also to innovative applications [6, 7], particularly in magnetic recording, actuators, and microwave devices. For these applications, however, a powder constituted by MNPs is not suitable; a high den‐ sity, consolidated solid is required. Consolidation by classic sintering methods requires high

> © 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.

**Sintering (SPS)**

and Raul Valenzuela

**Sintering (SPS)**

Raul Valenzuela

**Abstract**

**1. Introduction**

Giulia Franceschin, Nancy Flores‐Martínez, Gabriela Vázquez‐Victorio, Souad Ammar and

Giulia Franceschin, Nancy Flores‐Martínez, Gabriela Vázquez‐Victorio, Souad Ammar

Additional information is available at the end of the chapter

on magnetic oxide materials as functional solids.

solid‐state processing, magnetic properties

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.68871

**Provisional chapter**

## **Sintering and Reactive Sintering by Spark Plasma Sintering (SPS) Sintering (SPS)**

**Sintering and Reactive Sintering by Spark Plasma** 

DOI: 10.5772/intechopen.68871

Giulia Franceschin, Nancy Flores‐Martínez, Gabriela Vázquez‐Victorio, Souad Ammar and Raul Valenzuela Gabriela Vázquez‐Victorio, Souad Ammar and Raul Valenzuela Additional information is available at the end of the chapter

Giulia Franceschin, Nancy Flores‐Martínez,

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.68871

#### **Abstract**

A wide variety of technological applications, especially in electronics, requires high‐density nanostructured solids, consolidated by sintering from nanoparticles. A new sintering tech‐ nique known as spark plasma sintering (SPS) appears as the only method to reach high densities while preserving the final grain size within the nanometric range, with the added advantage of carrying out the process at significantly lower temperatures and shorter times as compared with the classical processes. Recent studies have revealed that in many cases, SPS can also accomplish the solid‐state reaction to achieve the desired compound, leading to reactive SPS (RSPS). In this chapter, a review of RSPS is presented, focusing particularly on magnetic oxide materials as functional solids.

**Keywords:** nanostructured solids, ceramic materials, reactive spark plasma sintering, solid‐state processing, magnetic properties

#### **1. Introduction**

For the past two decades, the synthesis and applications of magnetic nanoparticles (MNPs) have gained immense interest in a wide range of technologies, especially in the biomedical field [1–4]. These applications are based on the novel magnetic properties associated with the nanoscale [5]. In the electronics field, nanostructured materials point also to innovative applications [6, 7], particularly in magnetic recording, actuators, and microwave devices. For these applications, however, a powder constituted by MNPs is not suitable; a high den‐ sity, consolidated solid is required. Consolidation by classic sintering methods requires high

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. © 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.

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons

temperatures and long times—more than 1000°C for several hours—to reach densities above 90%. Such conditions lead to an excessive grain growth, which deteriorate the properties asso‐ ciated with the nanoscale, thus making such consolidation methods impractical.

In this review, we analyze the microstructure formation of the products of chemical reac‐ tions occurring in RSPS, in an attempt to directly produce nanostructured solids starting from the corresponding reactants, that is, an intermediate solid phase, or a mixture of precursors, containing the required elements to form the desired phase. We also discuss the possibility of fabricating nanocomposites, in which the interfaces between the constituting phases can be

Sintering and Reactive Sintering by Spark Plasma Sintering (SPS)

http://dx.doi.org/10.5772/intechopen.68871

125

Focusing on magnetic granular oxide nanostructures, we present successful syntheses with a special emphasis on their microstructure stability and attractive properties of the materi‐ als. We discuss the challenges of producing a dense nanostructured material when reaction and densification do not coincide during the SPS. Case examples in the fields of magneto‐ caloric materials (manganites), soft magnetic materials (garnets), and permanent magnets

We also discuss the limitations of such a technique, in relation to its reducing operating con‐ ditions and propose some alternatives to overcome main drawbacks. Indeed, RSPS is mainly performed using graphite‐made die and punches under dynamic vacuum, creating a reduc‐ ing atmosphere. In the case of oxide materials, this can lead to a partial reduction, sometimes even to a metal contamination, affecting the final physical properties of the consolidated sol‐ ids (electric conductivity, for instance). The replacement of carbon dies by tungsten carbide in

The composition of the environment inside the SPS chamber affects the material's diffusivity during sintering. For this reason, the processes that normally occur during a sintering cycle, such as phase constitution, densification, and grain growth, are strongly affected by the sin‐

Typical materials developed by the SPS technique are refractory metals and interm‐ etallics, oxide and nonoxide ceramics. The particles constituting the powders before consolidation tend to decrease their surface energy by desorption of chemical species, once introduced inside the SPS chamber. The released gas, water, or organic compounds in the atmosphere modify the thermodynamic driving force to surface reduction and

The atmosphere composition inside the pressing tools which contain the sample may dif‐ fer considerably from the atmosphere outside [9], making it difficult to control the sinter‐ ing atmosphere. Sample reduction during RSPS occurs when the thermodynamic conditions

improved by particular tailoring.

(hexaferrites) are specifically addressed.

such materials offers an interesting alternative.

tering atmosphere [13].

• Vacuum (10−4 to 10−5 bar), • Inert gas (up to 1.3 bar), or

• Reducing gas hydrogen‐based mixtures.

sintering.

**2. Reduction reactions during SPS processes**

Commonly, the atmospheres employed for sintering are:

Spark plasma sintering (SPS) [8] has recently been revealed as an extremely efficient sinter‐ ing technique for consolidating nanopowders into high density, nanostructured materials. In practice, the powders are heated in a conductive SPS die at very high rates by the action of electric pulses and maintained under uniaxial pressure (**Figure 1**), leading to their sintering with impressive shorter times and lower temperatures than in conventional methods. In addi‐ tion, to be a more efficient method, it allows a tight control of grain growth, thus permitting the production of nanostructured materials. The principle of SPS and convenient design of the facilities make it attractive for conducting materials. Recent results show, however, that it is equally powerful for nonconductive ones [9].

The current is applied and passed through the graphite die. If the sample is nonconductive, the heat generated inside the walls of the die assists in powder consolidation. If starting pow‐ ders are a conductive material, the current goes through the powder, and the first sparks are in the surface atoms as well as in surface defects. This punctual warming of atoms is known as hot spots. In these zones, the temperature increases thousands of degrees in a very short time, and nucleation and grain growth begin. If the sample is an insulating phase, the electric field associated with the electric pulses has also a strong effect on atomic diffusion, and sintering is enhanced [9].

SPS has been recently examined in a much broader perspective and has gained a strong reputa‐ tion as a versatile method of solid‐state synthesis, not only for sintering, but also for solid‐state reactions, as reported in relevant literature [11, 12]. SPS then becomes reactive SPS (RSPS).

**Figure 1.** Schematics of the vacuum chamber, electrodes, and pressing die of the SPS system (adapted from Ref. [10]).

In this review, we analyze the microstructure formation of the products of chemical reac‐ tions occurring in RSPS, in an attempt to directly produce nanostructured solids starting from the corresponding reactants, that is, an intermediate solid phase, or a mixture of precursors, containing the required elements to form the desired phase. We also discuss the possibility of fabricating nanocomposites, in which the interfaces between the constituting phases can be improved by particular tailoring.

Focusing on magnetic granular oxide nanostructures, we present successful syntheses with a special emphasis on their microstructure stability and attractive properties of the materi‐ als. We discuss the challenges of producing a dense nanostructured material when reaction and densification do not coincide during the SPS. Case examples in the fields of magneto‐ caloric materials (manganites), soft magnetic materials (garnets), and permanent magnets (hexaferrites) are specifically addressed.

We also discuss the limitations of such a technique, in relation to its reducing operating con‐ ditions and propose some alternatives to overcome main drawbacks. Indeed, RSPS is mainly performed using graphite‐made die and punches under dynamic vacuum, creating a reduc‐ ing atmosphere. In the case of oxide materials, this can lead to a partial reduction, sometimes even to a metal contamination, affecting the final physical properties of the consolidated sol‐ ids (electric conductivity, for instance). The replacement of carbon dies by tungsten carbide in such materials offers an interesting alternative.

## **2. Reduction reactions during SPS processes**

The composition of the environment inside the SPS chamber affects the material's diffusivity during sintering. For this reason, the processes that normally occur during a sintering cycle, such as phase constitution, densification, and grain growth, are strongly affected by the sin‐ tering atmosphere [13].

Typical materials developed by the SPS technique are refractory metals and interm‐ etallics, oxide and nonoxide ceramics. The particles constituting the powders before consolidation tend to decrease their surface energy by desorption of chemical species, once introduced inside the SPS chamber. The released gas, water, or organic compounds in the atmosphere modify the thermodynamic driving force to surface reduction and sintering.

Commonly, the atmospheres employed for sintering are:

• Vacuum (10−4 to 10−5 bar),

temperatures and long times—more than 1000°C for several hours—to reach densities above 90%. Such conditions lead to an excessive grain growth, which deteriorate the properties asso‐

Spark plasma sintering (SPS) [8] has recently been revealed as an extremely efficient sinter‐ ing technique for consolidating nanopowders into high density, nanostructured materials. In practice, the powders are heated in a conductive SPS die at very high rates by the action of electric pulses and maintained under uniaxial pressure (**Figure 1**), leading to their sintering with impressive shorter times and lower temperatures than in conventional methods. In addi‐ tion, to be a more efficient method, it allows a tight control of grain growth, thus permitting the production of nanostructured materials. The principle of SPS and convenient design of the facilities make it attractive for conducting materials. Recent results show, however, that it is

The current is applied and passed through the graphite die. If the sample is nonconductive, the heat generated inside the walls of the die assists in powder consolidation. If starting pow‐ ders are a conductive material, the current goes through the powder, and the first sparks are in the surface atoms as well as in surface defects. This punctual warming of atoms is known as hot spots. In these zones, the temperature increases thousands of degrees in a very short time, and nucleation and grain growth begin. If the sample is an insulating phase, the electric field associated with the electric pulses has also a strong effect on atomic diffusion, and sintering

SPS has been recently examined in a much broader perspective and has gained a strong reputa‐ tion as a versatile method of solid‐state synthesis, not only for sintering, but also for solid‐state reactions, as reported in relevant literature [11, 12]. SPS then becomes reactive SPS (RSPS).

**Figure 1.** Schematics of the vacuum chamber, electrodes, and pressing die of the SPS system (adapted from Ref. [10]).

ciated with the nanoscale, thus making such consolidation methods impractical.

equally powerful for nonconductive ones [9].

is enhanced [9].

124 Sintering of Functional Materials


The atmosphere composition inside the pressing tools which contain the sample may dif‐ fer considerably from the atmosphere outside [9], making it difficult to control the sinter‐ ing atmosphere. Sample reduction during RSPS occurs when the thermodynamic conditions are favorable to the imbalance of either one of the following reactions from the right to the left side:

$$\text{metal} + \text{O}\_2 \leftrightharpoons \text{oxide} + \Delta \text{H}^0 \text{l} \tag{1}$$

$$\text{metal} + 2\text{H}\_2\text{O} \leftrightharpoons \text{oxide} + 2\text{H}\_2 + \Delta\text{H}\_2\text{O} \tag{2}$$

$$\text{metal} + \text{2CO}\_2 \leftrightharpoons \text{oxide} + \text{2CO} + \Delta\text{H}\text{O}\_3 \tag{3}$$

Δ*H*<sup>0</sup> 1 , Δ*H*<sup>0</sup> 2 , and Δ*H*<sup>0</sup> 3 are the heat released per O2 mole in each oxidizing reaction (from left to right side). By definition, we can calculate the corresponding changes of free enthalpy:

$$
\Delta G\_\mathbf{l}^0 = -\Delta H\_\mathbf{l}^0 = -RT\ln\left(\frac{p\_{\text{metal}}}{p\_{\text{oxido}}} \cdot P\_{O\_\mathbf{t}}\right) \tag{4}
$$

$$
\Delta G\_2^0 = -\Delta H\_2^0 = -RT\ln\left(\frac{P\_{\text{total}}}{P\_{\text{oxid}}} \cdot \frac{P\_{H\_2O}^2}{P\_{H\_2}^2}\right) \tag{5}
$$

$$
\Delta G\_3^{\text{0}} = -\Delta H\_3^{\text{0}} = -RT\ln\left(\frac{P\_{\text{metal}}}{P\_{\text{ox}\text{tot}}} \cdot \frac{P\_{\text{CO}\_2}^2}{P\_{\text{CO}}^2}\right) \tag{6}
$$

A standard measure for the tendency of a metal (or a chemical element) to oxidize is given by Eqs. (4)–(6) when *P*O2 = 1 ; thus, we obtain Eq. (7), which is the heat released when 1 mole of O2 gas at 1 atm pressure combines with 1 mole of metal to form the oxide in function of temperature *T*:

$$
\Delta G^{\bullet} = -RT \ln \left( \frac{p\_{\text{mstal}}}{p\_{\text{oxide}}} \right) \tag{7}
$$

**Figure 2.** Ellingham‐Richardson diagram [16].

Sintering and Reactive Sintering by Spark Plasma Sintering (SPS)

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127

The graphic representation of Δ*G*<sup>0</sup> = *f* (*T*) is known as the Ellingham‐Richardson diagram (**Figure 2**). It allows the direct evaluation of the relative affinity between the elements and the oxidizing agent. Elements with a lower line in Δ*G*<sup>0</sup> in the diagram have a greater affinity to oxygen. This diagram is useful to understand the thermodynamics of the reactions between the sintering material and the atmosphere, and it can give information about the dissociation temperature, the dissociation pressure, and the effect of reducing agents [14].

Special attention must be given to the pressing tools' material. Standard pressing tools used in the RSPS process are graphite based, often internally covered with carbon sheets or foils, in order to ease the removal of the sample after sintering [15]. Thus, graphite components are in close con‐ tact with the sample and can become reactive with the oxygen eventually present in the sample itself at temperatures higher than 600°C. Other sources of oxygen are moisture or other gases in the sintering atmosphere. Such a chemical reaction causes the formation of CO and a continuous decrease of the oxygen partial pressure within the furnace, creating a reducing condition in the

Sintering and Reactive Sintering by Spark Plasma Sintering (SPS) http://dx.doi.org/10.5772/intechopen.68871 127

**Figure 2.** Ellingham‐Richardson diagram [16].

are favorable to the imbalance of either one of the following reactions from the right to the

(1)

(2)

(3)

(4)

(5)

(6)

(7)

mole in each oxidizing reaction (from left to

= *f* (*T*) is known as the Ellingham‐Richardson diagram

in the diagram have a greater affinity to

left side:

126 Sintering of Functional Materials

Δ*H*<sup>0</sup> 1 , Δ*H*<sup>0</sup> 2

of O2

temperature *T*:

The graphic representation of Δ*G*<sup>0</sup>

oxidizing agent. Elements with a lower line in Δ*G*<sup>0</sup>

, and Δ*H*<sup>0</sup>

3

are the heat released per O2

right side). By definition, we can calculate the corresponding changes of free enthalpy:

A standard measure for the tendency of a metal (or a chemical element) to oxidize is given by Eqs. (4)–(6) when *P*O2 = 1 ; thus, we obtain Eq. (7), which is the heat released when 1 mole

(**Figure 2**). It allows the direct evaluation of the relative affinity between the elements and the

oxygen. This diagram is useful to understand the thermodynamics of the reactions between the sintering material and the atmosphere, and it can give information about the dissociation

Special attention must be given to the pressing tools' material. Standard pressing tools used in the RSPS process are graphite based, often internally covered with carbon sheets or foils, in order to ease the removal of the sample after sintering [15]. Thus, graphite components are in close con‐ tact with the sample and can become reactive with the oxygen eventually present in the sample itself at temperatures higher than 600°C. Other sources of oxygen are moisture or other gases in the sintering atmosphere. Such a chemical reaction causes the formation of CO and a continuous decrease of the oxygen partial pressure within the furnace, creating a reducing condition in the

temperature, the dissociation pressure, and the effect of reducing agents [14].

gas at 1 atm pressure combines with 1 mole of metal to form the oxide in function of

sintering atmosphere. When an intense gas phase transport is established between the sample and the mold, reduction of oxides or even precipitation of carbon or carbide in the sample may occur [9].

for high‐pressure sintering, because of their high mechanical resistance to compression; the possibility to increase the pressure applied to pistons during consolidation allows to operate the process at even lower temperature, thus limiting the possibility of reaction between carbon and oxide inside the material. Tungsten carbide, steel, and refractory met‐ als, such as molybdenum alloys, copper‐beryllium, and alumina [25], have been also used as conductive sintering tools. Double‐walled tools with inner ceramic die and outer graph‐ ite mantle have been also employed [26]. Some works report the use of layers and foils of alumina or other different metals, such as molybdenum, tungsten, and tantalum, which are introduced inside the graphite die to cover the internal mold walls before introducing the sample [9]. By these less‐costly operating conditions, the sintering material is never in

Sintering and Reactive Sintering by Spark Plasma Sintering (SPS)

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129

In addition to heating, the effect of current pulses is to enhance mass transport during sin‐ tering, more specifically by one of the three mechanisms [27]: (a) increasing the point defect concentration; (b) a reduction in the activation energy for mobility of defects; or (c) electron

Temperature and current are not independent parameters; high heating rates are achieved by increasing pulsed direct current. This is the major difference between the conventional hot‐pressing and the SPS methods; in the RSPS method, both the die (typically graphite) and the sample are heated by Joule mechanism from a current passing through them (if the sample is conducting) [27]. Nonconductive materials are heated by means of heat conduction from the die walls. Pulsed direct current can enhance the reaction kinetics when the reactants are brought to interact in the SPS. This effect is, however, system dependent [28]. A change in the electrical conductivity of the materials in the die as reac‐ tion progresses can give rise to undesired results. Schmidt et al. [29] studied the decom‐

increases the electrical conductivity of the material in the SPS die. This example shows that as the reaction product accumulates, the conductivity of the material in the sintering die changes; if it increases, the reaction is self‐enhanced due to the presence of in‐situ– formed conductive particles inducing the formation of hot spots in the remaining, not yet

In the case of silica‐doped yttria‐stabilized zirconia sintering, the electrical resistivity of grain boundaries is often increased by the presence of impurity phases of siliceous compounds. SPS allowed a significant reduction in these compounds while leaving the grains unaffected [31]. This effect was attributed to the generation of electrical discharges between particles as the SPS electrical pulses are applied. The discharges expel the liquid silica phase to triple points

in RSPS; in order to increase electrical conductivity, graphite

powder. Metallic magnesium, the product of decomposition,

contact with graphite.

**3.1. Electrical current effect**

position behavior of MgH2

was added to the MgH2

fully reacted, mixture.

**3. Effects of current and pressure**

wind modification of the diffusion flux (electromigration).

in grain boundaries, thus reducing their effects on resistivity.

The materials obtained by the RSPS technique can be distinguished into two main categories in relation to their affinity to reductive atmosphere: (a) if oxide reduction is a desirable effect and (b) if reduction is a secondary effect that should be avoided during the ceramic forma‐ tion. Metals and nonoxide ceramics constitute the first class of materials, while oxide ceramics represent the second one.

For metals and nonoxide ceramics, reductive atmospheres, such as inert gas or reducing hydrogen gas mixture, are suggested during sintering, because they are effective on clean‐ ing the oxide naturally formed on the metallic surface of the starting powder during air exposition.

The benefit of oxide reduction in this kind of materials is related to the possibility to enhance the sample densification; in fact, oxide compounds possess a smaller density than the corre‐ sponding metal, and they hinder atomic diffusion during the densification step.

Such a reductive atmosphere has also been suggested for ultra‐high‐temperature ceramics (UHTC) to promote their densification. Sometimes, specific additives are mixed to the starting powder to promote reduction and thus enhance densification. For instance, C or B<sup>4</sup> C has been used as additives in TaC densification [17], while MoSi2 , TaSi2 , and SiC have been considered as additives for oxygen removal during HfB2 sintering [18–20].

Systems requiring a reductive atmosphere during their reactive consolidation are oxide/ metal nanocomposites. In practice, the reductive atmosphere can be specifically used for the in‐situ formation of metal component. As an example, Al2 O3 /Ni granular solids were produced by reacting and sintering a mixture of Al2 O3 and NiO powders inside a carbon die [15].

For functional oxide ceramics, including magnetics, reducing sintering atmosphere may have dramatic consequences on the final properties. It may modify the starting oxide composition. Typically, it generates oxygen vacancies, which in the case of transparent ceramics, such as yttrium‐aluminum‐garnet (YAG), induces light absorption and in‐line transmission decrease [21]. In the case of ferromagnetic *p*‐doped manganite ceramics, the formed oxygen vacancies decrease the average oxidation state of the paramagnetic manganese cations, inducing a net reduction of the Curie temperature of the final solid in comparison to its conventionally made bulk counterpart [22]. In Ni‐Zn ferrite, a well‐known resistive soft magnet, consolidation by SPS involves a Fe3+ into Fe2+reduction inside the ferrite grains compensated by a loss of Ni cations, which precipitate as Ni metal between the grains, increasing the total electrical con‐ ductivity [23].

As a consequence of oxide ceramic changes in reductive atmospheres, solutions to the tools' reactivity at high temperature have been considered. In other words, when the con‐ trol of the atmosphere is not enough to avoid secondary reduction reactions, new tool materials have been employed. Graphite reinforced with carbon fiber dies [24] is suitable for high‐pressure sintering, because of their high mechanical resistance to compression; the possibility to increase the pressure applied to pistons during consolidation allows to operate the process at even lower temperature, thus limiting the possibility of reaction between carbon and oxide inside the material. Tungsten carbide, steel, and refractory met‐ als, such as molybdenum alloys, copper‐beryllium, and alumina [25], have been also used as conductive sintering tools. Double‐walled tools with inner ceramic die and outer graph‐ ite mantle have been also employed [26]. Some works report the use of layers and foils of alumina or other different metals, such as molybdenum, tungsten, and tantalum, which are introduced inside the graphite die to cover the internal mold walls before introducing the sample [9]. By these less‐costly operating conditions, the sintering material is never in contact with graphite.

## **3. Effects of current and pressure**

#### **3.1. Electrical current effect**

sintering atmosphere. When an intense gas phase transport is established between the sample and the mold, reduction of oxides or even precipitation of carbon or carbide in the sample may

The materials obtained by the RSPS technique can be distinguished into two main categories in relation to their affinity to reductive atmosphere: (a) if oxide reduction is a desirable effect and (b) if reduction is a secondary effect that should be avoided during the ceramic forma‐ tion. Metals and nonoxide ceramics constitute the first class of materials, while oxide ceramics

For metals and nonoxide ceramics, reductive atmospheres, such as inert gas or reducing hydrogen gas mixture, are suggested during sintering, because they are effective on clean‐ ing the oxide naturally formed on the metallic surface of the starting powder during air

The benefit of oxide reduction in this kind of materials is related to the possibility to enhance the sample densification; in fact, oxide compounds possess a smaller density than the corre‐

Such a reductive atmosphere has also been suggested for ultra‐high‐temperature ceramics (UHTC) to promote their densification. Sometimes, specific additives are mixed to the starting

Systems requiring a reductive atmosphere during their reactive consolidation are oxide/ metal nanocomposites. In practice, the reductive atmosphere can be specifically used for

For functional oxide ceramics, including magnetics, reducing sintering atmosphere may have dramatic consequences on the final properties. It may modify the starting oxide composition. Typically, it generates oxygen vacancies, which in the case of transparent ceramics, such as yttrium‐aluminum‐garnet (YAG), induces light absorption and in‐line transmission decrease [21]. In the case of ferromagnetic *p*‐doped manganite ceramics, the formed oxygen vacancies decrease the average oxidation state of the paramagnetic manganese cations, inducing a net reduction of the Curie temperature of the final solid in comparison to its conventionally made bulk counterpart [22]. In Ni‐Zn ferrite, a well‐known resistive soft magnet, consolidation by SPS involves a Fe3+ into Fe2+reduction inside the ferrite grains compensated by a loss of Ni cations, which precipitate as Ni metal between the grains, increasing the total electrical con‐

As a consequence of oxide ceramic changes in reductive atmospheres, solutions to the tools' reactivity at high temperature have been considered. In other words, when the con‐ trol of the atmosphere is not enough to avoid secondary reduction reactions, new tool materials have been employed. Graphite reinforced with carbon fiber dies [24] is suitable

, TaSi2

O3

sintering [18–20].

O3

C has been

, and SiC have been considered

and NiO powders inside a carbon

/Ni granular solids were

sponding metal, and they hinder atomic diffusion during the densification step.

powder to promote reduction and thus enhance densification. For instance, C or B<sup>4</sup>

used as additives in TaC densification [17], while MoSi2

produced by reacting and sintering a mixture of Al2

the in‐situ formation of metal component. As an example, Al2

as additives for oxygen removal during HfB2

occur [9].

128 Sintering of Functional Materials

exposition.

die [15].

ductivity [23].

represent the second one.

In addition to heating, the effect of current pulses is to enhance mass transport during sin‐ tering, more specifically by one of the three mechanisms [27]: (a) increasing the point defect concentration; (b) a reduction in the activation energy for mobility of defects; or (c) electron wind modification of the diffusion flux (electromigration).

Temperature and current are not independent parameters; high heating rates are achieved by increasing pulsed direct current. This is the major difference between the conventional hot‐pressing and the SPS methods; in the RSPS method, both the die (typically graphite) and the sample are heated by Joule mechanism from a current passing through them (if the sample is conducting) [27]. Nonconductive materials are heated by means of heat conduction from the die walls. Pulsed direct current can enhance the reaction kinetics when the reactants are brought to interact in the SPS. This effect is, however, system dependent [28]. A change in the electrical conductivity of the materials in the die as reac‐ tion progresses can give rise to undesired results. Schmidt et al. [29] studied the decom‐ position behavior of MgH2 in RSPS; in order to increase electrical conductivity, graphite was added to the MgH2 powder. Metallic magnesium, the product of decomposition, increases the electrical conductivity of the material in the SPS die. This example shows that as the reaction product accumulates, the conductivity of the material in the sintering die changes; if it increases, the reaction is self‐enhanced due to the presence of in‐situ– formed conductive particles inducing the formation of hot spots in the remaining, not yet fully reacted, mixture.

In the case of silica‐doped yttria‐stabilized zirconia sintering, the electrical resistivity of grain boundaries is often increased by the presence of impurity phases of siliceous compounds. SPS allowed a significant reduction in these compounds while leaving the grains unaffected [31]. This effect was attributed to the generation of electrical discharges between particles as the SPS electrical pulses are applied. The discharges expel the liquid silica phase to triple points in grain boundaries, thus reducing their effects on resistivity.

#### **3.2. Pressure effect**

Mechanically, the pressure has a direct effect on particle rearrangement and the destruction of agglomerates, particularly in the case of nanometric powders. However, the significance of the pressure on sintering depends on the particle size. When the particle size is small, the relative contribution of the pressure is small but becomes significant as the particle size increases [27]. In a study on the sintering of nanometric pure zirconia, Skandan et al. [30] found that the pressure had no effect on the relative density of fine‐grained powder (6 nm) up to a pressure of about 35 MPa; in contrast, the density increased sharply when higher pressure Nygren [32], the grain growth is deeply related with the size of starting powder. When nano‐ metric size precursors are employed, most of the driving force to reduce specific area is des‐

Sintering and Reactive Sintering by Spark Plasma Sintering (SPS)

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In the case of micrometric precursors, pressure can enhance densification by four mechanisms [33]: particle rearrangement, localized deformation, bulk deformation, and neck growth. For large particle size of precursors (45–90 μm), density is enhanced through particle rearrange‐ ment and localized deformation. In contrast, neck growth increases for smaller size of precur‐

Grains can keep the memory of the synthesis process by which they were made, and their char‐

RSPS, varying the initial powder size and the nature of precursors could enhance consolidation, particularly for sol‐gel and co‐precipitation synthesis [34]. Co‐gelified alumina‐titania powders and mechanical mixtures of alumina and titania (both obtained by sol‐gel method) and alu‐ mina‐titania powders (recovered by co‐precipitation method) were treated thermally before the RSPS process. The grain sizes for the powders synthesized by sol‐gel and co‐precipitation method were 10 and 50 nm, respectively. Experimental conditions were the same for all the three samples during RSPS process. The final report showed a smaller increase in the grain size (0.5 μm) for the consolidated powders synthesized by sol‐gel as compared with the co‐pre‐ cipitated powders (8 μm). However, the final density for the co‐gelified alumina‐titania and

was obtained in co‐gelified alumina‐synthesized sample.

dated in similar conditions [35]. In the RSPS process, monodispersed initial powders around 5 nm, and clusters of ∼50 nm (made also of ∼5 nm particles) were rapidly heated to 600°C for 6 min before rapid cooling. Unexpectedly, the final grain size resulted larger for the mono‐ dispersed precursor than for the clustered case. This difference was interpreted on the basis that grain growth is an essentially surface process, and in the monodispersed case, particles

Preliminary works exploring the possibility of preparing nanostructured manganite ceramics by the RSPS process have evidenced the role of the precursor's nature. Starting from a mixture of raw bulk oxides required a higher reacting and sintering temperature, while starting from a mix‐ ture of their hydroxide counterparts allowed a decrease in this operating parameter. Typically,

the same pressure for almost the same sintering time (**Figure 4b**) [42]. The final density of both ceramics exceeded 90% of the theoretical value, and their average grain size was in the submi‐ crometer range, the finest grains being obtained at the lowest sintering temperature of course.

of 50 MPs for a couple of minutes (**Figure 4a**) [12]. In contrast, ceramic LaMnO3

O3

ceramics were prepared by RSPS starting from the La(OH)3

commercial powder mixture, working at 800°C under a uniaxial pressure

commercial powder mixture, working at 1000°C under

O4

TiO5 by 131

). Finally, the

, Na(OH),

was obtained

starting powders were consoli‐

acteristics are retained even in the sintered particles [34]. During the processing of Al2

alumina titania initial powders was very close to the theoretical density (3.7 g/cm3

tined to the densification process.

full phase of Al2

**4.2. Nature of precursors**

MnO3

La0,85Na0,15‐xKx

K(OH), and MnO2

by RSPS starting from La2

TiO5

sors (∼25 μm), while bulk deformation has no influence [33].

Different size and aggregation states of polyol‐made CoFe<sup>2</sup>

offered a larger free surface than in the clustered case.

O3

and Mn2

**Figure 3.** Relationship between hold temperature and the applied pressure required to obtain samples with a relative density of 95% in the case of nanometric fully stabilized zirconia (8% Y<sup>2</sup> O3 ). Hold time: 5 min. The grain size of the materials is also shown [30].

was used. For larger particle size powder (12 nm), the same behavior was observed except that the transition occurred at about 10 MPa. Another result of the application of pressure is a decrease in the sintering temperature. For the case of SPS densification of a nanometric cubic zirconia, Anselmi‐Tamburini et al. [27] showed that the combination of fast heating rate and high pressure produces a marked reduction in the sintering temperature. **Figure 3** shows the effect of pressure on the sintering temperature required to obtain a 95% relative density (with 5‐min hold time). The figure also shows the grain size obtained under these conditions. The temperature required to achieve 95% of density decreases linearly with the logarithm of the applied pressure. The grain size varied from about 200 to 15 nm.

#### **4. Precursors**

#### **4.1. Size of precursors**

The size of precursors plays an important role in the final consolidates. In this manner, den‐ sification hinges on the characteristics of initial powders inside the die. As mentioned by Nygren [32], the grain growth is deeply related with the size of starting powder. When nano‐ metric size precursors are employed, most of the driving force to reduce specific area is des‐ tined to the densification process.

In the case of micrometric precursors, pressure can enhance densification by four mechanisms [33]: particle rearrangement, localized deformation, bulk deformation, and neck growth. For large particle size of precursors (45–90 μm), density is enhanced through particle rearrange‐ ment and localized deformation. In contrast, neck growth increases for smaller size of precur‐ sors (∼25 μm), while bulk deformation has no influence [33].

Grains can keep the memory of the synthesis process by which they were made, and their char‐ acteristics are retained even in the sintered particles [34]. During the processing of Al2 TiO5 by RSPS, varying the initial powder size and the nature of precursors could enhance consolidation, particularly for sol‐gel and co‐precipitation synthesis [34]. Co‐gelified alumina‐titania powders and mechanical mixtures of alumina and titania (both obtained by sol‐gel method) and alu‐ mina‐titania powders (recovered by co‐precipitation method) were treated thermally before the RSPS process. The grain sizes for the powders synthesized by sol‐gel and co‐precipitation method were 10 and 50 nm, respectively. Experimental conditions were the same for all the three samples during RSPS process. The final report showed a smaller increase in the grain size (0.5 μm) for the consolidated powders synthesized by sol‐gel as compared with the co‐pre‐ cipitated powders (8 μm). However, the final density for the co‐gelified alumina‐titania and alumina titania initial powders was very close to the theoretical density (3.7 g/cm3 ). Finally, the full phase of Al2 TiO5 was obtained in co‐gelified alumina‐synthesized sample.

Different size and aggregation states of polyol‐made CoFe<sup>2</sup> O4 starting powders were consoli‐ dated in similar conditions [35]. In the RSPS process, monodispersed initial powders around 5 nm, and clusters of ∼50 nm (made also of ∼5 nm particles) were rapidly heated to 600°C for 6 min before rapid cooling. Unexpectedly, the final grain size resulted larger for the mono‐ dispersed precursor than for the clustered case. This difference was interpreted on the basis that grain growth is an essentially surface process, and in the monodispersed case, particles offered a larger free surface than in the clustered case.

#### **4.2. Nature of precursors**

**3.2. Pressure effect**

130 Sintering of Functional Materials

**4. Precursors**

**4.1. Size of precursors**

materials is also shown [30].

Mechanically, the pressure has a direct effect on particle rearrangement and the destruction of agglomerates, particularly in the case of nanometric powders. However, the significance of the pressure on sintering depends on the particle size. When the particle size is small, the relative contribution of the pressure is small but becomes significant as the particle size increases [27]. In a study on the sintering of nanometric pure zirconia, Skandan et al. [30] found that the pressure had no effect on the relative density of fine‐grained powder (6 nm) up to a pressure of about 35 MPa; in contrast, the density increased sharply when higher pressure

was used. For larger particle size powder (12 nm), the same behavior was observed except that the transition occurred at about 10 MPa. Another result of the application of pressure is a decrease in the sintering temperature. For the case of SPS densification of a nanometric cubic zirconia, Anselmi‐Tamburini et al. [27] showed that the combination of fast heating rate and high pressure produces a marked reduction in the sintering temperature. **Figure 3** shows the effect of pressure on the sintering temperature required to obtain a 95% relative density (with 5‐min hold time). The figure also shows the grain size obtained under these conditions. The temperature required to achieve 95% of density decreases linearly with the logarithm of the

**Figure 3.** Relationship between hold temperature and the applied pressure required to obtain samples with a relative

O3

). Hold time: 5 min. The grain size of the

The size of precursors plays an important role in the final consolidates. In this manner, den‐ sification hinges on the characteristics of initial powders inside the die. As mentioned by

applied pressure. The grain size varied from about 200 to 15 nm.

density of 95% in the case of nanometric fully stabilized zirconia (8% Y<sup>2</sup>

Preliminary works exploring the possibility of preparing nanostructured manganite ceramics by the RSPS process have evidenced the role of the precursor's nature. Starting from a mixture of raw bulk oxides required a higher reacting and sintering temperature, while starting from a mix‐ ture of their hydroxide counterparts allowed a decrease in this operating parameter. Typically, La0,85Na0,15‐xKx MnO3 ceramics were prepared by RSPS starting from the La(OH)3 , Na(OH), K(OH), and MnO2 commercial powder mixture, working at 800°C under a uniaxial pressure of 50 MPs for a couple of minutes (**Figure 4a**) [12]. In contrast, ceramic LaMnO3 was obtained by RSPS starting from La2 O3 and Mn2 O3 commercial powder mixture, working at 1000°C under the same pressure for almost the same sintering time (**Figure 4b**) [42]. The final density of both ceramics exceeded 90% of the theoretical value, and their average grain size was in the submi‐ crometer range, the finest grains being obtained at the lowest sintering temperature of course.

atmosphere. The main conclusion was that the final particle size increased with increasing milling speed. Also, in milling times shorter than 2 h at 800 rpm, the lattice parameter varia‐

Amorphous or poorly crystallized intermediate solid phases obtained by soft chemistry, combining precipitation in a liquid solution and moderate annealing, and containing all the desired elements were also used to form highly dense and fine‐grained oxide ceramics. This

ics. The precipitated solids were first annealed at 600 and 400°C, respectively, to remove the

from La3+to Lu3+. One of the most studied phases is the yttrium iron garnet (YIG), which is a remarkable ferrimagnetic material with many applications in microwave [44], magne‐ tooptical [45], and spintronic devices [46], most of them based on the fact that YIG has the smallest linewidth for ferromagnetic resonance (FMR) [44]. Its ferrimagnetism results from

**Figure 5.** Hysteresis loops of ball‐milled mixtures of iron and yttrium oxide for 5 h and annealed for 3 h at different

of 700 and 750°C, respectively, to obtain highly dense and fine‐grained ceramics.

**5. RSPS‐made magnetic ceramics: Synthesis and properties**

Magnetic garnets possess the crystal structure of mineral Mn3

(RE) and Fe3+ cations instead, leading to the general formula RE3

O and CO2

manganite [43] and Y<sup>3</sup>

Fe5

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Al2 Si3

Fe5

, and then SPS treated at a temperature

O12 garnet [49] ceram‐

133

O12, with rare‐earth

O12, RE is in the series

tion was insignificant.

**5.1. Soft magnets: Garnets**

temperatures (adapted from Ref. [48]).

is for instance the case of La0.65Ca0.20Na0.15MnO3

main noninorganic species in the form of H2

**Figure 4.** Shrinkage curve (temperature and piston displacement as a function of time) recorded during the RSPS process of La0.85Na0.15MnO3 (a) and LaMnO3 (b) ceramics and representative SEM micrographs of each ceramic (c and d, respectively) (adapted from Ref. [42]).

Starting from a mixture of raw oxides assisted by ball milling has been extensively used for the mechanical activation of the reactive powders before SPS treatment [36–38]. In order to achieve high‐density ceramics with low SPS temperatures, a small particle size and high reactivity must be taken into consideration [39]. As an example, Ni0.5Zn0.5Fe2 O4 ferrite was prepared by Song et al. [40]. Stoichiometric quantities of NiO, ZnO, and Fe2 O3 were milled in a high‐energy planetary ball mill. Different parameters of the grinding time were varied, for example 10, 20, and 40 h, at a speed of 400 rpm. As expected, they found the higher the grinding time, the smaller the average size for the starting powders (<100 nm). For the sin‐ tering process, they selected the powders with 40 h milling time and particle size below 100 nm. Different temperatures were chosen (850, 875, 900, and 925°C) for 5 min, a pressure of 48 MPa and 5‐min vacuum were applied. The best densification result was obtained at 925°C. A density of 5.23 g/cm3 , corresponding to 99% of the theoretical value, was reached. No second‐ ary phases were detected in structural characterization. Zehani et al. [41] studied NiZnCu ferrite at several grinding times and speeds. The stoichiometric proportions of precursor oxides (NiO, Fe2 O3 , CuO, ZnO) were ground in a planetary mill. RSPS was then performed at different temperatures and holding times, using a graphite die and working under argon atmosphere. The main conclusion was that the final particle size increased with increasing milling speed. Also, in milling times shorter than 2 h at 800 rpm, the lattice parameter varia‐ tion was insignificant.

Amorphous or poorly crystallized intermediate solid phases obtained by soft chemistry, combining precipitation in a liquid solution and moderate annealing, and containing all the desired elements were also used to form highly dense and fine‐grained oxide ceramics. This is for instance the case of La0.65Ca0.20Na0.15MnO3 manganite [43] and Y<sup>3</sup> Fe5 O12 garnet [49] ceram‐ ics. The precipitated solids were first annealed at 600 and 400°C, respectively, to remove the main noninorganic species in the form of H2 O and CO2 , and then SPS treated at a temperature of 700 and 750°C, respectively, to obtain highly dense and fine‐grained ceramics.

## **5. RSPS‐made magnetic ceramics: Synthesis and properties**

#### **5.1. Soft magnets: Garnets**

Starting from a mixture of raw oxides assisted by ball milling has been extensively used for the mechanical activation of the reactive powders before SPS treatment [36–38]. In order to achieve high‐density ceramics with low SPS temperatures, a small particle size and high

**Figure 4.** Shrinkage curve (temperature and piston displacement as a function of time) recorded during the RSPS

**Déplacement de pistons (mm)**

**Déplacement de pistons (mm)**

**c)**

**d)**

(b) ceramics and representative SEM micrographs of each ceramic (c and d,

in a high‐energy planetary ball mill. Different parameters of the grinding time were varied, for example 10, 20, and 40 h, at a speed of 400 rpm. As expected, they found the higher the grinding time, the smaller the average size for the starting powders (<100 nm). For the sin‐ tering process, they selected the powders with 40 h milling time and particle size below 100 nm. Different temperatures were chosen (850, 875, 900, and 925°C) for 5 min, a pressure of 48 MPa and 5‐min vacuum were applied. The best densification result was obtained at 925°C. A

ary phases were detected in structural characterization. Zehani et al. [41] studied NiZnCu ferrite at several grinding times and speeds. The stoichiometric proportions of precursor

at different temperatures and holding times, using a graphite die and working under argon

, corresponding to 99% of the theoretical value, was reached. No second‐

, CuO, ZnO) were ground in a planetary mill. RSPS was then performed

O4

**200 nm**

O3

ferrite was

were milled

reactivity must be taken into consideration [39]. As an example, Ni0.5Zn0.5Fe2

**0 5 10 15 20 25 30**

**Temps (minute)**

**0 5 10 15 20 25 30**

**Temps (minute)**

(a) and LaMnO3

**Température (°C) Déplacement de pistons (mm)**

**Température (°C) Déplacement de pistons (mm)**

prepared by Song et al. [40]. Stoichiometric quantities of NiO, ZnO, and Fe2

density of 5.23 g/cm3

**0**

**0**

process of La0.85Na0.15MnO3

respectively) (adapted from Ref. [42]).

**200**

**400**

**600**

**Température (°C)**

**800**

**1000**

b)

**200**

**400**

**Température (°C)**

**600**

**800**

132 Sintering of Functional Materials

a)

O3

oxides (NiO, Fe2

Magnetic garnets possess the crystal structure of mineral Mn3 Al2 Si3 O12, with rare‐earth (RE) and Fe3+ cations instead, leading to the general formula RE3 Fe5 O12, RE is in the series from La3+to Lu3+. One of the most studied phases is the yttrium iron garnet (YIG), which is a remarkable ferrimagnetic material with many applications in microwave [44], magne‐ tooptical [45], and spintronic devices [46], most of them based on the fact that YIG has the smallest linewidth for ferromagnetic resonance (FMR) [44]. Its ferrimagnetism results from

**Figure 5.** Hysteresis loops of ball‐milled mixtures of iron and yttrium oxide for 5 h and annealed for 3 h at different temperatures (adapted from Ref. [48]).

superexchange interactions [55] between octahedral and tetrahedral Fe3+ cations, which are antiparallel. As a bulk, YIG is commonly prepared by the classic solid‐state reaction tech‐ nique which involves temperatures as high as 1350°C, for a few hours [47].

Interestingly, the resulting dense and submicrometer‐grain‐sized ceramic exhibited the same magnetic properties as the conventionally made bulk counterpart: a saturation magnetization of 28 emu/g and a coercive field close to zero at room temperature. Clearly, the reduction of the grain size from the micrometer size range to the submicrometer one does not introduce major magnetic changes, the surface‐to‐volume atomic fraction remaining negligible in both

lent alkaline‐earth cation) phases with the perovskite structure have been extensively studied over the last 15 years in view of their remarkable physical properties, which can be used for a wide variety of applications, particularly for giant magnetoresistance devices and magne‐ tocalorics [50–53]. The correlation between magnetic and transport properties are interpreted on the basis of double exchange (DE) mechanism [54], the superexchange (SE) interactions [55], the electron‐phonon coupling due to the Jahn‐Teller effect of Mn3+ ions, and the mag‐

Mn3+ ions is an insulating antiferromagnet [56], while doped ones contain Mn3+ and Mn4+ ions and may be ferromagnetic conductors. The magnetic properties of the former are driven by SE interactions, while those of the latter are mainly due to DE interactions. Consequently, the

0 50 100 150 200 250 300

**T (K)**

(Ln: trivalent rare‐earth ion; X: monovalent alkaline, or diva‐

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135

compound, containing only

ceramic produced by (a) conventional

cases to induce significant magnetic changes.

MnO3

netic‐phase separation [51]. The undoped stoichiometric LaMnO3

**5.2. Magnetocalorics: Manganites**

0

5

10

15

**M (emu/g)**

20

(b)

500 mT

**Figure 7.** Temperature dependence of magnetization at 50 mT of La0.70Ca0.30MnO3

solid‐state route at 1300°C and (b) by RSPSP at 800°C (adapted from Ref. [59]).

(a)

25

30

35

*P*‐doped manganite Ln1−xX<sup>x</sup>

Nanostructured YIG (or other garnet) is typically prepared by combining soft chemistry, or ball milling, and annealing to complete the reaction before SPS sintering. This garnet crystal structure possesses a relatively large unit cell (160 atoms), making it difficult to achieve its synthesis at low temperature. A typical procedure can be high‐energy ball milling of Fe2 O3 + Y<sup>2</sup> O3 oxide reagents, followed by thermal annealing before SPS treatment [48]. The mag‐ netization of the recovered annealed powders increases with the increase in the annealing temperature, and it approaches its bulk value, namely 28 emu/g, only by samples annealed at *T* ≥ 900°C (**Figure 5**).

RSPS offers an excellent alternative to produce consolidated nanostructured YIG at low temperatures and very short sintering times. A convenient YIG precursor can be prepared by hydrolysis in a polyol method, followed by processing by RSPS to obtain a nanostruc‐ tured garnet phase [49], with the general magnetic properties of bulk YIG. The interme‐ diate solid phase is amorphous (**Figure 6**) with the required Y/Fe stoichiometric ratio. Its preannealing at a temperature of 400°C allows its decomposition and the removal of main organic contents, but it is unsuitable to form the desired garnet phase. A reaction/sintering RSPS treatment at 750°C for 15 min produced a nanostructured solid with high density and nanosize grains. XRD‐resolved patterns, **Figure 6**, showed that an amorphous phase leads first to the orthoferrite YFeO<sup>3</sup> phase (600–650°C) and then the transformation to the garnet Y3 Fe5 O12phasefrom 750°C.

**Figure 6.** Temperature‐resolved X‐ray diffraction patterns of the polyol‐synthesized YIG precursor. At about 600°C, yttrium orthoferrite (YFeO<sup>3</sup> ) is formed, which then transforms into YIG at higher temperatures [49].

Interestingly, the resulting dense and submicrometer‐grain‐sized ceramic exhibited the same magnetic properties as the conventionally made bulk counterpart: a saturation magnetization of 28 emu/g and a coercive field close to zero at room temperature. Clearly, the reduction of the grain size from the micrometer size range to the submicrometer one does not introduce major magnetic changes, the surface‐to‐volume atomic fraction remaining negligible in both cases to induce significant magnetic changes.

#### **5.2. Magnetocalorics: Manganites**

superexchange interactions [55] between octahedral and tetrahedral Fe3+ cations, which are antiparallel. As a bulk, YIG is commonly prepared by the classic solid‐state reaction tech‐

Nanostructured YIG (or other garnet) is typically prepared by combining soft chemistry, or ball milling, and annealing to complete the reaction before SPS sintering. This garnet crystal structure possesses a relatively large unit cell (160 atoms), making it difficult to achieve its synthesis at low temperature. A typical procedure can be high‐energy ball milling of Fe2

RSPS offers an excellent alternative to produce consolidated nanostructured YIG at low temperatures and very short sintering times. A convenient YIG precursor can be prepared by hydrolysis in a polyol method, followed by processing by RSPS to obtain a nanostruc‐ tured garnet phase [49], with the general magnetic properties of bulk YIG. The interme‐ diate solid phase is amorphous (**Figure 6**) with the required Y/Fe stoichiometric ratio. Its preannealing at a temperature of 400°C allows its decomposition and the removal of main organic contents, but it is unsuitable to form the desired garnet phase. A reaction/sintering RSPS treatment at 750°C for 15 min produced a nanostructured solid with high density and nanosize grains. XRD‐resolved patterns, **Figure 6**, showed that an amorphous phase leads

**Figure 6.** Temperature‐resolved X‐ray diffraction patterns of the polyol‐synthesized YIG precursor. At about 600°C,

) is formed, which then transforms into YIG at higher temperatures [49].

oxide reagents, followed by thermal annealing before SPS treatment [48]. The mag‐ netization of the recovered annealed powders increases with the increase in the annealing temperature, and it approaches its bulk value, namely 28 emu/g, only by samples annealed at

phase (600–650°C) and then the transformation to the garnet

O3

nique which involves temperatures as high as 1350°C, for a few hours [47].

+ Y<sup>2</sup> O3

Y3 Fe5

*T* ≥ 900°C (**Figure 5**).

134 Sintering of Functional Materials

first to the orthoferrite YFeO<sup>3</sup>

yttrium orthoferrite (YFeO<sup>3</sup>

O12phasefrom 750°C.

*P*‐doped manganite Ln1−xX<sup>x</sup> MnO3 (Ln: trivalent rare‐earth ion; X: monovalent alkaline, or diva‐ lent alkaline‐earth cation) phases with the perovskite structure have been extensively studied over the last 15 years in view of their remarkable physical properties, which can be used for a wide variety of applications, particularly for giant magnetoresistance devices and magne‐ tocalorics [50–53]. The correlation between magnetic and transport properties are interpreted on the basis of double exchange (DE) mechanism [54], the superexchange (SE) interactions [55], the electron‐phonon coupling due to the Jahn‐Teller effect of Mn3+ ions, and the mag‐ netic‐phase separation [51]. The undoped stoichiometric LaMnO3 compound, containing only Mn3+ ions is an insulating antiferromagnet [56], while doped ones contain Mn3+ and Mn4+ ions and may be ferromagnetic conductors. The magnetic properties of the former are driven by SE interactions, while those of the latter are mainly due to DE interactions. Consequently, the

**Figure 7.** Temperature dependence of magnetization at 50 mT of La0.70Ca0.30MnO3 ceramic produced by (a) conventional solid‐state route at 1300°C and (b) by RSPSP at 800°C (adapted from Ref. [59]).

atomic Mn4+/Mn3+ ratio is a key parameter in the achievement of the magnetoelectrical proper‐ ties of these oxides.

**5.3. Hard magnets: Hexaferrites**

smaller coercive field (**Figure 9**).

Hexaferrites have become extremely important materials since they have a large variety of applications due to their high magnetocrystalline anisotropy in relation with their Δ*M* hex‐ agonal structure. Their magnetic properties are mainly driven by SE interactions as most of insulating oxides. Technologically speaking, hexaferrites are mainly used in the form of bulk solids as permanent magnets in magnetic recording and magnetic data storage devices, and

Sintering and Reactive Sintering by Spark Plasma Sintering (SPS)

http://dx.doi.org/10.5772/intechopen.68871

137

As nanopowders, hexaferrites can be produced by different methods such as sol‐gel [61], hydrothermal [62], aerosol pyrolysis [63], or mechanochemical synthesis [64]. In most of the cases, a subsequent annealing is required to provide enough energy to complete phase forma‐

To the best of our knowledge, the RSPS process has been scarcely used to produce consolidated nanostructured hexaferrites. Bolarín‐Miró et al. [65] carried out a comparative study between M‐type strontium hexaferrite prepared from strontium and iron single oxides mechanically activated by high‐energy ball milling for 5 h followed by RSPS, and the same milled powder mixture sintered by conventional route. They showed that, in comparison with conventional heat treatment, RSPS process allows the formation of strontium hexaferrite single phase at lower temperatures with a higher magnetization. In contrast, the resulting ceramics exhibited

Stingaciu et al. [66] reported the preparation of strontium hexaferrite by SPS starting also from ball‐milling‐activated commercially available SrFe12O19 powder. They observed that the hexaferrite phase is maintained during the milling process (8 to 42 h), while it is not during the SPS treatment. Due to the reductive operating conditions, a nonnegligible amount of mag‐

**Figure 9.** *J(H)* Hysteresis loops of precursor's mixtures milled for 5 h and (a) pressed at 800 MPa and annealed at 800°C

(circle) and (b) SPS‐treated at 700°C under 80 MPa (square) (adapted from Ref. [65]).

O4

nanocomposite. Moreover, they

) and an increase

more recently, as systems operating at microwave/GHz frequencies [60].

tion and an SPS sintering is needed to achieve their consolidation.

netite is formed leading to the production of a SrFe12O19‐Fe3

evidenced a pronounced decrease in the room‐temperature coercive field (*H*<sup>c</sup>

To date, RSPS was successfully used to produce various manganite solid solutions, start‐ ing from raw oxides or hydroxides annealed at 800–1000°C under an uniaxial pressure of 50 MPa and under vacuum: LaMnO3 [42], La0,85Na0,15MnO3 [57], La0,85Na0,15‐xKx MnO3 [12], La0.67Ca0.33MnO3 [58], and La0.7àCa0.30‐xBax MnO3 [59] among others. Interestingly, all the pro‐ duced ceramics exhibited high densities over 90% of the theoretical values and submicrom‐ eter grain size, and a systematically broadened paramagnetic‐to‐ferromagnetic transition as a function of temperature, with a decreased Curie temperature (*T*C) value (**Figure 7**). These changes are due to an evolution of the chemical composition concerning the synthesis condi‐ tions. The conventionally made ceramics are assumed to be chemically homogeneous with a Mn4+/Mn3+atomic ratio fixed by the doping rate, while the RSPS‐made ones may suffer from heterogeneities related to their very rapid reacting/sintering kinetics. These heterogeneities can be associated to a Mn4+ concentration variation between the ceramic core and its surface in contact with graphite during RSPS experiments, with a total Mn4+/Mn3+ atomic ratio smaller than its theoretical value. This discrepancy was confirmed by K‐Mn edge X‐ray absorption spectroscopy (XANES) and iodometry chemical analysis and related to the reductive SPS pro‐ cessing conditions [58].

To evaluate the magnetocaloric properties of manganites, the variation of the magnetic entropy upon a given magnetic field change, Δ*M*, is usually inferred from the first mag‐ netization curves and plotted as a function of the temperature around *T*C value. Δ*M* (*T*) of RSPS‐processed manganites is systematically much more broadened than that of their con‐ ventionally made ceramics (**Figure 8**) for the reasons given above. The commercial applica‐ tions require a magnetocaloric effect extending on a broad temperature range. RSPS ceramics may just offer such an opportunity.

**Figure 8.** Magnetic entropy variation ΔSM as a function of the temperature for a magnetic field change of 1, 2, 3, 4, and 5T La0.70Ca0.30MnO3 ceramic produced by conventional solid‐state route at 1300°C (a) and by RSPS at 800°C (adapted from Ref. [59]).

#### **5.3. Hard magnets: Hexaferrites**

atomic Mn4+/Mn3+ ratio is a key parameter in the achievement of the magnetoelectrical proper‐

To date, RSPS was successfully used to produce various manganite solid solutions, start‐ ing from raw oxides or hydroxides annealed at 800–1000°C under an uniaxial pressure of

duced ceramics exhibited high densities over 90% of the theoretical values and submicrom‐ eter grain size, and a systematically broadened paramagnetic‐to‐ferromagnetic transition as a function of temperature, with a decreased Curie temperature (*T*C) value (**Figure 7**). These changes are due to an evolution of the chemical composition concerning the synthesis condi‐ tions. The conventionally made ceramics are assumed to be chemically homogeneous with a Mn4+/Mn3+atomic ratio fixed by the doping rate, while the RSPS‐made ones may suffer from heterogeneities related to their very rapid reacting/sintering kinetics. These heterogeneities can be associated to a Mn4+ concentration variation between the ceramic core and its surface in contact with graphite during RSPS experiments, with a total Mn4+/Mn3+ atomic ratio smaller than its theoretical value. This discrepancy was confirmed by K‐Mn edge X‐ray absorption spectroscopy (XANES) and iodometry chemical analysis and related to the reductive SPS pro‐

To evaluate the magnetocaloric properties of manganites, the variation of the magnetic entropy upon a given magnetic field change, Δ*M*, is usually inferred from the first mag‐ netization curves and plotted as a function of the temperature around *T*C value. Δ*M* (*T*) of RSPS‐processed manganites is systematically much more broadened than that of their con‐ ventionally made ceramics (**Figure 8**) for the reasons given above. The commercial applica‐ tions require a magnetocaloric effect extending on a broad temperature range. RSPS ceramics

> 1T 2T 3T 4T 5T

> > **∆S**

**Figure 8.** Magnetic entropy variation ΔSM as a function of the temperature for a magnetic field change of 1, 2, 3, 4, and

**- M (J/**

**k**

**g**

**K)**

0.0

ceramic produced by conventional solid‐state route at 1300°C (a) and by RSPS at 800°C (adapted

0.5

1.0

1.5

2.0 **b)**

MnO3

[42], La0,85Na0,15MnO3

[57], La0,85Na0,15‐xKx

100 120 140 160 180 200 220

**T (K)**

[59] among others. Interestingly, all the pro‐

MnO3

[12],

1T 2T 3T 4T 5T

ties of these oxides.

136 Sintering of Functional Materials

La0.67Ca0.33MnO3

cessing conditions [58].

5T La0.70Ca0.30MnO3

from Ref. [59]).

**- ∆S**

**M (J/**

**kgK)**

may just offer such an opportunity.

200 220 240 260 280 300 320 340

**T (K)**

50 MPa and under vacuum: LaMnO3

[58], and La0.7àCa0.30‐xBax

Hexaferrites have become extremely important materials since they have a large variety of applications due to their high magnetocrystalline anisotropy in relation with their Δ*M* hex‐ agonal structure. Their magnetic properties are mainly driven by SE interactions as most of insulating oxides. Technologically speaking, hexaferrites are mainly used in the form of bulk solids as permanent magnets in magnetic recording and magnetic data storage devices, and more recently, as systems operating at microwave/GHz frequencies [60].

As nanopowders, hexaferrites can be produced by different methods such as sol‐gel [61], hydrothermal [62], aerosol pyrolysis [63], or mechanochemical synthesis [64]. In most of the cases, a subsequent annealing is required to provide enough energy to complete phase forma‐ tion and an SPS sintering is needed to achieve their consolidation.

To the best of our knowledge, the RSPS process has been scarcely used to produce consolidated nanostructured hexaferrites. Bolarín‐Miró et al. [65] carried out a comparative study between M‐type strontium hexaferrite prepared from strontium and iron single oxides mechanically activated by high‐energy ball milling for 5 h followed by RSPS, and the same milled powder mixture sintered by conventional route. They showed that, in comparison with conventional heat treatment, RSPS process allows the formation of strontium hexaferrite single phase at lower temperatures with a higher magnetization. In contrast, the resulting ceramics exhibited smaller coercive field (**Figure 9**).

Stingaciu et al. [66] reported the preparation of strontium hexaferrite by SPS starting also from ball‐milling‐activated commercially available SrFe12O19 powder. They observed that the hexaferrite phase is maintained during the milling process (8 to 42 h), while it is not during the SPS treatment. Due to the reductive operating conditions, a nonnegligible amount of mag‐ netite is formed leading to the production of a SrFe12O19‐Fe3 O4 nanocomposite. Moreover, they evidenced a pronounced decrease in the room‐temperature coercive field (*H*<sup>c</sup> ) and an increase

**Figure 9.** *J(H)* Hysteresis loops of precursor's mixtures milled for 5 h and (a) pressed at 800 MPa and annealed at 800°C (circle) and (b) SPS‐treated at 700°C under 80 MPa (square) (adapted from Ref. [65]).

in the magnetization (*M*s) at maximum applied field of 1 T, for powders milled for a longer time and consolidated (**Figure 10**). They concluded that the magnetic properties of the studied nanocomposites are largely conditioned by the extrinsic properties of the secondary phase,

at maximum applied field of 1 T (and hence the energy product *BH*) was measured at room temperature, on the nanocomposite. The hysteresis loops also appeared closer to a rectangu‐ lar form, which is the best shape for applications, as this leads to a well‐defined coercive field

**b)**

O3

*H* (kOe)


**a) c)**

Currently, spark plasma sintering appears as the only method capable to consolidate nanopowders into high‐density nanostructured solids; in this chapter, we have briefly reviewed its application to carry out also the solid‐state reaction needed to achieve a par‐ ticular phase, starting from precursors synthesized by diverse methods. Many challenges remain, especially in the cases of reaction by precursor decomposition, when reaction and sintering temperatures are significantly different. RSPS is still a very young tech‐ nique, with many potential capabilities, which will certainly be developed in the near

, Gabriela Vázquez‐Victorio1,2, Souad Ammar1

**200 nm**

139

nanostructures produced by combining polyol

Sintering and Reactive Sintering by Spark Plasma Sintering (SPS)

http://dx.doi.org/10.5772/intechopen.68871

, Nancy Flores‐Martínez<sup>1</sup>

1 ITODYS, Université Paris Diderot, Sorbonne Paris Cité, Paris Cedex, France

2 Institute for Materials Research, National Autonomous University of Mexico, Mexico City,

\*Address all correspondence to: monjaras@unam.mx

and remanent magnetization.

**200 nm**

**Figure 11.** SEM micrographs of (a) BaFe12O19 and (c) BaFe12O19‐Fe2


process to SPS (800°C, 100 MPa, 5 min), and their room temperature hysteresis loops (b) [67].

0.0

*M* (emu/g)

0.5

BFO

BFO + oxide

1.0

**6. Conclusions**

future.

México

**Author details**

Giulia Franceschin<sup>1</sup>

and Raul Valenzuela1,2\*

**Figure 10.** Evolution of the coercive field *H*c (a), the remanent magnetization *M*r (a), the magnetization at maximum applied field 1 T, *M*s, and the maximum magnetic energy product (*BH*) max of SPS‐consolidated SrFe12O19 powder after different milling times. These data obtained from the hysteresis loops recorded at room temperature, applying the external magnetic field perpendicular to the uniaxial SPS pressing direction [66].

Fe3 O4 , a soft ferrimagnet, formed after the SPS, rather than the hard SrFe12O19phaseparticle‐ size‐reduction effect. Additionally, they reported the highest maximal magnetic energy val‐ ues (*BH*) max, 4.0–4.6 kJ/m3 , for the samples with the lowest Fe3 O4 content, underlining the complexity of the involved demagnetization mechanism.

Vázquez‐Victorio [67] combined soft chemistry synthesis (polyol process) and consolidation by SPS to produce nanostructured BaFe12O19 barium hexaferrite. Typically, they produced an intermediate solid phase by reaction of the metallic salts in a polyol within an appropriate Ba/Fe atomic ratio; they were annealed at 800°C to complete the desired crystalline phase before SPS sintering at 800°C for 5–10 min and 100 MPa, under vacuum. Varying the nature of the metallic salts and the polyol solvent, they succeeded to produce highly dense (density > 95%) and ultrafine‐grained (∼100 nm) pure BaFe12O19and BaFe12O19 with a small content of iron oxide. A direct dependency of the magnetic properties of the produced solids on their iron oxide content was observed (**Figure 11**). The highest coercive field and magnetization

**Figure 11.** SEM micrographs of (a) BaFe12O19 and (c) BaFe12O19‐Fe2 O3 nanostructures produced by combining polyol process to SPS (800°C, 100 MPa, 5 min), and their room temperature hysteresis loops (b) [67].

at maximum applied field of 1 T (and hence the energy product *BH*) was measured at room temperature, on the nanocomposite. The hysteresis loops also appeared closer to a rectangu‐ lar form, which is the best shape for applications, as this leads to a well‐defined coercive field and remanent magnetization.

#### **6. Conclusions**

in the magnetization (*M*s) at maximum applied field of 1 T, for powders milled for a longer time and consolidated (**Figure 10**). They concluded that the magnetic properties of the studied nanocomposites are largely conditioned by the extrinsic properties of the secondary phase,

, a soft ferrimagnet, formed after the SPS, rather than the hard SrFe12O19phaseparticle‐ size‐reduction effect. Additionally, they reported the highest maximal magnetic energy val‐

**Figure 10.** Evolution of the coercive field *H*c (a), the remanent magnetization *M*r (a), the magnetization at maximum

different milling times. These data obtained from the hysteresis loops recorded at room temperature, applying the

O4

content, underlining the

max of SPS‐consolidated SrFe12O19 powder after

, for the samples with the lowest Fe3

Vázquez‐Victorio [67] combined soft chemistry synthesis (polyol process) and consolidation by SPS to produce nanostructured BaFe12O19 barium hexaferrite. Typically, they produced an intermediate solid phase by reaction of the metallic salts in a polyol within an appropriate Ba/Fe atomic ratio; they were annealed at 800°C to complete the desired crystalline phase before SPS sintering at 800°C for 5–10 min and 100 MPa, under vacuum. Varying the nature of the metallic salts and the polyol solvent, they succeeded to produce highly dense (density > 95%) and ultrafine‐grained (∼100 nm) pure BaFe12O19and BaFe12O19 with a small content of iron oxide. A direct dependency of the magnetic properties of the produced solids on their iron oxide content was observed (**Figure 11**). The highest coercive field and magnetization

Fe3 O4

ues (*BH*)

138 Sintering of Functional Materials

max, 4.0–4.6 kJ/m3

complexity of the involved demagnetization mechanism.

applied field 1 T, *M*s, and the maximum magnetic energy product (*BH*)

external magnetic field perpendicular to the uniaxial SPS pressing direction [66].

Currently, spark plasma sintering appears as the only method capable to consolidate nanopowders into high‐density nanostructured solids; in this chapter, we have briefly reviewed its application to carry out also the solid‐state reaction needed to achieve a par‐ ticular phase, starting from precursors synthesized by diverse methods. Many challenges remain, especially in the cases of reaction by precursor decomposition, when reaction and sintering temperatures are significantly different. RSPS is still a very young tech‐ nique, with many potential capabilities, which will certainly be developed in the near future.

## **Author details**

Giulia Franceschin<sup>1</sup> , Nancy Flores‐Martínez<sup>1</sup> , Gabriela Vázquez‐Victorio1,2, Souad Ammar1 and Raul Valenzuela1,2\*

\*Address all correspondence to: monjaras@unam.mx

1 ITODYS, Université Paris Diderot, Sorbonne Paris Cité, Paris Cedex, France

2 Institute for Materials Research, National Autonomous University of Mexico, Mexico City, México

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**Chapter 7**

**Provisional chapter**

**Selective Laser Sintering of Nanoparticles**

**Selective Laser Sintering of Nanoparticles**

DOI: 10.5772/intechopen.68872

Selective laser sintering of nanoparticles has received much attention recently as it enables rapid fabrication of functional layers including metal conductors and metal‐oxide elec‐ trodes on heat‐sensitive polymer substrate in ambient conditions. Photothermal reactions induced by lasers rapidly increase the local temperature of the target nanoparticle in a highly selective manner, and subsequent sintering steps including melting and coales‐ cence between nanoparticles occur to fabricate interconnected sintered films for various future applications. The mechanism of laser sintering, as well as possible target materials subject to laser sintering, together with experimental schemes developed to improve the

In this chapter, we focus on a specific type of sintering that utilizes laser and nanoparticle as its heat source and target material, respectively. Laser and nanoparticle on their own have interesting properties which are advantageous for conventional sintering process. Laser is a tool with a broad range of parameters and enables numerous responses such as remote temperature manipulation and rapid processing speed which cannot be achieved by other mechanical tools. Nanoparticles, having controllable sizes and shapes, find their application in various fields, and melting temperature depression due to their size effect is one of the key properties for sintering as it reduces the temperature required for the sintering process to a great extent. Laser sintering of nanoparticles—combining these two elements—not only possesses both the abovementioned features but also provides additional virtues and allows facile, damage‐free fabrication of functional layers on heat‐sensitive substrate to bring novel applications in the form of flexible electronics. Selective laser sintering of nanoparticles is

process and potential applications, is briefly summarized in this chapter.

**Keywords:** laser, optics, nanoparticle, metal, metal‐oxide, flexible electronics

© 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.

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.68872

Sukjoon Hong

**Abstract**

**1. Introduction**

Sukjoon Hong

**Provisional chapter**
