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

**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 m* 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, Si3

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.% Al2

, 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.% Si3 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.% Si3 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.% Si3

*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*

*m* 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.% Al2 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, Si3 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.% Si3 N4 , and Al-15 vol.% Al2 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

√

\_\_\_\_\_\_\_\_\_ 24 *Vf* Δ*α*Δ*T* \_\_\_\_\_\_\_\_\_ (1 − *Vf* ) *b rp*

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

*σOrowan* = \_\_\_\_\_\_ 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

*λ* ln \_\_*r*

*b* (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.% Si3 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.% Si3

Moreover, **Table 1** also shows that the tensile properties of the Al-1. 5 vol.% Si3

and Al-X (X = SiC, Si3

34 Sintering of Functional Materials

O3

vol.% Al2

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.% Si3 N4 , and micronsized 15 vol.% Al2 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, Si3 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-ZrO2

[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<EMAIL_ADDRESS>
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.% Si3

particles can be attributed to: (i) the lower CTE values of ceramic (SiC, Si3

sized 15 vol.% Al2

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, Si3

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<EMAIL_ADDRESS>
1 Center for Advanced Materials, Qatar University, Doha, Qatar

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

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[26] Penchal Reddy M, Ubaid F, Shakoor A, Mohamed AMA, Madhuri W. Structural and mechanical properties of microwave sintered Al-Ni50Ti50 composites. Journal of Science

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**Chapter 3**

**Provisional chapter**

) can be mixed with tricalcium phosphate

to the TCP‐40 wt% TiO2

= 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 MgF2

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

‐MgF2

**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 (MgF2

nia at 1200°C. Moreover, the addition of 4 wt% MgF2

(PO4 ) 2

leads to better mechanical properties (σr

obtained performances of the TCP‐TiO2

tissue and especially as enamel.

patibility of the β‐TCP and the high tribological properties of TiO2

30, 40 and 50 wt%) was studied. The performances of the TCP‐TiO2

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‐Ca3

Titania (TiO2

phosphate

**1. Introduction**

Tricalcium phosphate (β‐Ca3

(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 (MgF2 ) can be mixed with tricalcium phosphate (β‐Ca3 (PO4 ) 2 , β‐TCP) to make bioceramic composites, which would combine the biocom‐ patibility of the β‐TCP and the high tribological properties of TiO2 and MgF2 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‐TiO2 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% MgF2 to the TCP‐40 wt% TiO2 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‐TiO2 ‐MgF2 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‐Ca3 (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 (Al2 O3 ) or zirconia (ZrO2 ) 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 Al2 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% TiO2 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‐TiO2 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 TiO2 (2.5, 5, 7.5, 10, 20, 30, 40 and 50 wt%). In this study, the performances of the TCP‐TiO2 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% TiO2

ion (F−

tion of MgF2

tution of F−

in the HAp/ZrO2

for OH‐

of the TCP‐40 wt% TiO2

Brazilian test (σr

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

crystal structure [37]. Many studies show that, the addition of MgF2

cisely the enamel. Hence, MgF2

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, MgF2 crystallizes in the rutile‐type structure [26, 27]. Thanks to the higher compressibility of MgF2 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 MgF2

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, MgF2

study is to ameliorate the performances of the TCP‐40 wt% TiO2

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

TCP‐40 wt% TiO2 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% TiO2 composites by the inclusion of several additives to reach mechanical property similar to those of the bone tissues and pre‐ cisely the enamel. Hence, MgF2 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, MgF2 crystallizes in the rutile‐type structure [26, 27]. Thanks to the higher compressibility of MgF2 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 MgF2 in the HAp/ZrO2 composites [37]. As a result, the MgF2 (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 MgF2 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, MgF2 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% TiO2 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 (Al2

ducing the Al2

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 TiO2

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‐TiO2

5, 7.5, 10, 20, 30, 40 and 50 wt%). In this study, the performances of the TCP‐TiO2

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% TiO2

) or zirconia (ZrO2

) 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 (CaCO3 : Fluka, purity ≥ 98.5%) and calcium phosphate dibasic anhydrous (CaHPO4 : 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 MgF2

of β‐TCP, anatase‐TiO2

spectrometer.

with a gold layer.

 **(wt%) B′a**

A′, B′ and C′ are the weight rates of β‐TCP, TiO2

**Table 2.** The weight ratios of the different TCP‐TiO2

2010) [50].

**A′a**

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 (N2

 **(wt%) C′a**

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 MgF2

‐MgF2

, 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‐TiO2 ‐MgF2 composites were synthesized trical‐ cium phosphate (β‐TCP), magnesium fluoride MgF2 (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‐TiO2 , respectively. The calculated theoretical densities of each composite are illustrated in **Table 1**.

Different amounts of MgF2 (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 TiO2 and β‐TCP, respectively.

**Table 1.** The weight ratios of different TCP‐TiO2 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‐TiO2 and MgF2 , 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

‐MgF2

, 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%

) +H2

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 (CaCO3

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

(s) +CaCO3

98.5%) and calcium phosphate dibasic anhydrous (CaHPO4

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‐TiO2

cium phosphate (β‐TCP), magnesium fluoride MgF2

and the X‐ray diffraction (XRD).

48 Sintering of Functional Materials

**2. Materials and methods**

2CaHPO4

TiO2

TiO2

**Aa**

a

b

and anatase‐TiO2

following Eq. (3):

illustrated in **Table 1**.

Different amounts of MgF2

 **(wt%) Ba**

A and B are the weight rates of TiO2

**Table 1.** The weight ratios of different TCP‐TiO2

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 (N2 ) as the adsorbed gas (ASAP 2010) [50].

a A′, B′ and C′ are the weight rates of β‐TCP, TiO2 and MgF2 , respectively. bTheoretical density.

**Table 2.** The weight ratios of the different TCP‐TiO2 ‐MgF2 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 MgF2 37.00 0.05 0.10 3.15

liquid phase formation. Furthermore, the melting points of both β‐TCP and TiO2

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

and (c) TCP‐10 wt% TiO2

composites.

and 1855°C, respectively [22]. Thus, the liquid phase is formed between TiO2

found in the previous studies [12, 22]. They show that in the TiO2

**Figure 1.** DTA curves of: (a) β‐TCP, (b) anatase‐TiO2

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 TiO2

a

b

Mean diameter.

Theoretical density.

**Compounds SSA (m2**

(**Figure 1b**). The thermogram of the TCP‐10 wt% TiO2

**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, TiO2 and MgF2 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 TiO2 shows that no evolution can be obtained with the sintering temperature (**Figure 1b**). The thermogram of the TCP‐10 wt% TiO2 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, TiO2

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

Sρ (4)

π × D × t (5)

and MgF2

powders.

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

50 Sintering of Functional Materials

rupture strength (σr

using (Eq. (5)):

Eq. (6) [54]:

**3. Results**

DBET = \_\_\_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 = \_\_\_\_\_\_\_ 2 × P

**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‐TiO2 and (c) TCP‐10 wt% TiO2 composites.

liquid phase formation. Furthermore, the melting points of both β‐TCP and TiO2 are 1756 and 1855°C, respectively [22]. Thus, the liquid phase is formed between TiO2 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 TiO2 /β‐TCP system, there is an

**Figure 2.** Phase equilibrium diagram of the TiO2 ‐Ca3 (PO4 )2 system [22].

eutectic with a composition of 63 wt% TCP‐37 wt% TiO2 at 1380°C [22]. **Figure 2** confirms the presence of the liquid phase which is showed in the binary diagram of the Ca3 (PO4 )2 / TiO2 [22].

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

, (c) MgF2

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% TiO2

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% TiO2 (**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 TiO2

in the tricalcium phosphate matrix

composites.

and β‐TCP.

composites

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

and (d) TCP‐10 wt% TiO2

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‐TiO2

with 40 wt% TiO2

The mechanical strength of the TCP‐TiO2

**Figure 3.** Linear shrinkage curves of: (a) β‐TCP, (b) TiO2

results appears especially after the addition of 50 wt% TiO2

could be according to the formation of a liquid phase between TiO2

TiO2

TiO2

**Figure 3** shows the dilatometric measurements of different powders used in this study (β‐TCP, TiO2 , TCP‐10 wt% TiO2 composites and MgF2 ). The sintering temperature of the initial powder began at about 1000, 900 and 800°C for the β‐TCP, the TiO2 and the MgF2 , 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 MgF2 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% TiO2 composites. In fact, no evolution was reported with the TCP‐10 wt% TiO2 composites (**Figure 3d**). As a result, the content of 10 wt% TiO2 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 TiO2 .

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) TiO2 , (c) MgF2 and (d) TCP‐10 wt% TiO2 composites.

eutectic with a composition of 63 wt% TCP‐37 wt% TiO2

, TCP‐10 wt% TiO2

**Figure 2.** Phase equilibrium diagram of the TiO2

a result, the content of 10 wt% TiO2

TiO2

[22].

52 Sintering of Functional Materials

(β‐TCP, TiO2

the presence of the liquid phase which is showed in the binary diagram of the Ca3

**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 MgF2

initial powder began at about 1000, 900 and 800°C for the β‐TCP, the TiO2

‐Ca3 (PO4 )2

at about 1000°C [57], which confirms our results. The curve of the pure MgF2

(**Figure 3c**). **Figure 3d** presents the shrinkage curve of the TCP‐10 wt% TiO2

well‐explained by the exchange of Ca2+ and Ti4+ ions derived from β‐TCP and TiO2

In fact, no evolution was reported with the TCP‐10 wt% TiO2

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 MgF2

#### **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% TiO2 (**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‐TiO2 composites improves with both the content of TiO2 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‐TiO2 composites was hindered abruptly (**Figure 5**). The discrepancy of results appears especially after the addition of 50 wt% TiO2 in the tricalcium phosphate matrix (**Figure 5h**). The amelioration of these performances of the TCP‐40 wt% TiO2 composites could be according to the formation of a liquid phase between TiO2 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% TiO2 . 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 TiO2 : (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 TiO2 (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% TiO2 . Beyond 40 wt% TiO2 , the properties of the composites were hold up with the increase of the sintering temperature (**Figures 7** and **8**).

MgF2

full densification.

with 10 wt% MgF2

4 wt% of MgF2

, 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 TiO2

(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% MgF2

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

(**Figure 11**). The addition of MgF2

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‐

Sintering of the Tricalcium Phosphate-Titania-Magnesium Fluoride Composites

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

composites sintered at 1200°C

:

55

composites was obtained

outstandingly enhanced

of the densification was about 94% with 4 wt% MgF2

for 1 h with different percentages of MgF2

(**Figure 10**).

(**Figure 10**). Above 4 wt% MgF2

with the same percentages of MgF2

Vickers hardness of the TCP‐40 wt% TiO2

**Figure 10** shows the rupture strength of the TCP‐40 wt% TiO2

able amelioration of the rupture strength of the TCP‐40 wt% TiO2

#### **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% TiO2 composites sintered at 1200°C for 1 h with different amounts of MgF2 (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‐TiO2 composites increases from 3 wt% MgF2 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 TiO2 : (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 TiO2

(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% TiO2

(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 TiO2

54 Sintering of Functional Materials

TiO2

TiO2

the TCP‐TiO2

reached 33.1 and 15.7 GPa, respectively, with 40 wt% TiO2

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

strength. The densification behavior of the TCP‐40 wt% TiO2

composites increases from 3 wt% MgF2

for 1 h with different amounts of MgF2

MgF2 , then the density of the composites decreases abruptly (**Figure 9**). The optimum value of the densification was about 94% with 4 wt% MgF2 . This result is obviously close to the full densification.

**Figure 10** shows the rupture strength of the TCP‐40 wt% TiO2 composites sintered at 1200°C for 1 h with different percentages of MgF2 (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% MgF2 (**Figure 10**). Above 4 wt% MgF2 , the rupture strength decreases sharply (**Figure 10**). A remark‐ able amelioration of the rupture strength of the TCP‐40 wt% TiO2 composites was obtained with 10 wt% MgF2 (**Figure 10**).

The influence of the magnesium fluoride addition on Vickers hardness was studied at 1200°C with the same percentages of MgF2 (**Figure 11**). The addition of MgF2 outstandingly enhanced Vickers hardness of the TCP‐40 wt% TiO2 composites (**Figure 11**). Thus, the incorporation of 4 wt% of MgF2 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 TiO2 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% MgF2 , Vickers hardness of the composites was hindered (**Figure 11**). However, Vickers hardness of the TCP‐40 wt% TiO2 composites increases slowly with 10 wt% MgF2 (**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% MgF2

at: (a) 1000°C; (b) 1100°C, (c)

‐4 wt% MgF2

com‐

compos‐

**Figure 15** shows the evolution of Vickers hardness of the TCP‐38 wt% TiO2

**Figure 7.** Young's modulus of the β‐TCP sintered for 1 h with different percentages of TiO2

The evolution of the elastic moduli (E and G) of the TCP‐38 wt% TiO2

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% TiO2 composites sintered with different contents of MgF2 is shown in **Figure 12A** and **B**, respectively. The elastic modu‐ lus of the TCP‐40 wt% TiO2 composites increase with the addition of MgF2 (**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% MgF2 and these values reach 51 and 20 GPa, respectively. The performances of the samples were hindered with the increase of the percent‐ ages of MgF2 in the TCP‐40 wt% TiO2 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% TiO2 ‐4 wt% MgF2 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‐TiO2 ‐MgF2

**Figure 7.** Young's modulus of the β‐TCP sintered for 1 h with different percentages of TiO2 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% TiO2 ‐4 wt% MgF2 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% TiO2 ‐4 wt% MgF2 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% MgF2

composites are obtained by adding 4 wt% MgF2

in the TCP‐40 wt% TiO2

the densification of the TCP‐38 wt% TiO2

increases slowly with 10 wt% MgF2

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

56 Sintering of Functional Materials

with different contents of MgF2

lus of the TCP‐40 wt% TiO2

TiO2

ages of MgF2

time (**Figure 13**).

was hindered (**Figure 11**). However, Vickers hardness of the TCP‐40 wt% TiO2

The evolution of the elastic modulus (E and G) of the TCP‐40 wt% TiO2

(**Figure 11**).

**Figure 6.** Vickers hardness versus temperature of the β‐TCP sintered for 1 h with different percentages of TiO2

, 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 MgF2

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% MgF2

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‐TiO2

composites

at: (a)

composites sintered

(**Figure 12A** and

‐MgF2

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 TiO2 at: (a) 1000°C, (b) 1100°C, (c) 1200°C and (d) 1300°C.

**Figure 10.** Rupture strength of the TCP‐40 wt% TiO2

**Figure 11.** Vickers hardness of the TCP‐40 wt% TiO2

of MgF2 .

of MgF2 . 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% TiO2 composites sintered at 1200°C for 1 h with different percentages of MgF2 .

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% TiO2 composites sintered at 1200°C for 1 h with different percentages of MgF2 .

**Figure 8.** Shear modulus of the β‐TCP sintered for 1 h with different percentages of TiO2

1200°C and (d) 1300°C.

58 Sintering of Functional Materials

**Figure 9.** Relative density of the TCP‐40 wt% TiO2

MgF2 . 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% TiO2 composites sintered at 1200°C for 1 h with different percentages of MgF2 .

**Figure 13.** Relative density of the TCP‐38 wt% TiO2

**Figure 14.** Rupture strength of the TCP‐38 wt% TiO2

the sintering time.

the sintering time.

‐4 wt% MgF2

‐4 wt% MgF2

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% TiO2 composites sintered at 1200°C for 1 h with different percentages of MgF2 : (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 MgF2 (1, 2.5, 4, 4.5, 5, 7.5 and 10 wt%). The diffraction pattern of the TCP‐40 wt% TiO2 composites sintered without MgF2 displays the α‐TCP phase traces pres‐ ence (ICDD data file no. 43‐1484), the calcium titanate (CaTiO3 ) (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% MgF2 to the TCP‐40 wt% TiO2 led to the presence of traces of both calcium titanate (CaTiO3 ) 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 MgF2 to the TCP‐TiO2 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% TiO2 ‐4 wt% MgF2 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 MgF2

ence (ICDD data file no. 43‐1484), the calcium titanate (CaTiO3

no. 65‐1119) (**Figure 17a**). The introduction of 1 wt% MgF2

the presence of traces of both calcium titanate (CaTiO3

(**Figure 17b**). Further increasing the percentage of MgF2

(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% TiO2

to the TCP‐TiO2

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 MgF2

**Figure 12.** Elastic modulus of the TCP‐40 wt% TiO2

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

TCP‐40 wt% TiO2

60 Sintering of Functional Materials

MgF2

**Figure 14.** Rupture strength of the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites sintered at 1200°C for different length of the sintering time.

**Figure 15.** Vickers hardness of the TCP‐38 wt% TiO2 ‐4 wt% MgF2 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 Mg2 (PO4 )F (ICDD data file no. 42‐0582) was detected after the addi‐ tion of 2.5, 4 and 4.5 wt% MgF2 (**Figure 17c**–**e**). At higher contents of MgF2 (after 4.5 wt%), the Mg2 (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% TiO2 composites sintered at 1200°C for 1 h without and with different percentages of MgF2 (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% TiO2 composites sintered at 1200°C for 1 h without and with different amounts of MgF2 (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% TiO2 composites sintered without MgF2 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% TiO2 ‐1 wt% MgF2 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

Mg2 (PO4

wt% MgF2

composites (**Figure 19f**–**h**).

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

**Figure 16.** Elastic modulus of the TCP‐38 wt% TiO2

shows the micrographs of the TCP‐40 wt% TiO2

)F (**Figure 19b**–**e**). The intensity of the first peak increases with the amount of MgF2

‐4.5

composites sintered at 1200°C for 1 h

composites sintered at 1200°C for different length of the

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63

while the intensity of the second peak decreases and disappears with the TCP‐40 wt% TiO2

‐4 wt% MgF2

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% MgF2

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% TiO2

)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 MgF2

(**Figure 17c**–**e**). At higher contents of MgF2

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

a new phase relative to Mg2

(PO4

(**Figure 17f**–**h**).

the sintering time.

62 Sintering of Functional Materials

the Mg2

MgF2

tion of 2.5, 4 and 4.5 wt% MgF2

**Figure 15.** Vickers hardness of the TCP‐38 wt% TiO2

(PO4

**Figure 18** displays the FTIR spectra of the TCP‐40 wt% TiO2

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 MgF2

The 31P MAS‐NMR spectra of the TCP‐40 wt% TiO2

without and with different amounts of MgF2

MAS‐NMR spectrum of the TCP‐40 wt% TiO2

**Figure 16.** Elastic modulus of the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites sintered at 1200°C for different length of the sintering time: (A) Young's modulus and (B) the Shear modulus.

Mg2 (PO4 )F (**Figure 19b**–**e**). The intensity of the first peak increases with the amount of MgF2 while the intensity of the second peak decreases and disappears with the TCP‐40 wt% TiO2 ‐4.5 wt% MgF2 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% TiO2 composites sintered at 1200°C for 1 h

**Figure 17.** XRD patterns of the TCP‐40 wt% TiO2 composites sintered at 1200°C for 1 h without and with different percentages of MgF2 : (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, +: CaTiO3 , \*: rutile, ♣: fluorapatite, ●: Mg2 (PO4 )F).

without and with different percentages of MgF2

**Figure 19.** 31P MAS‐NMR spectra of the TCP‐40 wt% TiO2

composites sintered with 2.5 wt% MgF2

composites sintered without MgF2

the formation of a new lamella form (**Figure 20b**). At higher amount of MgF2

sification (94%) was registered after the addition of 4 wt% MgF2

‐7.5 wt% MgF2

form (**Figure 20e**–**f**). Moreover, the micrographs of the TCP‐35 wt% TiO2

Moreover, the microstructure of the TCP‐38 wt% TiO2

Both the addition of 7.5 and 10 wt% of MgF2

grains of the TCP and the titania (**Figure 20a**). The microstructure of the TCP‐40 wt% TiO2 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% TiO2

different percentages of MgF2

the TCP‐36.25 wt% TiO2

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

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65

was completely transformed (**Figure 20b**).

‐4 wt% MgF2

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% TiO2

composites shows

‐10 wt% MgF2

(4 wt%),

**Figure 18.** Infrared spectra of the TCP‐40 wt% TiO2 composites sintered at 1200°C for 1 h without and with different percentages of MgF2 : (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% TiO2 composites sintered at 1200°C for 1 h without and with different percentages of MgF2 : (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% TiO2

**Figure 18.** Infrared spectra of the TCP‐40 wt% TiO2

percentages of MgF2

, \*: rutile, ♣: fluorapatite, ●: Mg2

percentages of MgF2

(o: α‐TCP, ♦: β‐TCP, +: CaTiO3

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 MgF2 (2.5, 4, 7.5 and 10 wt%). The micrograph of the TCP‐40 wt% TiO2 composites sintered without MgF2 depicts the coalescence of the grains of the TCP and the titania (**Figure 20a**). The microstructure of the TCP‐40 wt% TiO2 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% MgF2 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 MgF2 (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% MgF2 to the TCP‐40 wt% TiO2 composites (**Figure 20c**–**d**). This result can be explicated by the reduction in the porosity. Moreover, the microstructure of the TCP‐38 wt% TiO2 ‐4 wt% MgF2 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 MgF2 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% TiO2 ‐7.5 wt% MgF2 composites sintered at 1200°C presents the bubbles form (**Figure 20e**–**f**). Moreover, the micrographs of the TCP‐35 wt% TiO2 ‐10 wt% MgF2 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‐TiO2

addition of a reinforcement agent like magnesium fluoride (MgF2

on the mechanical properties of the TCP‐40 wt% TiO2

ison between the previous work derived from the TCP‐40 wt% TiO2

‐4 wt% MgF2

(the TCP‐38 wt% TiO2

. 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]. MgF2

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% MgF2

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 TiO2

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% MgF2

and H = 360 Hv) of the TCP‐40 wt% TiO2

tively. The mechanical properties of the TCP‐40 wt% TiO2

of MgF2

the amounts of MgF2

ent work, with adding MgF2

the TCP‐38 wt% TiO2

**Figure 20.** SEM micrographs of the TCP‐40 wt% TiO2 composites sintered at 1200°C for 1 h without and with different percentages of MgF2 : (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 TiO2 . The properties of the TCP‐TiO2 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 (MgF2 ). 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]. MgF2 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 MgF2 on the mechanical properties of the TCP‐40 wt% TiO2 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% TiO2 composites increase as in terms of the amounts of MgF2 . The optimum values (ρ = 94%, σr = 27 MPa, E = 51 GPa, G = 20 GPa and H = 360 Hv) of the TCP‐40 wt% TiO2 composites are attained after the sintering process at 1200°C for 1 h with the addition of 4 wt% MgF2 . 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% TiO2 composites and the pres‐ ent work, with adding MgF2 (the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites) leads to the same conclusion described between enamel and dentin. However, as far as hardness is considered, the TCP‐38 wt% TiO2 ‐4 wt% MgF2 composites are considered harder than the TCP‐40 wt%

**4. Discussion**

percentages of MgF2

66 Sintering of Functional Materials

**Figure 20.** SEM micrographs of the TCP‐40 wt% TiO2

(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% TiO2 ‐4 wt% MgF2 composites in biomedical applications and specifically as enamel.

The values of the elastic modulus (E and G) of the TCP‐38 wt% TiO2

values with the elastic modulus of the TCP‐40 wt% TiO2

The evolution of the mechanical properties (E, G, H and σr

MgF2

MgF2

wt% TiO2

(Ca5 (PO4 )3

[Mg2 (PO4

wt% MgF2

wt% TiO2

wt% MgF2

the Mg2

phase (Mg2

to the TCP‐TiO2

percentages of MgF2

(PO4

amounts of MgF2

‐4 wt% MgF2

FAp with more quantities in the TCP‐TiO2

proves the presence of traces of CaTiO3