High-Entropy Alloy Composites

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

Cao Z, Ruan H, Li T. A promising new class of high-temperature alloys: Eutectic high-entropy alloys. Sci. Rep. 2015;4:6200. DOI: 10.1038/srep06200

[54] Vaidya M, Muralikrishna G M, Murty B S. High-entropy alloys by mechanical alloying: A review. J. Mater. Res. 2019;34:664-686. DOI: 10.1557/

jmr.2019.37

[45] Corbett J D, Garcia E, Guloy A M, Hurng W M, Kwon Y U, Leon-Escamilla E A. Widespread interstitial chemistry of Mn5Si3-type and related phases. Hidden impurities and opportunities. Chem. Mater. 1998;10:2824-2836. DOI:

10.1021/cm980223c

[46] Kitagawa J, Hamamoto S.

10.7566/JPSCP.30.011055

actamat.2016.08.081

net/SSP.130.15

s41598-017-15766-y

PhysRev.125.837

PhysRev.92.874

actamat.2013.01.042

[48] Izumi F, Momma K. Threedimensional visualization in powder diffraction. Solid State Phenom. 2007; 130:15-20. DOI: 10.4028/www.scientific.

Superconductivity in Nb5Ir3−xPtxO. JPS Conf. Proc. 2020;30:011055. DOI:

[47] Miracle D B, Senkov O N. A critical review of high entropy alloys and related concepts. Acta Mater. 2017;122:448-511. DOI: 10.1016/j.

[49] Tsubota M, Kitagawa J. A necessary criterion for obtaining accurate lattice parameters by Rietveld method. Sci. Rep. 2017;7:15381. DOI: 10.1038/

[50] Giorgi A L, Szklarz E G, Storms E K, Bowman A L, Matthias B T. Effect of Composition on the superconducting transition temperature of tantalum carbide and niobium carbide. Phys. Rev.

temperatures of superconductors. Phys. Rev. 1953;92:874-876. DOI: 10.1103/

[52] Otto F, Yang Y, Bei H, George E P. Relative effects of enthalpy and entropy on the phase stability of equiatomic high-entropy alloys. Acta Materialia 2013;61:2628-2638. DOI: 10.1016/j.

[53] Lu Y, Dong Y, Guo S, Jiang L, Kang H, Wang T, Wen B, Wang Z, Jie J,

1962;125:837-838. DOI: 10.1103/

[51] Matthias B T. Transition

**36**

**39**

**Chapter 3**

**Abstract**

gations are reported in this Chapter.

microstructural characterization

**1. Introduction**

Why Al-B4C Metal Matrix

*Mohamed F. Ibrahim, Hany R. Ammar, Agnes M. Samuel,* 

*Mahmoud S. Soliman, Victor Songmene and Fawzy H. Samuel*

The Al-B4C metal matrix composite (MMC) is characterized by its ability to absorb neutrons which makes it the most suitable shielding material for nuclear reactors. The present work was performed on two series of Al-B4C metal matrix composites made using a powder injection apparatus. In one series, commercially pure aluminum (A5) served as the matrix. For the second set, 6063 alloy was used. In all cases the volume fraction of B4C reinforcement particles (grit size 400 mesh, purity 99.5%) was approximately 15%. The volume fraction of the injected B4C particles was determined using a computer driven image analyzer. Measured amounts of Ti, Zr, and Ti + Zr, were added to the molten composites of both series. Microstructural characterization was carried out employing a field emission scanning electron microscope operating at 20 kV and equipped with an electron dispersive x-ray spectroscopic system (EDS). The same technique was applied to characterize the fracture behavior of the tested composites. Mechanical properties of these composites were investigated using impact testing, and ambient and high temperature tensile testing methods. Almost 1000 impact and tensile samples were tested following different heat treatments. The obtained results from these investi-

**Keywords:** MMC, precipitation hardening, FESEM, tensile testing, impact testing,

The Al-B4C metal matrix composite (MMC) is characterized by its high thermal conductivity and its ability to absorb neutrons which makes it a suitable shielding material [1]. Increasing the concentration of B4C (>30%) increases the composite strength as well as its neutron absorption capacity. Roy et al. [2] suggested the use of 7xxx alloys as base material for the MMC due to its low density and its hardening ability caused by heat treatment which would contribute to the strength of the MMC. The use of 2124 Al alloy composites reinforced with B4C particulates has been proposed by Öksüz and Oskay [3]. The authors claim that the volumetric wear rates of the 2124 Al alloy and its composites are increased with increase in the applied load. Singla et al. [4] proposed the use of molten technique for the production of Al-B4C MMC. The authors studied an MMC made of Al-7075 alloy as the matrix and B4C 32 μm particulate as the reinforcement agent. Mohan

Composites? A Review

## **Chapter 3**

## Why Al-B4C Metal Matrix Composites? A Review

*Mohamed F. Ibrahim, Hany R. Ammar, Agnes M. Samuel, Mahmoud S. Soliman, Victor Songmene and Fawzy H. Samuel*

#### **Abstract**

The Al-B4C metal matrix composite (MMC) is characterized by its ability to absorb neutrons which makes it the most suitable shielding material for nuclear reactors. The present work was performed on two series of Al-B4C metal matrix composites made using a powder injection apparatus. In one series, commercially pure aluminum (A5) served as the matrix. For the second set, 6063 alloy was used. In all cases the volume fraction of B4C reinforcement particles (grit size 400 mesh, purity 99.5%) was approximately 15%. The volume fraction of the injected B4C particles was determined using a computer driven image analyzer. Measured amounts of Ti, Zr, and Ti + Zr, were added to the molten composites of both series. Microstructural characterization was carried out employing a field emission scanning electron microscope operating at 20 kV and equipped with an electron dispersive x-ray spectroscopic system (EDS). The same technique was applied to characterize the fracture behavior of the tested composites. Mechanical properties of these composites were investigated using impact testing, and ambient and high temperature tensile testing methods. Almost 1000 impact and tensile samples were tested following different heat treatments. The obtained results from these investigations are reported in this Chapter.

**Keywords:** MMC, precipitation hardening, FESEM, tensile testing, impact testing, microstructural characterization

#### **1. Introduction**

The Al-B4C metal matrix composite (MMC) is characterized by its high thermal conductivity and its ability to absorb neutrons which makes it a suitable shielding material [1]. Increasing the concentration of B4C (>30%) increases the composite strength as well as its neutron absorption capacity. Roy et al. [2] suggested the use of 7xxx alloys as base material for the MMC due to its low density and its hardening ability caused by heat treatment which would contribute to the strength of the MMC. The use of 2124 Al alloy composites reinforced with B4C particulates has been proposed by Öksüz and Oskay [3]. The authors claim that the volumetric wear rates of the 2124 Al alloy and its composites are increased with increase in the applied load. Singla et al. [4] proposed the use of molten technique for the production of Al-B4C MMC. The authors studied an MMC made of Al-7075 alloy as the matrix and B4C 32 μm particulate as the reinforcement agent. Mohan

and Kennedy [5] investigated the machinability of Al-(7 and 14) wt.% Si alloys reinforced with B4C. The MMCs were developed using the stir casting technique. Their results show that the composite reinforced with B4C with a particle size of 100 nanometers has better mechanical properties and wear behavior compared to those reinforced with 24-micron or 6-micron sized particulates. Vaidya et al. [6] found that the strength of B4C particle reinforced Al 6061 composite was significantly greater than the unreinforced alloy.

Drilling experiments were conducted by Kumar et al. [7] on 6061 alloy-15%B4C (220 μm particulate diameter) using a vertical machine with High Speed Steel drills of 6 mm, 9 mm and 12 mm diameter under dry drilling conditions. It was found that speed, design of the experiment and drill diameter have a marked influence on the Over-Cut (half the difference of the diameter of the hole produced to the tool diameter). Topcu [8] and Manjunatha et al. [9] used the powder atomization technique to produce Al-5% B4C and Al-15%B4C MMCs. The authors reported that the wear resistance increased in proportion to the amount of the boron carbide reinforced. Tribo-surface characteristics of two aluminum metal matrix composites (Al-MMC) of compositions Al–13 vol%B4C and Al–13 vol%SiC sliding against a commercial phenolic brake pad under dry conditions were investigated by Shorowordi et al. [10–12]. The friction coefficient was found to decrease slightly at high contact pressure.

The wear rate and friction coefficient of Al–B4C was lower than that of Al–SiC. Several studies on friction behavior involving Al-MMC friction against ferrous materials revealed that during sliding, a layer, termed as mechanically mixed layer (MML), was formed on the worn surface of the Al-MMC [13–17]. Such layer, however, was not found to form on unreinforced aluminum. Several researchers [18–22] studied the production of Al-11%B4C using stir melt technique. The 6061 alloy was the matrix to which B4C particles were added. Prior to addition, the B4C particles were preheated along with K2TiF6 halide salt. The resulting composite was found to have improved mechanical properties compared to the base alloy. Uthayakumar et al. [23] performed a study on the wear performance of Al–5%SiC–5%B4C hybrid composites under dry sliding conditions using a pin on disc tribometer method. The main conclusion was that the hybrid composites can retain the wear resistance properties up to 60 N load and sliding speed ranges of 1–4 m/s.

Comparison of microstructural and mechanical properties of Al–10 vol% TiC, Al–10 vol% B4C and Al–5 vol% TiC–5 vol% B4C composites prepared by casting techniques was made by Mazaheri et al. [24]. The results show that the wear behavior of Al-B4C MMC is the best among the three composites studied. The wettability of B4C particulates was investigated by Toptan et al. [25]. They found that addition of Ti leads to formation of thin layers (80–180 nm in thickness) of Ti-C and Ti-B around the B4C particulates which would solve the wettability issue. Similar observation on titanium as one of the reactive metals that can be used to increase wettability in Al-B4C system was reported by other researchers [26–31]. According to Wang et al. [32] and Yang et al. [33], the stress distribution within a particlereinforced composite subjected to external loading is non-uniform. Nanostructured Al–B4C composite sheets were processed by accumulative roll bonding (ARB), and the effect of the number of ARB cycles on the distribution of the B4C particles in the Al matrix was evaluated by Yazdani and Salahinejad [34] who noted an improvement in the reinforcement distribution by increasing the ARB cycles.

The present chapter summarizes the work that was carried out by the present authors using two types of Al-B4C composites: (i) a mechanically alloyed composite supplied by Ceradyne Canada ULC, a 3 M Company, Chicoutimi, Québec, Canada, and (ii) an in-house made composite using powder injection at the Université du Québec a Chicoutimi [35–44].

**41**

**Figure 1.**

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

Reinforcement powder (grit size 400 and 95.4% purity) additions of 15 vol. % were made using a powder injection apparatus (**Figure 1**). The B4C particulate was injected into molten Al. Fe, Ti and Zr additions were introduced into the molten bath, using Al-25%Fe, Al-10% Ti and Al-15% Zr master alloys, respectively, whereas Mg and Si were added as pure elements. Chemical compositions of the investigated composites are listed in **Table 1**. A general view of the powder injection set-up showing a schematic of the injection system is shown in **Figure 1**. It consists

In order to ensure uniform distribution of the B4C particulates, the molten composite melt was stirred vigorously (300 rpm) at 730 ± 5°C. Thereafter, the molten composite was poured in two different metallic molds preheated at 450°C, as shown in **Figure 2**: an L-shaped mold (3.5x 3.8 x 30.5 cm) which was used for microstructure characterization, and a book-type mold (4 x 17 x 34 cm). In order to determine

**2. Experimental procedure**

**2.1 Composite preparation**

of the following components:

i. a fluidizer tube

iii. resistance heating coils

vii.flow diversion baffles.

ii. a carrier tube and a quartz nozzle

iv. an adjustable two-dimensional movable stand

vi. an impeller (stirrer) with adjustable rotation speed

v. a melting unit with resistance heating

*A schematic diagram of the powder injector used in the present work.*

## **2. Experimental procedure**

## **2.1 Composite preparation**

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

cantly greater than the unreinforced alloy.

speed ranges of 1–4 m/s.

Québec a Chicoutimi [35–44].

and Kennedy [5] investigated the machinability of Al-(7 and 14) wt.% Si alloys reinforced with B4C. The MMCs were developed using the stir casting technique. Their results show that the composite reinforced with B4C with a particle size of 100 nanometers has better mechanical properties and wear behavior compared to those reinforced with 24-micron or 6-micron sized particulates. Vaidya et al. [6] found that the strength of B4C particle reinforced Al 6061 composite was signifi-

Drilling experiments were conducted by Kumar et al. [7] on 6061 alloy-15%B4C (220 μm particulate diameter) using a vertical machine with High Speed Steel drills of 6 mm, 9 mm and 12 mm diameter under dry drilling conditions. It was found that speed, design of the experiment and drill diameter have a marked influence on the Over-Cut (half the difference of the diameter of the hole produced to the tool diameter). Topcu [8] and Manjunatha et al. [9] used the powder atomization technique to produce Al-5% B4C and Al-15%B4C MMCs. The authors reported that the wear resistance increased in proportion to the amount of the boron carbide reinforced. Tribo-surface characteristics of two aluminum metal matrix composites (Al-MMC) of compositions Al–13 vol%B4C and Al–13 vol%SiC sliding against a commercial phenolic brake pad under dry conditions were investigated by Shorowordi et al. [10–12]. The friction coefficient was found to decrease slightly at high contact pressure. The wear rate and friction coefficient of Al–B4C was lower than that of Al–SiC. Several studies on friction behavior involving Al-MMC friction against ferrous materials revealed that during sliding, a layer, termed as mechanically mixed layer (MML), was formed on the worn surface of the Al-MMC [13–17]. Such layer, however, was not found to form on unreinforced aluminum. Several researchers [18–22] studied the production of Al-11%B4C using stir melt technique. The 6061 alloy was the matrix to which B4C particles were added. Prior to addition, the B4C particles were preheated along with K2TiF6 halide salt. The resulting composite was found to have improved mechanical properties compared to the base alloy. Uthayakumar et al. [23] performed a study on the wear performance of Al–5%SiC–5%B4C hybrid composites under dry sliding conditions using a pin on disc tribometer method. The main conclusion was that the hybrid composites can retain the wear resistance properties up to 60 N load and sliding

Comparison of microstructural and mechanical properties of Al–10 vol% TiC, Al–10 vol% B4C and Al–5 vol% TiC–5 vol% B4C composites prepared by casting techniques was made by Mazaheri et al. [24]. The results show that the wear behavior of Al-B4C MMC is the best among the three composites studied. The wettability of B4C particulates was investigated by Toptan et al. [25]. They found that addition of Ti leads to formation of thin layers (80–180 nm in thickness) of Ti-C and Ti-B around the B4C particulates which would solve the wettability issue. Similar observation on titanium as one of the reactive metals that can be used to increase wettability in Al-B4C system was reported by other researchers [26–31]. According to Wang et al. [32] and Yang et al. [33], the stress distribution within a particlereinforced composite subjected to external loading is non-uniform. Nanostructured Al–B4C composite sheets were processed by accumulative roll bonding (ARB), and the effect of the number of ARB cycles on the distribution of the B4C particles in the Al matrix was evaluated by Yazdani and Salahinejad [34] who noted an improve-

ment in the reinforcement distribution by increasing the ARB cycles.

The present chapter summarizes the work that was carried out by the present authors using two types of Al-B4C composites: (i) a mechanically alloyed composite supplied by Ceradyne Canada ULC, a 3 M Company, Chicoutimi, Québec, Canada, and (ii) an in-house made composite using powder injection at the Université du

**40**

Reinforcement powder (grit size 400 and 95.4% purity) additions of 15 vol. % were made using a powder injection apparatus (**Figure 1**). The B4C particulate was injected into molten Al. Fe, Ti and Zr additions were introduced into the molten bath, using Al-25%Fe, Al-10% Ti and Al-15% Zr master alloys, respectively, whereas Mg and Si were added as pure elements. Chemical compositions of the investigated composites are listed in **Table 1**. A general view of the powder injection set-up showing a schematic of the injection system is shown in **Figure 1**. It consists of the following components:


In order to ensure uniform distribution of the B4C particulates, the molten composite melt was stirred vigorously (300 rpm) at 730 ± 5°C. Thereafter, the molten composite was poured in two different metallic molds preheated at 450°C, as shown in **Figure 2**: an L-shaped mold (3.5x 3.8 x 30.5 cm) which was used for microstructure characterization, and a book-type mold (4 x 17 x 34 cm). In order to determine

**Figure 1.** *A schematic diagram of the powder injector used in the present work.*

#### *Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*


#### **Table 1.**

*Codes and compositions of the MMCs used in this study.*

#### **Figure 2.**

*(a) L-shaped mold, (b) book-type mold, (c) L-shaped casting, (d) book-mold casting.*

the solidification rate obtained from each mold, trials were made using Al-7%Si. **Figure 3** depicts the dendrite arm spacing (DAS) and grain size corresponding to each mold. The castings made using the book-mold were hot rolled into slabs of 1-3 mm thickness, depending on the type of test carried out.

#### **2.2 Microstructural investigation**

Samples for microstructural characterization were prepared from the L-shaped mold casting in the as cast condition using 5A composite. The volume fraction and average size of the B4C particles was measured using Clemex image analyzer. Fracture surfaces were examined of samples sectioned from both tensile- and impact tested bars. The samples were examined using Hitachi S-7000 and Hitachi SU-8000 FE-SEM microscopes equipped with EDS facilities at McGill University, Montreal.

#### **2.3 Mechanical testing**

Charpy impact testing was carried out on un-notched test specimens (10x 10 x 55 mm). The samples were sectioned from the L-shaped mold castings and heat treated in an electrical air forced furnace. An instrumented Charpy impact testing machine, equipped with a data acquisition unit was employed to measure the load,

**43**

**Figure 5.**

**Figure 3.**

**Figure 4.** *Hot rolled sheets.*

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

for each composite/condition were reported.

total absorbed energy (Et) to fracture. The mean values of 6 impact-tested samples

Slabs (25x 20x 400 mm) were prepared from the book mold castings. Prior to rolling using a four cylinder mill, the slabs were annealed at 500°C for 16 h. The last two passes were carried out at room temperature to straighten the rolled slabs (sheets) – see **Figure 4**. **Figure 5** shows the dimensions of samples prepared from

*(a, c) Optical micrographs of Al-7%Si alloy for (a) block casting, 60* μ*m; (c) L-shaped casting, 30* μ*m;* 

*(b, d) Macrographs showing grain size in (b) block casting; (d) L-shaped casting.*

*Typical sample for room and high temperature tensile testing (dimensions are in mm).*

total absorbed energy (Et) to fracture. The mean values of 6 impact-tested samples for each composite/condition were reported.

Slabs (25x 20x 400 mm) were prepared from the book mold castings. Prior to rolling using a four cylinder mill, the slabs were annealed at 500°C for 16 h. The last two passes were carried out at room temperature to straighten the rolled slabs (sheets) – see **Figure 4**. **Figure 5** shows the dimensions of samples prepared from

#### **Figure 3.**

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

**Alloy Code Composition (Ti, Zr, Sc in wt%)**

3A Al-15v/oB4C + 0.45%Ti + 0.25%Zr 4A Al-15v/oB4C + 0.45%Ti + 0.15%Sc

3B 6063-15v/oB4C + 0.45%Ti + 0.25%Zr 4B 6063-15v/oB4C + 0.45%Ti + 0.15%Sc

5A Al-15v/oB4C + 0.45%Ti + 0.15%Sc + 0.25%Zr

5B 6063-15v/oB4C + 0.45%Ti + 0.15%Sc + 0.25%Zr

1A Al-15v/oB4C

1B 6063-15v/oB4C

*Codes and compositions of the MMCs used in this study.*

2A Al-15v/oB4C + 0.45%Ti

2B 6063-15v/oB4C + 0.45%Ti

the solidification rate obtained from each mold, trials were made using Al-7%Si. **Figure 3** depicts the dendrite arm spacing (DAS) and grain size corresponding to each mold. The castings made using the book-mold were hot rolled into slabs of

Samples for microstructural characterization were prepared from the L-shaped

Charpy impact testing was carried out on un-notched test specimens (10x 10 x 55 mm). The samples were sectioned from the L-shaped mold castings and heat treated in an electrical air forced furnace. An instrumented Charpy impact testing machine, equipped with a data acquisition unit was employed to measure the load,

mold casting in the as cast condition using 5A composite. The volume fraction and average size of the B4C particles was measured using Clemex image analyzer. Fracture surfaces were examined of samples sectioned from both tensile- and impact tested bars. The samples were examined using Hitachi S-7000 and Hitachi SU-8000 FE-SEM microscopes equipped with EDS facilities at McGill University, Montreal.

1-3 mm thickness, depending on the type of test carried out.

*(a) L-shaped mold, (b) book-type mold, (c) L-shaped casting, (d) book-mold casting.*

**2.2 Microstructural investigation**

**2.3 Mechanical testing**

**42**

**Figure 2.**

**Table 1.**

*(a, c) Optical micrographs of Al-7%Si alloy for (a) block casting, 60* μ*m; (c) L-shaped casting, 30* μ*m; (b, d) Macrographs showing grain size in (b) block casting; (d) L-shaped casting.*

**Figure 4.** *Hot rolled sheets.*

#### **Figure 5.**

*Typical sample for room and high temperature tensile testing (dimensions are in mm).*

the rolled sheets and used for room and high temperature tensile testing. Tensile samples (matrix is aluminum) were solutionised at 620°C for 24 h. In spite of the fact that pure aluminum normally is not heat treatable, it could benefit from the precipitation of Zr-rich particles during aging. The 6063/B4C/15p composite samples were solutionized at 540°C to minimize surface oxidation (MgO). After solution heat treatment, the tensile bars were quenched in warm water (60° C), followed by aging for 10 h at 200, 300 and 400° C, and then air cooling. Room temperature testing was carried out using an MTS Servohydraulic mechanical testing machine at a strain rate of 4 x 10−4/s.

High temperature testing was done at strain rate of 5 x 10−4/s in a temperature range 25–500°C. In all cases, tensile properties were measured: ultimate tensile strength (UTS), the 0.2% offset yield strength (YS) and percentage elongation (%El). For each working condition, at least five specimens were tested and mean values were reported (SD ±5%). Microstructure and fracture behavior of selected samples were examined using optical microscopy and Field Emission Scanning Electron Microscopy (FESEM) techniques.

## **3. Results and discussion**

#### **3.1 Microstructural characterization (as cast condition)**

The main function of the addition of Zr and Ti, is to protect the B4C particles from reacting with the molten Al [45–48]. **Figure 6(a)** depicts the microstructure of a specimen sectioned from the L-shaped castings, revealing a uniform distribution of B4C particles throughout the matrix. From such micrographs, the volume fraction of B4C particles was determined (~15 vol.%). According to the Al–Ti binary diagram [49], at 730°C an amount of 0.5 wt-%Ti could be added to the molten composite. From the reported findings of Tahiri et al. [50–54] it was reported that increasing the concentration of Ti in the molten alloy above 0.5 wt.% will increase the temperature of the molten alloy. As a result, fluidity of the composite will be markedly reduced caused by the segregation of B4C particles as exhibited in **Figure 6(b)**.

FESEM examination of 5A composite treated with Ti revealed that in addition to B4C particles, possible precipitation of several intermetallics mainly, TiB2, TiC and traces of AlB24C4, Al4C3, Al3BC and AlB12, along with the primary intermetallic phases TiAl, Ti3Al and TiAl3 could also occur. It is expected that the formation of these phases in layers would lead to an improvement in the adhesion between the matrix and the B4C reinforcement [21–23]. **Figure 7** displays electron micrographs of alloy B,

#### **Figure 6.**

*Secondary electron micrographs showing: (a) a uniform distribution of B4C in matrix of base alloy using in-house powder injecting technique, (b) segregation of B4C (white circle).*

**45**

**Figure 8.**

**Figure 7.**

*towards matrix.*

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

*SE images from composite B showing (a) regular B4C particles protected by layers of Zr–Ti rich particles, (II) irregular forms of B4C showing partial reaction with matrix forming AlBC compound; (b) B4C particle surrounded by two layers of Zr–Ti rich particles; and (c) B4C particle showing progress of cumulative reaction* 

*Element distribution: (a) backscattered electron image-note formation of several particles around a B4C* 

*particle, white arrows, (b) boron, (c) alumium, (d) carbon, (e) titanium, (f) zirconium.*

#### **Figure 7.**

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

machine at a strain rate of 4 x 10−4/s.

Electron Microscopy (FESEM) techniques.

**3.1 Microstructural characterization (as cast condition)**

**3. Results and discussion**

exhibited in **Figure 6(b)**.

the rolled sheets and used for room and high temperature tensile testing. Tensile samples (matrix is aluminum) were solutionised at 620°C for 24 h. In spite of the fact that pure aluminum normally is not heat treatable, it could benefit from the precipitation of Zr-rich particles during aging. The 6063/B4C/15p composite samples were solutionized at 540°C to minimize surface oxidation (MgO). After solution heat treatment, the tensile bars were quenched in warm water (60° C), followed by aging for 10 h at 200, 300 and 400° C, and then air cooling. Room temperature testing was carried out using an MTS Servohydraulic mechanical testing

High temperature testing was done at strain rate of 5 x 10−4/s in a temperature range 25–500°C. In all cases, tensile properties were measured: ultimate tensile strength (UTS), the 0.2% offset yield strength (YS) and percentage elongation (%El). For each working condition, at least five specimens were tested and mean values were reported (SD ±5%). Microstructure and fracture behavior of selected samples were examined using optical microscopy and Field Emission Scanning

The main function of the addition of Zr and Ti, is to protect the B4C particles from reacting with the molten Al [45–48]. **Figure 6(a)** depicts the microstructure of a specimen sectioned from the L-shaped castings, revealing a uniform distribution of B4C particles throughout the matrix. From such micrographs, the volume fraction of B4C particles was determined (~15 vol.%). According to the Al–Ti binary diagram [49], at 730°C an amount of 0.5 wt-%Ti could be added to the molten composite. From the reported findings of Tahiri et al. [50–54] it was reported that increasing the concentration of Ti in the molten alloy above 0.5 wt.% will increase the temperature of the molten alloy. As a result, fluidity of the composite will be markedly reduced caused by the segregation of B4C particles as

FESEM examination of 5A composite treated with Ti revealed that in addition to B4C particles, possible precipitation of several intermetallics mainly, TiB2, TiC and traces of AlB24C4, Al4C3, Al3BC and AlB12, along with the primary intermetallic phases TiAl, Ti3Al and TiAl3 could also occur. It is expected that the formation of these phases in layers would lead to an improvement in the adhesion between the matrix and the B4C reinforcement [21–23]. **Figure 7** displays electron micrographs of alloy B,

*Secondary electron micrographs showing: (a) a uniform distribution of B4C in matrix of base alloy using* 

*in-house powder injecting technique, (b) segregation of B4C (white circle).*

**44**

**Figure 6.**

*SE images from composite B showing (a) regular B4C particles protected by layers of Zr–Ti rich particles, (II) irregular forms of B4C showing partial reaction with matrix forming AlBC compound; (b) B4C particle surrounded by two layers of Zr–Ti rich particles; and (c) B4C particle showing progress of cumulative reaction towards matrix.*

#### **Figure 8.**

*Element distribution: (a) backscattered electron image-note formation of several particles around a B4C particle, white arrows, (b) boron, (c) alumium, (d) carbon, (e) titanium, (f) zirconium.*

**Figure 9.**

*B4C-matrix interactions: (a) formation of Al-C, (b) formation of Zr rich phases-see Figure 7(c).*

**Figure 10.**

*EDS spectra obtained from Figure 9(b) confirming the B4C-matrix interaction and the dependence of the composition of outcome on its position with respect to B4C particles.*

containing Ti and Zr. As can be seen in **Figure 7(a)**, the B4C particles are surrounded by several layers of Zr–Ti rich phases (area marked I). Area marked II shows B4C particles that have partially reacted with the matrix due to formation of the layer of AlBC existing in the matrix. In area III, fine B4C particles are found be transformed completely into AlBC compounds. **Figure 7(b)** reveals a B4C particle surrounded by a thin layer of Ti-rich phase followed by several layers of Zr–Ti rich phases. Some of these Zr–Ti rich phase particles are seen to grow into the aluminum matrix.

**Figure 8** is produced from 5A composite remelted for multiple times at 730°C. **Figure 8(a)** shows the distribution of the B4C particles. The B, Al and C distribution are presented in **Figures 8(b), 8(c), 8(d)**, respectively. Another point to be considered is that Ti covers the entire surface of the B4C particle (similar to C and B) whereas Zr is limited to the layer decorating the B4C particle. It is inferred from **Figure 9** that the layers surrounding the B4C particles are a mixture of Al-Ti, and Al-C-Ti compounds. The EDS spectra obtained from **Figure 9(b)** are presented in **Figure 10**. Based on these EDSs, areas near the B4C particles could be made of Al-B-Zr compound whereas those away are probably Al-Ti-Zr compound [55].

#### **4. Mechanical properties**

#### **4.1 Impact testing**

Total energies (Et) produced from the ten studied composites in the solutionized condition are shown in **Figure 11(a)**. Apparently the absorbed energy of the composite depends on what matrix is used and the volume fraction of the undissolved intermetallics. Following aging at 200°C for 10 h (**Figure 11(b)**), the precipitation of Zr-and Sc containing phases [55–58] led to significant decrease in the values of Et which may be attributed to precipitation of Mg2Si phase particles

**47**

and Zr elements.

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

during aging in particular in B-series. It is inferred from **Figure 11(b)** that precipitation of Zr and/or Sc phases has an insignificant effect on the absorbed energy in series A-composites. According to Fuller et al. [59] aging 6063 alloy at 300°C would result in alloy softening due to coarsening of Mg2Si phases particles. Simultaneous precipitation of Zr-rich phase may lead to balancing the composite toughness to some extent (**Figure 11(c)**). Aging at higher temperatures i.e. 400°C for 10 h resulted in coarsening of all types of precipitated phases causing important improvement in the composite toughness, regardless of the type of the matrix

*Total absorbed energies of the present composites: (a) SHT, (b) aging at 200°C, (c) aging at 300°C, (d) aging* 

Fracture mechanism of Al-2%Cu composite was investigated by Miserez [60]. The study showed that the fracture may occur in two stages: (i) particle fracture leading to void nucleation in the matrix, and (ii) voids nucleated in the matrix in areas of high stress concentrations. The blue arrow in **Figure 12(a)** shows that the crack is propagating through the protecting layer surrounded by stacking faults (white arrows). On the left hand side of the micrograph several stacking faults appear in the form of steps (white arrows). Two distinctive types of cracks were observed in **Figure 12(b)**: cracks that took place at the interior of the B4C particles and continued through the protecting layers i.e. intergranular, or those occurring at the B4C/matrix interfaces (black arrow). No particle debonding was observed due to the existence of the protecting layers as displayed in **Figure 12(c)**. The microstructure beneath the fracture surface (vertical section-loading direction) shown in **Figure 12(d)** demonstrates the coherency between the B4C particles and the surrounding matrix. Aging the composite at 400°C led to marked coarsening of the Al3Zr phase particles as seen in **Figure 13(a)**, which explains the improvement in the composite toughness in **Figure 11(d)**. The EDS spectrum in **Figure 13(b)** corresponds to the circled area in **Figure 13(a)** revealing strong reflections from Al

used, as exhibited in **Figure 11(d)**.

**Figure 11.**

*at 400°C.*

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

**Figure 11.**

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

*B4C-matrix interactions: (a) formation of Al-C, (b) formation of Zr rich phases-see Figure 7(c).*

*EDS spectra obtained from Figure 9(b) confirming the B4C-matrix interaction and the dependence of the* 

containing Ti and Zr. As can be seen in **Figure 7(a)**, the B4C particles are surrounded by several layers of Zr–Ti rich phases (area marked I). Area marked II shows B4C particles that have partially reacted with the matrix due to formation of the layer of AlBC existing in the matrix. In area III, fine B4C particles are found be transformed completely into AlBC compounds. **Figure 7(b)** reveals a B4C particle surrounded by a thin layer of Ti-rich phase followed by several layers of Zr–Ti rich phases. Some of

**Figure 8** is produced from 5A composite remelted for multiple times at 730°C. **Figure 8(a)** shows the distribution of the B4C particles. The B, Al and C distribution are presented in **Figures 8(b), 8(c), 8(d)**, respectively. Another point to be considered is that Ti covers the entire surface of the B4C particle (similar to C and B) whereas Zr is limited to the layer decorating the B4C particle. It is inferred from **Figure 9** that the layers surrounding the B4C particles are a mixture of Al-Ti, and Al-C-Ti compounds. The EDS spectra obtained from **Figure 9(b)** are presented in **Figure 10**. Based on these EDSs, areas near the B4C particles could be made of Al-B-Zr compound whereas those away are probably Al-Ti-Zr compound [55].

Total energies (Et) produced from the ten studied composites in the solutionized condition are shown in **Figure 11(a)**. Apparently the absorbed energy of the composite depends on what matrix is used and the volume fraction of the undissolved intermetallics. Following aging at 200°C for 10 h (**Figure 11(b)**), the precipitation of Zr-and Sc containing phases [55–58] led to significant decrease in the values of Et which may be attributed to precipitation of Mg2Si phase particles

these Zr–Ti rich phase particles are seen to grow into the aluminum matrix.

*composition of outcome on its position with respect to B4C particles.*

**46**

**4. Mechanical properties**

**4.1 Impact testing**

**Figure 10.**

**Figure 9.**

*Total absorbed energies of the present composites: (a) SHT, (b) aging at 200°C, (c) aging at 300°C, (d) aging at 400°C.*

during aging in particular in B-series. It is inferred from **Figure 11(b)** that precipitation of Zr and/or Sc phases has an insignificant effect on the absorbed energy in series A-composites. According to Fuller et al. [59] aging 6063 alloy at 300°C would result in alloy softening due to coarsening of Mg2Si phases particles. Simultaneous precipitation of Zr-rich phase may lead to balancing the composite toughness to some extent (**Figure 11(c)**). Aging at higher temperatures i.e. 400°C for 10 h resulted in coarsening of all types of precipitated phases causing important improvement in the composite toughness, regardless of the type of the matrix used, as exhibited in **Figure 11(d)**.

Fracture mechanism of Al-2%Cu composite was investigated by Miserez [60]. The study showed that the fracture may occur in two stages: (i) particle fracture leading to void nucleation in the matrix, and (ii) voids nucleated in the matrix in areas of high stress concentrations. The blue arrow in **Figure 12(a)** shows that the crack is propagating through the protecting layer surrounded by stacking faults (white arrows). On the left hand side of the micrograph several stacking faults appear in the form of steps (white arrows). Two distinctive types of cracks were observed in **Figure 12(b)**: cracks that took place at the interior of the B4C particles and continued through the protecting layers i.e. intergranular, or those occurring at the B4C/matrix interfaces (black arrow). No particle debonding was observed due to the existence of the protecting layers as displayed in **Figure 12(c)**. The microstructure beneath the fracture surface (vertical section-loading direction) shown in **Figure 12(d)** demonstrates the coherency between the B4C particles and the surrounding matrix. Aging the composite at 400°C led to marked coarsening of the Al3Zr phase particles as seen in **Figure 13(a)**, which explains the improvement in the composite toughness in **Figure 11(d)**. The EDS spectrum in **Figure 13(b)** corresponds to the circled area in **Figure 13(a)** revealing strong reflections from Al and Zr elements.

#### **Figure 12.**

*Fracture characteristics of the present composites: (a) stacking faults-SHT, (b) cracks- aging at 200°C/10 h, (c) B4C/matrix coherency, (d) vertical section beneath (c) confirming particle/matrix bonding-note the severe reaction around some of the B4C particles - circled areas.*

#### **4.2 Tensile testing**

#### *4.2.1 Room temperature testing*

The stress–strain curves of two aluminum matrix composites in the SHT condition and after aging at 200°C/h are shown in **Figure 14**. The main observation to be made is the slow working hardening rate illustrated by low work hardening and the slow increase in the composite UTS (**Figure 14(a)**). As a consequence of aging at 200°C/10 h, the UTS increased by approximately 80 MPa which may be attributed to the precipitation of Al3Zr phase particles [61]. Considering the solutionizing treatment of 5B composite the maximum attainable strength is about 280 MPa- **Figure 14(b)**. Using a heat treatable matrix i.e. 6063 alloy, the composite revealed significant improvement in both the UTS levels as well as work hardening rate as displayed in **Figure 14(c)** which may be caused by the precipitation of Mg2Si phase particles during the storing period prior to testing (~10 minutes at room temperature). As expected, aging at 200°C/10 h resulted in increasing the composite strength from 280 MPa to 500 MPa, **Figure 14(d)**, which may be interpreted in terms of simultaneous precipitation of both Mg2Si and Al3Zr phase particles.

**49**

**Figure 13.**

*circle in (a).*

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

Following the solution heat treatment of composite 5A, the fracture surface is characterized by the formation of deep dimple network as demonstrated in **Figure 15(a)**. Some of these dimples revealed the presence of deformation bands (arrowed) due to composite ductility. The marking seen on the surface of the B4C particles in **Figure 15(b)** may be caused by gradual fracture of the particles, maintaining at the same time their coherency with the aluminum matrix. Precipitation of Al3Zr phase particles during aging at 200°C/10 h is clearly seen in **Figure 15(c)**. Due to reduction in the composite ductility, some of the B4C particles were cracked as shown by the white arrows in the same figure. In the case of 6063 alloy matrix, with the significant increase in the composite UTS level following aging at 200 C/10 h (500 MPa), cracks are seen to initiate and propagate through the Zr-Ti protecting layer as demonstrated in **Figure 15(d)** – see blue arrow. No B4C particle

*(a) Fracture surface of composite 5B aged at 400°C/10 h, (b) EDS spectrum corresponding to white* 

The 5A and 5B composites (**Table 1**) were tested in the temperature range of 25–500°C and their corresponding stress–strain curves are displayed in **Figure 16**. Composite 5B showed a slightly higher strength compared to composite 5A. It should

debonding is observed to take place under axial loading.

*4.2.2 High temperature testing*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

The stress–strain curves of two aluminum matrix composites in the SHT condition and after aging at 200°C/h are shown in **Figure 14**. The main observation to be made is the slow working hardening rate illustrated by low work hardening and the slow increase in the composite UTS (**Figure 14(a)**). As a consequence of aging at 200°C/10 h, the UTS increased by approximately 80 MPa which may be attributed to the precipitation of Al3Zr phase particles [61]. Considering the solutionizing treatment of 5B composite the maximum attainable strength is about 280 MPa- **Figure 14(b)**. Using a heat treatable matrix i.e. 6063 alloy, the composite revealed significant improvement in both the UTS levels as well as work hardening rate as displayed in **Figure 14(c)** which may be caused by the precipitation of Mg2Si phase particles during the storing period prior to testing (~10 minutes at room temperature). As expected, aging at 200°C/10 h resulted in increasing the composite strength from 280 MPa to 500 MPa, **Figure 14(d)**, which may be interpreted in terms of simultaneous precipitation of both Mg2Si and Al3Zr phase particles.

*Fracture characteristics of the present composites: (a) stacking faults-SHT, (b) cracks- aging at 200°C/10 h, (c) B4C/matrix coherency, (d) vertical section beneath (c) confirming particle/matrix bonding-note the severe* 

**48**

**4.2 Tensile testing**

**Figure 12.**

*4.2.1 Room temperature testing*

*reaction around some of the B4C particles - circled areas.*

#### **Figure 13.**

*(a) Fracture surface of composite 5B aged at 400°C/10 h, (b) EDS spectrum corresponding to white circle in (a).*

Following the solution heat treatment of composite 5A, the fracture surface is characterized by the formation of deep dimple network as demonstrated in **Figure 15(a)**. Some of these dimples revealed the presence of deformation bands (arrowed) due to composite ductility. The marking seen on the surface of the B4C particles in **Figure 15(b)** may be caused by gradual fracture of the particles, maintaining at the same time their coherency with the aluminum matrix. Precipitation of Al3Zr phase particles during aging at 200°C/10 h is clearly seen in **Figure 15(c)**. Due to reduction in the composite ductility, some of the B4C particles were cracked as shown by the white arrows in the same figure. In the case of 6063 alloy matrix, with the significant increase in the composite UTS level following aging at 200 C/10 h (500 MPa), cracks are seen to initiate and propagate through the Zr-Ti protecting layer as demonstrated in **Figure 15(d)** – see blue arrow. No B4C particle debonding is observed to take place under axial loading.

#### *4.2.2 High temperature testing*

The 5A and 5B composites (**Table 1**) were tested in the temperature range of 25–500°C and their corresponding stress–strain curves are displayed in **Figure 16**. Composite 5B showed a slightly higher strength compared to composite 5A. It should

#### **Figure 14.**

*Stress–strain diagrams corresponding to: (a) Al/B4C 5A composite - SHT, (b) Al/B4C composite aged at 200°C/ 10 h, (c) 5B 6063/B4C composite - SHT, (d) 6063/B4C composite aged at 200°C/ 10 h.*

**51**

**Figure 19(c)**.

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

relationships as described by Eq. 1 [66–70]:

Hollomon parameter. At a constant *ε*, *σ<sup>n</sup>*

slope of (Qa/nR) as shown in **Figure 17** [66–68].

show a mixture of ductile and brittle types of fracture [70].

be mentioned here that samples of composite 5B were tested in the T4 condition which involves natural aging. Increasing the testing temperature up to 450°C resulted in significant increase in the composite pct. Elongation to failure. Aging at further higher temperature would lead to precipitation of Al3Zr which would result in reduc-

For aging at temperatures higher than 0.5 of the melting temperature (Tm), there is a similarity between creep and hot deformation. Under this condition, the relation between the measured parameters can be expressed using power law

> exp *Qa <sup>n</sup> A Z RT* = =

where *ε* = strain rate, *σ* = flow stress, n = stress exponent, *Qa* = activation energy, *R* = gas constant, *T* = absolute temperature, A is a constant and *Z* = Zener-

exp *<sup>n</sup> Qa <sup>B</sup>*

where *B* = constant. Differentiation of Eq. 2 coupled with (1/T), gives *Qa* as:

( ) ln 1

 σ

can expressed as:

*RT* <sup>=</sup>

*Qa*

*nR <sup>T</sup>* <sup>∂</sup> <sup>=</sup> <sup>∂</sup> σ

Applying these equations, the plot of ln*σ vs* 1/*T*, will give a straight line with a

The fracture surface of composite 5A tested at 25°C was characterized by the presence of deformation bands covering the internal surface of the dimples as shown previously. **Figure 18(a)** exhibits the fracture surface of 5A composite tested at 250°C revealing multiple contour-type markings (white arrow) due to the high ductility ~15%. Testing at 350°C resulted in major increase in the deformation bands in the form of steps (blue arrows) as shown in **Figure 18(b)**. The thin white arrows point to cracked B4C particles. The black arrow in **Figure 18(c)** - 5A composite pulled to fracture at 450°C- indicates the presence of a long crack within the protective layer. In addition, the 5A composite exhibited elongated dimples as depicted in **Figure 18(c)**. **Figure 18(d)** is an enlarged portion of the crack in **Figure 18(c)**.From the associated EDS spectrum in **Figure 18(e)**, the possibility of precipitation of a large amount of Zr-rich particles, which would explain the reduction in the composite ductility when tested at this temperature. Fractographic observations made by Zhang et al. [69] on 6092/(B4C)p indicated the possibility of several interfacial bonding characteristics such as good bonding (extruded composites) and weak bonding (hot isostatic pressing). The fracture surfaces of composites would also

The fracture surface of 5B composite pulled to fracture at 250°C revealed that in addition to the deep dimples, some stacking faults could also be seen in the fracture surface (**Figure 19(a)**-white arrows). As in the case of 5A composite, at 350°C, the fracture surface exhibited a well-defined dimple structure as a result of the increase in the composite % elongation to fracture, **Figure 19(b)**-black arrows point to the thickness of the protection layer. Due to the strong particle/matrix interface adhesion, some of the B4C particles have been cracked at their interior. In this case, the crack was initiated at the particle/matrix interface and propagated through the particle. When the sample was tested at 450°C, the fracture surface revealed the formation of very large and deep dimple network as shown in

(1)

(2)

(3)

ing the composite ductility, as shown in **Figure 16(c)** [6, 62–65].

ε

σ

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

**50**

**Figure 15.**

**Figure 14.**

*Fracture surface of tensile tested samples of 5A composite: (a,b) SHT, (c,d) aging at 200 C/10h.*

*Stress–strain diagrams corresponding to: (a) Al/B4C 5A composite - SHT, (b) Al/B4C composite aged at* 

*200°C/ 10 h, (c) 5B 6063/B4C composite - SHT, (d) 6063/B4C composite aged at 200°C/ 10 h.*

be mentioned here that samples of composite 5B were tested in the T4 condition which involves natural aging. Increasing the testing temperature up to 450°C resulted in significant increase in the composite pct. Elongation to failure. Aging at further higher temperature would lead to precipitation of Al3Zr which would result in reducing the composite ductility, as shown in **Figure 16(c)** [6, 62–65].

For aging at temperatures higher than 0.5 of the melting temperature (Tm), there is a similarity between creep and hot deformation. Under this condition, the relation between the measured parameters can be expressed using power law relationships as described by Eq. 1 [66–70]:

$$
\dot{\varepsilon} \exp\left(\frac{Q\_u}{RT}\right) = A\sigma^n = Z \tag{1}
$$

where *ε* = strain rate, *σ* = flow stress, n = stress exponent, *Qa* = activation energy, *R* = gas constant, *T* = absolute temperature, A is a constant and *Z* = Zener-Hollomon parameter. At a constant *ε*, *σ<sup>n</sup>* can expressed as:

$$
\sigma^\* = B \exp\left(\frac{Q\_a}{RT}\right) \tag{2}
$$

where *B* = constant. Differentiation of Eq. 2 coupled with (1/T), gives *Qa* as:

$$\frac{\partial \ln \sigma}{\partial \left(\bigvee\_{T} \right)} = \frac{Q\_{\omega}}{nR} \tag{3}$$

Applying these equations, the plot of ln*σ vs* 1/*T*, will give a straight line with a slope of (Qa/nR) as shown in **Figure 17** [66–68].

The fracture surface of composite 5A tested at 25°C was characterized by the presence of deformation bands covering the internal surface of the dimples as shown previously. **Figure 18(a)** exhibits the fracture surface of 5A composite tested at 250°C revealing multiple contour-type markings (white arrow) due to the high ductility ~15%. Testing at 350°C resulted in major increase in the deformation bands in the form of steps (blue arrows) as shown in **Figure 18(b)**. The thin white arrows point to cracked B4C particles. The black arrow in **Figure 18(c)** - 5A composite pulled to fracture at 450°C- indicates the presence of a long crack within the protective layer. In addition, the 5A composite exhibited elongated dimples as depicted in **Figure 18(c)**. **Figure 18(d)** is an enlarged portion of the crack in **Figure 18(c)**.From the associated EDS spectrum in **Figure 18(e)**, the possibility of precipitation of a large amount of Zr-rich particles, which would explain the reduction in the composite ductility when tested at this temperature. Fractographic observations made by Zhang et al. [69] on 6092/(B4C)p indicated the possibility of several interfacial bonding characteristics such as good bonding (extruded composites) and weak bonding (hot isostatic pressing). The fracture surfaces of composites would also show a mixture of ductile and brittle types of fracture [70].

The fracture surface of 5B composite pulled to fracture at 250°C revealed that in addition to the deep dimples, some stacking faults could also be seen in the fracture surface (**Figure 19(a)**-white arrows). As in the case of 5A composite, at 350°C, the fracture surface exhibited a well-defined dimple structure as a result of the increase in the composite % elongation to fracture, **Figure 19(b)**-black arrows point to the thickness of the protection layer. Due to the strong particle/matrix interface adhesion, some of the B4C particles have been cracked at their interior. In this case, the crack was initiated at the particle/matrix interface and propagated through the particle. When the sample was tested at 450°C, the fracture surface revealed the formation of very large and deep dimple network as shown in **Figure 19(c)**.

**Figure 16.**

*(a) Stress–strain curves obtained from 5A composite, (b) stress–strain curves obtained from 5B composite, (c) % elongation as a function of testing temperatire (composite 1 is 5A, composite 2 is 5B).*

**53**

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

*4.2.3 Effect of strain rate*

*Relationship of flow stress vs 1/T (T = in kelvin degree).*

follows

**Figure 17.**

The main characterized parameters of hot deformation or creep behavior of commercial Al alloys and Al-based composites are by high values of *na* (> 5) as well as the activation energy *Qa*. These values are higher than those for solute diffusion [71–75]. This behavior can be explained in terms of interaction of dislocations with the dispersed strengthening particles resulting in a threshold stress *σo*. In this case, the deformation process is related to an effective stress, *σe* = (*σ* -*σo*) not the applied stress *σ*. Therefore, equations 1-2 can be rewritten to take into consideration *σo* as

( ) <sup>0</sup> exp /

*Gb Q RT Z A kT G* <sup>−</sup> = = ′

where: *A* = constant, *k* = Boltzmann's constant, *b* = magnitude of Burgers vector,

The true stress–strain curves obtained from testing the 5A composite tested at 300°C (a), 400°C (b) and 500°C (c), are respectively depicted in **Figure 20**. These curves can be divided into three stages; strain hardening where the stress increases with strain until reaches a steady state. Stage 2 represents maximum stress, followed stage 3 where necking takes place leading to failure. Generally speaking, with increasing the strain rate would result in an increase in the flow stress. The effect of testing temperature on the behavior of the stress- strain curves at a constant strain rate of 10−3 s−1 is displayed in **Figure 20(d)**. Increasing the testing temperature led to an increase in the composite ductility at 500°C and higher strain rates higher than

10−2 s−1. The ductility was decreased in temperature range of 350° C–450°C.

The relationship between the strain rate, *ε*̇ and stress *σ*, at a constant temperature, is governed by plotting *ε̇ vs. σ* applying a log log scale (**Figure 21**) for different testing temperatures. The results reported in **Figure 20** may suggest that the data points at each testing temperature fall on a straight line with a constant na that increases from 5.8 at 500°C to ~7 at 350–450°C, thereafter to 10.4 at 300°C. The high values of *na* are close to those obtained for commercial aluminum alloys [76–80] and metal matrix composites. As mentioned before, dislocations -second phase

*t*

ε

*G* = shear modulus and *Qt* = true activation energy.

*nt*

σ σ

(4)

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

*(a) Stress–strain curves obtained from 5A composite, (b) stress–strain curves obtained from 5B composite,* 

*(c) % elongation as a function of testing temperatire (composite 1 is 5A, composite 2 is 5B).*

**52**

**Figure 16.**

**Figure 17.** *Relationship of flow stress vs 1/T (T = in kelvin degree).*

#### *4.2.3 Effect of strain rate*

The main characterized parameters of hot deformation or creep behavior of commercial Al alloys and Al-based composites are by high values of *na* (> 5) as well as the activation energy *Qa*. These values are higher than those for solute diffusion [71–75]. This behavior can be explained in terms of interaction of dislocations with the dispersed strengthening particles resulting in a threshold stress *σo*. In this case, the deformation process is related to an effective stress, *σe* = (*σ* -*σo*) not the applied stress *σ*. Therefore, equations 1-2 can be rewritten to take into consideration *σo* as follows

$$\dot{\varepsilon} \exp(Q\_{\omega} \prime RT) = Z = A \prime \frac{Gb}{kT} \left(\frac{\sigma - \sigma\_0}{G}\right)^{u\_t} \tag{4}$$

where: *A* = constant, *k* = Boltzmann's constant, *b* = magnitude of Burgers vector, *G* = shear modulus and *Qt* = true activation energy.

The true stress–strain curves obtained from testing the 5A composite tested at 300°C (a), 400°C (b) and 500°C (c), are respectively depicted in **Figure 20**. These curves can be divided into three stages; strain hardening where the stress increases with strain until reaches a steady state. Stage 2 represents maximum stress, followed stage 3 where necking takes place leading to failure. Generally speaking, with increasing the strain rate would result in an increase in the flow stress. The effect of testing temperature on the behavior of the stress- strain curves at a constant strain rate of 10−3 s−1 is displayed in **Figure 20(d)**. Increasing the testing temperature led to an increase in the composite ductility at 500°C and higher strain rates higher than 10−2 s−1. The ductility was decreased in temperature range of 350° C–450°C.

The relationship between the strain rate, *ε*̇ and stress *σ*, at a constant temperature, is governed by plotting *ε̇ vs. σ* applying a log log scale (**Figure 21**) for different testing temperatures. The results reported in **Figure 20** may suggest that the data points at each testing temperature fall on a straight line with a constant na that increases from 5.8 at 500°C to ~7 at 350–450°C, thereafter to 10.4 at 300°C. The high values of *na* are close to those obtained for commercial aluminum alloys [76–80] and metal matrix composites. As mentioned before, dislocations -second phase

#### **Figure 18.**

*Fracture surface of tensile tested samples of 5A composite: (a) 250°C, (b) 350°C, (c) 450 C, (d) an enlarged portion of (c) showing the crack, (e) EDS corresponding to (c) revealing reflections due to Al, Zr, Ti elements.*

**55**

**Figure 22.**

**Figure 20.**

**Figure 21.**

*Strain rate and stress relationship in the temperature range 300–500°C.*

*Fracture surface of samples tested at 300°C: (a) 10−2 s−1, (b) 10−4 s−1.*

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

*True stress–strain curves at different strain rates at (a) 300°C, (b) 400°C, (c) 500°C, (d) strain rate 10−3 s−1.*

**Figure 19.** *Fracture behavior of 5B composite tested at: (a) 250°C, (b) 350°C, (c) 450°C.*

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

*Fracture surface of tensile tested samples of 5A composite: (a) 250°C, (b) 350°C, (c) 450 C, (d) an enlarged portion of (c) showing the crack, (e) EDS corresponding to (c) revealing reflections due to Al, Zr, Ti elements.*

**54**

**Figure 19.**

**Figure 18.**

*Fracture behavior of 5B composite tested at: (a) 250°C, (b) 350°C, (c) 450°C.*

**Figure 20.** *True stress–strain curves at different strain rates at (a) 300°C, (b) 400°C, (c) 500°C, (d) strain rate 10−3 s−1.*

**Figure 21.** *Strain rate and stress relationship in the temperature range 300–500°C.*

**Figure 22.** *Fracture surface of samples tested at 300°C: (a) 10−2 s−1, (b) 10−4 s−1.*

particles interaction would lead to high values of *na* and *Qa* threshold stress in the composite materials.

**Figure 22 (a)** shows the fracture surface of the samples tested at 300°C (strain rate of 10−2 s−1), consisting of a mixture of small dimples and intragranular fracture. **Figure 22 (b)** is the fracture surface at strain rate of 10−4 s−1 at 300°C, exhibiting larger dimples with precipitation of Al3Zr in particles at their interiors (circled areas) [81].

## **5. Summary**

The results obtained from the present investigations revealed that the powder injection technique used in our study proved to be effective in producing composites with a uniform distribution of B4C particles throughout the matrix (commercial aluminum or 6063 alloy). The combined addition of Zr and Ti improved the possibility of increasing the number of B4C particles in the matrix by improving the particulate wettability. The precipitating Al3(Zr1–xTix) particles decorating the B4C particles were found to grow into the surrounding matrix. Precipitation of Mg2Si in 6063/B4C was more effective in controlling the composite toughness than Al3Zr in the under-aging conditions. Overaging occurred at 400°C for prolonged aging times (i.e. 10 h), resulting in a significant improvement in the composite toughness regardless the type of the matrix. Cracks were always initiated at the particle/matrix interfaces and propagated either through the B4C particles or along the protecting Al3(Zr1–xTix) layer. No particle debonding was observed regardless the type of matrix or the testing method. Formation of the Zr/Ti rich layers surrounding the B4C particles strengthen their adhesion to the surrounding matrix. Increasing the testing temperature leads to rapid decrease in the composite strength in an exponential pattern which appeared in the gradual fracturing of the reinforcement B4C particles. The plots of flow stress as a function of testing temperature are linear with a fitting factor of 0.955. The value of *nt* ~ 5 and *Qt* of 130 kJ mol−1 along with subgrain formation may conclude that dislocation climb is the main controlling process. Similar observation was made in pure Al with dispersed particles. The pct elongation to failure reached a maximum value at intermediate value of *Z*, which can determine the optimum conditions for the composite formability.

**57**

**Author details**

Saudi Arabia

Canada

Mohamed F. Ibrahim1

Mahmoud S. Soliman3

\*, Hany R. Ammar2

\*Address all correspondence to: fhsamuel@uqac.ca

provided the original work is properly cited.

, Victor Songmene4

1 Université du Québec à Chicoutimi, Chicoutimi, Québec, G7H 2B1, Canada

2 Department of Mechanical Engineering, Qassim University, Saudi Arabia

3 Department of Mechanical Engineering, King Saud University, Riyadh,

4 Department of Mechanical Eng., École de Technologie Supérieure, Québec,

© 2021 The Author(s). Licensee IntechOpen. 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,

, Agnes M. Samuel1

and Fawzy H. Samuel1

,

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

composite materials.

areas) [81].

**5. Summary**

particles interaction would lead to high values of *na* and *Qa* threshold stress in the

**Figure 22 (a)** shows the fracture surface of the samples tested at 300°C (strain rate of 10−2 s−1), consisting of a mixture of small dimples and intragranular fracture. **Figure 22 (b)** is the fracture surface at strain rate of 10−4 s−1 at 300°C, exhibiting larger dimples with precipitation of Al3Zr in particles at their interiors (circled

The results obtained from the present investigations revealed that the powder injection technique used in our study proved to be effective in producing composites with a uniform distribution of B4C particles throughout the matrix (commercial

possibility of increasing the number of B4C particles in the matrix by improving the particulate wettability. The precipitating Al3(Zr1–xTix) particles decorating the B4C particles were found to grow into the surrounding matrix. Precipitation of Mg2Si in 6063/B4C was more effective in controlling the composite toughness than Al3Zr in the under-aging conditions. Overaging occurred at 400°C for prolonged aging times (i.e. 10 h), resulting in a significant improvement in the composite toughness regardless the type of the matrix. Cracks were always initiated at the particle/matrix interfaces and propagated either through the B4C particles or along the protecting Al3(Zr1–xTix) layer. No particle debonding was observed regardless the type of matrix or the testing method. Formation of the Zr/Ti rich layers surrounding the B4C particles strengthen their adhesion to the surrounding matrix. Increasing the testing temperature leads to rapid decrease in the composite strength in an exponential pattern which appeared in the gradual fracturing of the reinforcement B4C particles. The plots of flow stress as a function of testing temperature are linear with a fitting factor of 0.955. The value of *nt* ~ 5 and *Qt* of 130 kJ mol−1 along with subgrain formation may conclude that dislocation climb is the main controlling process. Similar observation was made in pure Al with dispersed particles. The pct elongation to failure reached a maximum value at intermediate value of *Z*, which

aluminum or 6063 alloy). The combined addition of Zr and Ti improved the

can determine the optimum conditions for the composite formability.

**56**

## **Author details**

Mohamed F. Ibrahim1 \*, Hany R. Ammar2 , Agnes M. Samuel1 , Mahmoud S. Soliman3 , Victor Songmene4 and Fawzy H. Samuel1

1 Université du Québec à Chicoutimi, Chicoutimi, Québec, G7H 2B1, Canada

2 Department of Mechanical Engineering, Qassim University, Saudi Arabia

3 Department of Mechanical Engineering, King Saud University, Riyadh, Saudi Arabia

4 Department of Mechanical Eng., École de Technologie Supérieure, Québec, Canada

\*Address all correspondence to: fhsamuel@uqac.ca

© 2021 The Author(s). Licensee IntechOpen. 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.

## **References**

[1] K. Kalaiselvan, N. Murugan, S. Parameswaran, Production and characterization of AA6061–B4C stir cast composite, Materials and Design, 32 (2011), pp. 4004-4009.

[2] S. Roy, S. Sharma, S. Sharma, Review on Fabrication of Aluminum 7075+ B4C Composites and its Testing, Int. J. of Advanced Research in Science and Eng., 6(2017), pp. 647-652.

[3] K.E. Öksüz and K.O. Oskay, The Effects of Aging on the Hardness and Wear Behaviour of 2124 Al Alloy/B4C Composites, Proceedings of the 4th International Congress APMAS2014, April 24-27, 2014, Fethiye, Turkey, 127 (2015), pp. 1367-1369.

[4] A. Singla, S. Shandilya, P. Gera, A. Gupta, Process Parameter Optimization using DOE Methodology on Al- MMC to Maximize Mechanical Properties, International Journal of Emerging Science and Engineering (IJESE), ISSN: 2319-6378, Volume-5 Issue-2, January 2018, pp. 4-7.

[5] S. Mohan, E. Kennedy, MACHINABILITY STUDIES OF ALSIB4C DEVELOPED MMC DEVELOPED USING STIR CASTING, Journal of Engineering & Technology, 4(2014), pp. 25-28.

[6] R.U. Vaidya, S.G. Song, A.K. Zurek, Dynamic mechanical response and thermal expansion of ceramic particle reinforced aluminum 6061 matrix composites. Philos Mag A, 70 (1994): pp. 819-836.

[7] Y. Kumar, G. Anil Kumar, J. Satheesh, T. Madhusudhan, A Review On Properties Of Al-B4C Composite Of Different Routes, International Research Journal of Engineering and Technology, 3 (2016), pp. 860-865.

[8] İ. Topcu, Investigation of Wear Behavior of Particle Reinforced AL/B4C Composites under Different Sintering Conditions, TEHNIČKI GLASNIK 14, (2020), pp. 7-14.

[9] B. Manjunatha, H. B. Niranjan, K.G.Satyanarayana, Effect of amount of boron carbide on wear loss of Al-6061 matrix composite by Taguchi technique and Response surface analysis. Materials Science and Engineering, 376(2018), https://doi. org/10.1088/1757-899X/376/1/012071.

[10] K.M. Shorowordi , A.S.M.A. Haseeb J.P. Celis, Tribo-surface characteristics of Al–B4C and Al–SiC composites worn under different contact pressures, Wear, 261 (2006), pp. 634-641.

[11] K.M. Shorowordi, T. Laoui, A.S.M.A. Haseeb, J.P. Celis, L. royen, Microstructure and interface characteristics of B4C, SiC and Al2O3 reinforced Al matrix composites: a comparative study, J.Mater. Process. Technol. 142 (2003), pp. 738-743.

[12] K.M. Shorowordi, A.S.M.A. Haseeb, J.P. Celis, Velocity effects on the wear, friction and tribochemistry of aluminium MMC sliding against phenolic brake pad, Wear 256 (2004)pp. 1176-1181.

[13] G. Straffellini, M. Pellizzari, A. Molinari, Influence of load and temperature on the dry sliding behaviour of Al-based metal-matrixcomposites against friction material, Wear 256 (2004), pp.754-763.

[14] M. Sing, D.P. Mondal, O.P. Modi, A.K. Jha, Two-body abrasive wear behaviour of aluminum alloy–sillimanite particle reinforced composite, Wear, 253 (2002),pp. 357-368.

[15] R.M. Mohanty, K. Balasubramaniam, S.K. Seshadri,

**59**

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

"Boron-Carbide Reinforced Aluminium 1100 matrix composites: Fabrications and properties", Materials Science and Engineering A, 498, (2008), pp. 42-52.

[24] Y. Mazaheri, M. Meratian, R. Emadi, A.R. Najarian, "Comparison of microstructural and mechanical properties of Al–TiC, Al–B4C and Al–TiC–B4C composites prepared by casting techniques" Materials Science & Engineering A, 560 (2013),pp.

[25] F. Toptan, A. Kilicarslan, I. Kerti, The Effect of Ti Addition on the Properties of Al-B4C Interface: A Microstructural Study, Materials Science Forum, 636-637 (2010), pp.

[26] R. Ipek, Adhesive wear behaviour of B4C and SiC reinforced 4147 Al matrix composites (Al/B4C–Al/SiC), J. Mater. Procng. Technol, 162-163 (2005),

[27] F. Bedir, Characteristic properties of Al–Cu–SiCp and Al–Cu–B4Cp composites produced by hot pressing method under nitrogen atmosphere, Mater. & Design 28 (2007) 1238-1244.

[28] M. Aizenshtein, N.Froumin, E. Shapiro-Tsoref, M.P. Dariel, N. Frage, Wetting and interface phenomena in the B4C/(Cu–B–Si) system, Scripta Mater.

[29] J. Jung and S. Kang, Advances in Manufacturing Boron Carbide– Aluminum Composites, J. Am. Ceram.

[30] P. Shen, B. Zou, S. Jin, Q. Jiang, Reaction mechanism in self-propagating high temperature synthesis of TiC-TiB2/Al composites from an Al-Ti-B4C system, Mater. Sci. & Eng. A 454-455

53 (2005), pp. 1231-1235.

Soc., 87(2004), pp.47-54.

(2007), pp. 300-309.

(2016), 330-348.

[31] P.Kumar and R. Parkash, Experimental investigation and optimization of EDM process

parameters for machining of aluminum boron carbide (Al–B4C) composite, Machining Science and Technology, 20

278-287.

192-197.

pp. 71-75.

[16] A. Canakci "Microstructure and abrasive wear behavior of B4C particle reinforced 2014 Al matrix composites", J. Mater. Sci., 46 (2011),pp.

[17] I. Kerti and F. Toptan , Microstructural variations in cast B4C-reinforced aluminum matrix composites (AMCs), Mater Lett, 62,

[18] V. Auradia, G.L. Rajesh, S. A. Kori, Processing of B4C Particulate Reinforced 6061Aluminum Matrix Composites by melt stirring involving two-step addition, Procedia Materials Science, 6 ( 2014 ), pp.1068-1076.

[19] A.R. Kennedy, B. Brampto, The reactive wetting and incorporation of B4C particles into molten aluminum, Scripta Mater, 44, (2014), pp.1077-1082.

[20] H.O. Topcua, N. Gulsoy A.N. Kadiogluc, Gulluoglua, "Processing and mechanical properties of

Alloys Compd. 482 (2009), 516.

[21] I. Topcua, H.O. Gulsoy, N.

[22] K. Kalaiselvan, N.Murugan, S. Parameswaran, Production and characterization of AA6061–B4C stir cast composite, Materials & Design,

[23] M. Uthayakumar, S. Aravindan, K. Rajkumar, Wear performance of Al– SiC–B4C hybrid composites under dry sliding conditions, Materials and Design

32(2011), pp. 4004-4009.

47 (2013), pp. 456-464.

and mechanical properties of B4C reinforced Al matrix composites, Journal of Alloys and Compounds,482(2009), pp. 516-521.

B4Creinforced Al matrix composites", J.

Kadioglu, A.N. Gulluoglu, Processing

2805-2813.

(2008). P.1215.

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

"Boron-Carbide Reinforced Aluminium 1100 matrix composites: Fabrications and properties", Materials Science and Engineering A, 498, (2008), pp. 42-52.

[16] A. Canakci "Microstructure and abrasive wear behavior of B4C particle reinforced 2014 Al matrix composites", J. Mater. Sci., 46 (2011),pp. 2805-2813.

[17] I. Kerti and F. Toptan , Microstructural variations in cast B4C-reinforced aluminum matrix composites (AMCs), Mater Lett, 62, (2008). P.1215.

[18] V. Auradia, G.L. Rajesh, S. A. Kori, Processing of B4C Particulate Reinforced 6061Aluminum Matrix Composites by melt stirring involving two-step addition, Procedia Materials Science, 6 ( 2014 ), pp.1068-1076.

[19] A.R. Kennedy, B. Brampto, The reactive wetting and incorporation of B4C particles into molten aluminum, Scripta Mater, 44, (2014), pp.1077-1082.

[20] H.O. Topcua, N. Gulsoy A.N. Kadiogluc, Gulluoglua, "Processing and mechanical properties of B4Creinforced Al matrix composites", J. Alloys Compd. 482 (2009), 516.

[21] I. Topcua, H.O. Gulsoy, N. Kadioglu, A.N. Gulluoglu, Processing and mechanical properties of B4C reinforced Al matrix composites, Journal of Alloys and Compounds,482(2009), pp. 516-521.

[22] K. Kalaiselvan, N.Murugan, S. Parameswaran, Production and characterization of AA6061–B4C stir cast composite, Materials & Design, 32(2011), pp. 4004-4009.

[23] M. Uthayakumar, S. Aravindan, K. Rajkumar, Wear performance of Al– SiC–B4C hybrid composites under dry sliding conditions, Materials and Design 47 (2013), pp. 456-464.

[24] Y. Mazaheri, M. Meratian, R. Emadi, A.R. Najarian, "Comparison of microstructural and mechanical properties of Al–TiC, Al–B4C and Al–TiC–B4C composites prepared by casting techniques" Materials Science & Engineering A, 560 (2013),pp. 278-287.

[25] F. Toptan, A. Kilicarslan, I. Kerti, The Effect of Ti Addition on the Properties of Al-B4C Interface: A Microstructural Study, Materials Science Forum, 636-637 (2010), pp. 192-197.

[26] R. Ipek, Adhesive wear behaviour of B4C and SiC reinforced 4147 Al matrix composites (Al/B4C–Al/SiC), J. Mater. Procng. Technol, 162-163 (2005), pp. 71-75.

[27] F. Bedir, Characteristic properties of Al–Cu–SiCp and Al–Cu–B4Cp composites produced by hot pressing method under nitrogen atmosphere, Mater. & Design 28 (2007) 1238-1244.

[28] M. Aizenshtein, N.Froumin, E. Shapiro-Tsoref, M.P. Dariel, N. Frage, Wetting and interface phenomena in the B4C/(Cu–B–Si) system, Scripta Mater. 53 (2005), pp. 1231-1235.

[29] J. Jung and S. Kang, Advances in Manufacturing Boron Carbide– Aluminum Composites, J. Am. Ceram. Soc., 87(2004), pp.47-54.

[30] P. Shen, B. Zou, S. Jin, Q. Jiang, Reaction mechanism in self-propagating high temperature synthesis of TiC-TiB2/Al composites from an Al-Ti-B4C system, Mater. Sci. & Eng. A 454-455 (2007), pp. 300-309.

[31] P.Kumar and R. Parkash, Experimental investigation and optimization of EDM process parameters for machining of aluminum boron carbide (Al–B4C) composite, Machining Science and Technology, 20 (2016), 330-348.

**58**

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

Composites under Different Sintering Conditions, TEHNIČKI GLASNIK 14,

[9] B. Manjunatha, H. B. Niranjan, K.G.Satyanarayana, Effect of amount of boron carbide on wear loss of Al-6061 matrix composite by Taguchi technique and Response surface analysis. Materials Science and Engineering, 376(2018), https://doi. org/10.1088/1757-899X/376/1/012071.

[10] K.M. Shorowordi , A.S.M.A. Haseeb J.P. Celis, Tribo-surface characteristics of Al–B4C and Al–SiC composites worn under different

[11] K.M. Shorowordi, T. Laoui, A.S.M.A. Haseeb, J.P. Celis, L. royen, Microstructure and interface characteristics of B4C, SiC and Al2O3 reinforced Al matrix composites: a comparative study, J.Mater. Process. Technol. 142 (2003), pp. 738-743.

[12] K.M. Shorowordi, A.S.M.A. Haseeb, J.P. Celis, Velocity effects on the wear, friction and tribochemistry of aluminium MMC sliding against phenolic brake pad, Wear 256 (2004)pp.

[13] G. Straffellini, M. Pellizzari, A. Molinari, Influence of load and temperature on the dry sliding behaviour of Al-based metal-matrixcomposites against friction material, Wear 256 (2004), pp.754-763.

[14] M. Sing, D.P. Mondal, O.P. Modi, A.K. Jha, Two-body abrasive wear behaviour of aluminum alloy–sillimanite particle reinforced composite, Wear, 253 (2002),pp.

Balasubramaniam, S.K. Seshadri,

contact pressures, Wear, 261 (2006), pp.

(2020), pp. 7-14.

634-641.

1176-1181.

357-368.

[15] R.M. Mohanty, K.

[1] K. Kalaiselvan, N. Murugan, S. Parameswaran, Production and characterization of AA6061–B4C stir cast composite, Materials and Design,

[2] S. Roy, S. Sharma, S. Sharma, Review on Fabrication of Aluminum 7075+ B4C Composites and its Testing, Int. J. of Advanced Research in Science and Eng.,

[3] K.E. Öksüz and K.O. Oskay, The Effects of Aging on the Hardness and Wear Behaviour of 2124 Al Alloy/B4C Composites, Proceedings of the 4th International Congress APMAS2014, April 24-27, 2014, Fethiye, Turkey, 127

[4] A. Singla, S. Shandilya, P. Gera, A. Gupta, Process Parameter Optimization using DOE Methodology on Al- MMC to Maximize Mechanical Properties, International Journal of Emerging Science and Engineering (IJESE), ISSN: 2319-6378, Volume-5 Issue-2, January

DEVELOPED USING STIR CASTING, Journal of Engineering & Technology,

[6] R.U. Vaidya, S.G. Song, A.K. Zurek, Dynamic mechanical response and thermal expansion of ceramic particle reinforced aluminum 6061 matrix composites. Philos Mag A, 70 (1994):

[7] Y. Kumar, G. Anil Kumar, J. Satheesh, T. Madhusudhan, A Review On Properties Of Al-B4C Composite Of Different Routes, International Research Journal of Engineering and Technology,

[8] İ. Topcu, Investigation of Wear Behavior of Particle Reinforced AL/B4C

3 (2016), pp. 860-865.

32 (2011), pp. 4004-4009.

**References**

6(2017), pp. 647-652.

(2015), pp. 1367-1369.

2018, pp. 4-7.

[5] S. Mohan, E. Kennedy, MACHINABILITY STUDIES OF ALSIB4C DEVELOPED MMC

4(2014), pp. 25-28.

pp. 819-836.

[32] Wang HY, Jiang QC, Wang Y, Ma BX, Zhao F. Fabrication of TiB2 particulate reinforced magnesium matrix composites by powder metallurgy. Mater Lett 58(2004), pp.3509-3513.

[33] H. Yang, T. D. Topping, K. Wehage, L. Jiang, E. J. Lavernia, J. M. Schoenung, Tensile behavior and strengthening mechanisms in a submicron B4Creinforced Al trimodal composite, Materials Science and Engineering: A, 616 (2014), pp. 35-43.

[34] A. Yazdani and E. Salahinejad, Evolution of reinforcement distribution in Al–B4C composites during accumulative roll bonding, Materials and Design 32 (2011), pp. 3137-3142.

[35] H. Junaedi, M.F.A. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman, A.A. Almajid, and F.H. Samuel, Effect of testing temperature on the strength and fracture behavior of Al-B4C composites. Journal of Composite Materials, 50(2016): pp. 2871-2880.

[36] M.F. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman and F.H. Samuel, On the impact toughness of Al-15 vol.% B4C metal matrix composites. Composites Part B: Engineering, 2015. 79 (2015): pp. 83-94.

[37] M.F. Ibrahim, A.M. Samuel, F.H. Samuel, H.R. Ammar and M.S. Soliman, On the Impact Toughness of Al-B4C MMC: The Role of Minor Additives and Heat Treatment, in 119th Metalcasting Congress, AFS 2015, Columbus, OH, April 21-23, 2015. 2015, American Foudry Society: Columbus, Ohio.

[38] M.F. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman, A.A. Almajid, and F.H. Samuel, Mechanical properties and fracture of Al-15 vol.-%B4C based metal matrix composites. International Journal of Cast Metals Research, 27(2014): pp. 7-14.

[39] M.F. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman and F.H. Samuel, Neue technologie zur produktion von Al-B4C Metalmatrixverbundwerkstoffen. Giesserei Praxis, 6(2014): pp. 263-271.

[40] M.F. Ibrahim, A.M. Samuel, H.R. Ammar, M.S. Soliman and F.H. Samuel, A new technology for the production of Al-B4C metal matrix composites. Giesserei-Praxis, 65(2014): pp. 263-271.

[41] M.F. Ibrahim, H.R. Ammar, S.A. Alkahtani and F.H. Samuel, Metallographic investigation of tensile- and impact-tested aluminum composites. Journal of Composite Materials, 2016. 50(20): p. 2793-2805.

[42] M.F. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman and F.H. Samuel, Metallurgical parameters controlling matrix/B4C particulate interaction in aluminium-boron carbide metal matrix composites. International Journal of Cast Metals Research, 26(2013): pp. 364-373.

[43] M.F. Ibrahim, A.M. Samuel, H.R. Ammar, M.S. Soliman and F.H. Samuel, A New Technology for the Production of Al-B4C Metal Matrix Composites, in Transactions of the American Foundry Society volume 121; 117th Metalcasting Congress ; April 6-9, 2013, [St. Louis, Missouri ; in conjunction with CastExpo'13]. 2013, American Foundry Society: [St. Louis, Missouri; in conjunction with CastExpo'13]. p. 99-110.

[44] M. S. Soliman, M. M. El Rayes, A. T. Abbas, D. Yu. Pimenov, I. N. Erdakov, H. Junaedi, Effect of tensile strain rate on high-temperature deformation and fracture of rolled Al-15 vol% B4C composite, Materials Science & Engineering A, 749(2019), pp.129-136.

[45] A. J. Pyzik, R. A. Newman, A. Wetzel and E. Dubensky: 'Composition control in aluminum boron carbide

**61**

307-320.

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

> Mechanical Behavior and Applications, S. Sivasankaran, Editor. 2017, InTech.

[53] H. Tahiri, F.H. Samuel and A.M. Samuel, Grain refinement in modified A356 Alloy, in 18th Canadian Materials Science Conference. 2006: McGill University, Montreal, QC,

[54] H. Tahiri, A.M. Samuel, F.H. Samuel, H.W. Doty and S. Valtierra, Mécanismes d'affinage des grains dans les alliages Al-Si en utilisant des alliages mères, in Int. Symposium on Aluminium : From Raw Materials to Applications, 45th annual Conf. of Metallurgists of CIM, C. 2006, Editor. 2006: Montreal, QC, Canada. pp.

[55] K.B. Lee, H.S. Sim, S.Y. Cho Reaction products of Al-Mg/B4C composite fabricated by pressureless infiltration technique. Mater Sci Eng A,

[56] K.E. Knipling, R.A. Karnesky, C.P. Lee, D.C. Dunand, D.N. Seidman, Precipitation evolution in Al-0.1Sc, Al-0.1Zr and Al-0.1Sc-0.1Zr (at. %) alloys during isochronal aging. Acta Mater, 58(2010), pp.5184-5195.

[57] K.E. Knipling, D.C. Dunand, D.N. Seidman, Criteria for developing castable, creep resistant aluminumbased alloys - A review. Z. Metallkunde,

[58] K.E. Knipling , D.N. Seidman , D.C. Dunand, Ambient- and hightemperature Mechanical properties of isochronally aged Al-0.06Sc, Al-0.06Zr and Al-0.06Sc-0.06Zr (at.%) alloys. Acta Mater, 59 (2011);

[59] C.B. Fuller, D.N. Seidman, D.C. Dunand Mechanical properties of Al (Sc, Zr) alloys at ambient and elevated temperatures. Acta Mater, 51(2003):pp.

A302 (2001): pp. 227-234.

97(2006):pp.246-265.

pp. 943-954.

4803-4814.

Canada.

863-880.

composites,' in 'Mechanical properties and performance of engineering ceramics II: ceramic engineering and science proceedings', (ed. R. Tandon et al.), Vol. 27; 2007, Westerville, OH,

[46] X. G. Chen: 'Application of Al–B4C metal matrix composites in the nuclear industry for neutron absorber materials', Proc. On 'Solidification processing of metal matrix composites', San Antonio, TX, USA, March 13-16, 2006 TMS,

[47] M. C. Flemings: 'Fluidity of metals – techniques for producing ultra-thin section castings', Br. Foundryman,

[48] M. Rosso, Ceramic and metal matrix composites: Routes and properties, Journal of Materials Processing Technology, 175(2006),

[49] H. Tahiri: 'Affinement des grains des alliages Al-(0-17%)Si', PhD thesis, Université du Québec à Chicoutimi, Chicoutimi, Que., Canada, 2007.

[50] H. Tahiri, S.S. Mohamed, H.W. Doty, S. Valtierra and F.H. Samuel, Effect of Sr–Grain Refining–Si Interactions on the Microstructural Characteristics of Al–Si Hypoeutectic Alloys. International Journal of Metalcasting, 12(2018): pp. 343-361.

[51] H. Tahiri, A.M. Samuel, H.W. Doty, S. Valtierra and F.H. Samuel, Effect of Sr–Grain Refiner–Si Interactions on the Microstructure Characteristics of Al–Si Hypereutectic Alloys. International Journal of Metalcasting, 12(2018): p.

[52] H. Tahiri, S.S. Mohamed, H.W. Doty, S. Valtierra and F.H. Samuel, Effects of Grain Refining on Columnarto-Equiaxed Transition in Aluminum Alloys, in Aluminum Alloys - Recent Trends in Processing, Characterization,

American Ceramic Society.

pp.343-350.

pp.364-375.

57(1964), pp.312-325.

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

[39] M.F. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman and F.H. Samuel, Neue technologie zur produktion von Al-B4C Metalmatrixverbundwerkstoffen. Giesserei Praxis,

[40] M.F. Ibrahim, A.M. Samuel, H.R. Ammar, M.S. Soliman and F.H. Samuel, A new technology for the production of Al-B4C metal matrix composites. Giesserei-Praxis, 65(2014): pp. 263-271.

[41] M.F. Ibrahim, H.R. Ammar, S.A. Alkahtani and F.H. Samuel, Metallographic investigation of tensile- and impact-tested aluminum composites. Journal of Composite Materials, 2016. 50(20): p. 2793-2805.

[42] M.F. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman and F.H. Samuel, Metallurgical parameters controlling matrix/B4C particulate interaction in aluminium-boron carbide metal matrix composites. International Journal of Cast Metals Research, 26(2013): pp.

[43] M.F. Ibrahim, A.M. Samuel, H.R. Ammar, M.S. Soliman and F.H. Samuel, A New Technology for the Production of Al-B4C Metal Matrix Composites, in Transactions of the American Foundry Society volume 121; 117th Metalcasting Congress ; April 6-9, 2013, [St. Louis, Missouri ; in conjunction with CastExpo'13]. 2013, American Foundry Society: [St. Louis, Missouri; in conjunction with CastExpo'13]. p.

[44] M. S. Soliman, M. M. El Rayes, A. T. Abbas, D. Yu. Pimenov, I. N. Erdakov, H. Junaedi, Effect of tensile strain rate on high-temperature deformation and fracture of rolled Al-15 vol% B4C composite, Materials Science & Engineering A, 749(2019), pp.129-136.

[45] A. J. Pyzik, R. A. Newman, A. Wetzel and E. Dubensky: 'Composition control in aluminum boron carbide

6(2014): pp. 263-271.

364-373.

99-110.

[32] Wang HY, Jiang QC, Wang Y, Ma BX, Zhao F. Fabrication of TiB2 particulate reinforced magnesium matrix composites by powder metallurgy. Mater Lett 58(2004),

[33] H. Yang, T. D. Topping, K. Wehage, L. Jiang, E. J. Lavernia, J. M. Schoenung, Tensile behavior and strengthening mechanisms in a submicron B4Creinforced Al trimodal composite, Materials Science and Engineering: A,

[34] A. Yazdani and E. Salahinejad, Evolution of reinforcement distribution

[35] H. Junaedi, M.F.A. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman, A.A. Almajid, and F.H. Samuel, Effect of testing temperature on the strength and fracture behavior of Al-B4C composites. Journal of Composite Materials, 50(2016): pp. 2871-2880.

[36] M.F. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman and F.H. Samuel, On the impact toughness of Al-15 vol.% B4C metal matrix composites. Composites Part B: Engineering, 2015.

[37] M.F. Ibrahim, A.M. Samuel, F.H. Samuel, H.R. Ammar and M.S. Soliman, On the Impact Toughness of Al-B4C MMC: The Role of Minor Additives and Heat Treatment, in 119th Metalcasting Congress, AFS 2015, Columbus, OH, April 21-23, 2015. 2015, American Foudry Society: Columbus,

[38] M.F. Ibrahim, H.R. Ammar, A.M. Samuel, M.S. Soliman, A.A. Almajid, and F.H. Samuel, Mechanical properties and fracture of Al-15 vol.-%B4C based metal matrix composites. International Journal of Cast Metals Research,

79 (2015): pp. 83-94.

in Al–B4C composites during accumulative roll bonding, Materials and Design 32 (2011), pp. 3137-3142.

pp.3509-3513.

616 (2014), pp. 35-43.

**60**

27(2014): pp. 7-14.

Ohio.

composites,' in 'Mechanical properties and performance of engineering ceramics II: ceramic engineering and science proceedings', (ed. R. Tandon et al.), Vol. 27; 2007, Westerville, OH, American Ceramic Society.

[46] X. G. Chen: 'Application of Al–B4C metal matrix composites in the nuclear industry for neutron absorber materials', Proc. On 'Solidification processing of metal matrix composites', San Antonio, TX, USA, March 13-16, 2006 TMS, pp.343-350.

[47] M. C. Flemings: 'Fluidity of metals – techniques for producing ultra-thin section castings', Br. Foundryman, 57(1964), pp.312-325.

[48] M. Rosso, Ceramic and metal matrix composites: Routes and properties, Journal of Materials Processing Technology, 175(2006), pp.364-375.

[49] H. Tahiri: 'Affinement des grains des alliages Al-(0-17%)Si', PhD thesis, Université du Québec à Chicoutimi, Chicoutimi, Que., Canada, 2007.

[50] H. Tahiri, S.S. Mohamed, H.W. Doty, S. Valtierra and F.H. Samuel, Effect of Sr–Grain Refining–Si Interactions on the Microstructural Characteristics of Al–Si Hypoeutectic Alloys. International Journal of Metalcasting, 12(2018): pp. 343-361.

[51] H. Tahiri, A.M. Samuel, H.W. Doty, S. Valtierra and F.H. Samuel, Effect of Sr–Grain Refiner–Si Interactions on the Microstructure Characteristics of Al–Si Hypereutectic Alloys. International Journal of Metalcasting, 12(2018): p. 307-320.

[52] H. Tahiri, S.S. Mohamed, H.W. Doty, S. Valtierra and F.H. Samuel, Effects of Grain Refining on Columnarto-Equiaxed Transition in Aluminum Alloys, in Aluminum Alloys - Recent Trends in Processing, Characterization, Mechanical Behavior and Applications, S. Sivasankaran, Editor. 2017, InTech.

[53] H. Tahiri, F.H. Samuel and A.M. Samuel, Grain refinement in modified A356 Alloy, in 18th Canadian Materials Science Conference. 2006: McGill University, Montreal, QC, Canada.

[54] H. Tahiri, A.M. Samuel, F.H. Samuel, H.W. Doty and S. Valtierra, Mécanismes d'affinage des grains dans les alliages Al-Si en utilisant des alliages mères, in Int. Symposium on Aluminium : From Raw Materials to Applications, 45th annual Conf. of Metallurgists of CIM, C. 2006, Editor. 2006: Montreal, QC, Canada. pp. 863-880.

[55] K.B. Lee, H.S. Sim, S.Y. Cho Reaction products of Al-Mg/B4C composite fabricated by pressureless infiltration technique. Mater Sci Eng A, A302 (2001): pp. 227-234.

[56] K.E. Knipling, R.A. Karnesky, C.P. Lee, D.C. Dunand, D.N. Seidman, Precipitation evolution in Al-0.1Sc, Al-0.1Zr and Al-0.1Sc-0.1Zr (at. %) alloys during isochronal aging. Acta Mater, 58(2010), pp.5184-5195.

[57] K.E. Knipling, D.C. Dunand, D.N. Seidman, Criteria for developing castable, creep resistant aluminumbased alloys - A review. Z. Metallkunde, 97(2006):pp.246-265.

[58] K.E. Knipling , D.N. Seidman , D.C. Dunand, Ambient- and hightemperature Mechanical properties of isochronally aged Al-0.06Sc, Al-0.06Zr and Al-0.06Sc-0.06Zr (at.%) alloys. Acta Mater, 59 (2011); pp. 943-954.

[59] C.B. Fuller, D.N. Seidman, D.C. Dunand Mechanical properties of Al (Sc, Zr) alloys at ambient and elevated temperatures. Acta Mater, 51(2003):pp. 4803-4814.

[60] A.G.T. Miserez, Fracture and Toughening of High Volume Fraction Ceramic Particle Reinforced Metals. PhD Thesis, École Polytechnique Federale de Lausanne 2003.

[61] H. Zhang, M.W. Chen, K.T. Ramesh, J. Yed , J.M. Schoenung, S.C. Chin, Tensile behavior and dynamic failure of aluminum 6092/B4C composites. Mater. Sci. Eng A, A433(2006):pp.70-82.

[62] X. Kai, Y. Zhao, A. Wang, Hot deformation behavior of in situ nano ZrB2 reinforced 2024Al matrix composites. Compos Sci Technol, 116 (2015):pp. 1-8.

[63] M. Kouzeli and A. Mortensen A. Size dependent strengthening in particle reinforced aluminum. Acta Mater, 50 (2002): 39-51.

[64] M.C. Shankar, J. P. Ka, R. Shettya, Individual and combined effect of reinforcements on stir cast aluminum metal matrix composites-a review. Int J Curr Eng Technol, 3(2013): pp.922-934.

[65] X. Kai, Y. Zhao, A. Wang, Hot deformation behavior of in situ nano ZrB2 reinforced 2024Al matrix composites. Compos Sci Technol, 116(2015):pp. 1-8.

[66] M.E. Kassner, M-T Perez-Prado, Five-power-law creep in single phase metals and alloys. Prog Mater Sci, 45 (2000):pp. 1-102.

[67] M. Abo-Elkhier M.S. Soliman, Superplastic characteristics of finegrained 7475 aluminum alloy. J Mater Eng Perform 15(2006): 76-80.

[68] M.J. McQueen M.E. Kassner Elevated temperature deformation: Hot working amplifies creep. Mater Sci Eng A, 58 (2005), pp. 410-411.

[69] Zhang H, Ramesh KT and Chin ESC. High strain rate response of aluminum 6092/B4C composites. Mater Sci Eng A, A384 (2004): 26-34.

[70] T.M. Lilo, Enhancing ductility of Al6061-10 wt.% B4C through equalchannel angular extrusion processing. Mater Sci Eng A, A410 (2005): pp.43-446.

[71] F.A. Mohamed, Correlation between creep behavior in Al-based solid solution alloys and powder metallurgy Al alloys, Mater. Sci. Eng. A 245 (1998) 242-256.

[72] E. Evangelista, S. Spigarelli, Constitutive equations for creep and plasticity of aluminum alloys produced by powder metallurgy and aluminumbased metal matrix composites, Metall. Mater. Trans. A 33A (2002) 373-381.

[73] S. Spigarelli, E. Evangelista, S. Cucchieri, Analysis of the creep response of an Al–17Si–4Cu–0.55Mg alloy, Mater. Sci. Eng. A 387-389 (2004) 702-705.

[74] F.A. Mohamed, K.T. Park, E.J. Lavernia, Creep Behavior of Discontinuous SiC-Al Composites, Mater. Sci. Eng. A 150 (1992) 21-35.

[75] R. Fernandez, G. Gonzalez-Doncel, Creep behavior of ingot and powder metallurgy 6061Al, J. Alloy Compd. 440 (2007) 158-167.

[76] K.T. Park, F.A. Mohamed, Creep strengthening in a discontinuous SiC-Al composite, Metall. Mater. Trans. A 26 (1995) 3119-3129.

[77] S. Spigarelli, E. Evangelista, H.J. McQueen, Study of hot workability of a heat treated AA6082 aluminum alloy, Scr. Mater. 49 (2003) 179-183.

[78] E.A. El-Danaf, A.A. Almajid, M.S. Soliman, High-temperature deformation and ductility of a modified 5083 Al alloy, J. Mater. Eng. Perform. 17 (2008) 572-579.

**63**

*Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

AA6082-T4 aluminum alloy, J. Mater.

[80] E.A. El-Danaf, M.S. Soliman, A.A. Almajid, Effect of solution heat treatment on the hot workability of Al-Mg-Si Alloy, Mater. Manuf. Process.

[81] S. Spigarelli, C. Paoletti, A new model for the description of creep behavior of aluminum-based composites reinforced with nanosized particles, Compos. Part A:Appl. Sci.

Manuf. 112 (2018) 346-355.

Sci. 43 (2008) 6324-6330.

24 (2009) 637-643.

[79] E.A. El-Danaf, A.A. Almajid, M.S. Soliman, Hot deformation of *Why Al-B4C Metal Matrix Composites? A Review DOI: http://dx.doi.org/10.5772/intechopen.95772*

AA6082-T4 aluminum alloy, J. Mater. Sci. 43 (2008) 6324-6330.

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

pp.43-446.

242-256.

702-705.

(2007) 158-167.

(1995) 3119-3129.

(2008) 572-579.

[70] T.M. Lilo, Enhancing ductility of Al6061-10 wt.% B4C through equalchannel angular extrusion processing. Mater Sci Eng A, A410 (2005):

[71] F.A. Mohamed, Correlation between

creep behavior in Al-based solid solution alloys and powder metallurgy Al alloys, Mater. Sci. Eng. A 245 (1998)

[72] E. Evangelista, S. Spigarelli, Constitutive equations for creep and plasticity of aluminum alloys produced by powder metallurgy and aluminumbased metal matrix composites, Metall. Mater. Trans. A 33A (2002) 373-381.

[73] S. Spigarelli, E. Evangelista, S. Cucchieri, Analysis of the creep response of an Al–17Si–4Cu–0.55Mg alloy, Mater. Sci. Eng. A 387-389 (2004)

[74] F.A. Mohamed, K.T. Park, E.J. Lavernia, Creep Behavior of Discontinuous SiC-Al Composites, Mater. Sci. Eng. A 150 (1992) 21-35.

[75] R. Fernandez, G. Gonzalez-Doncel, Creep behavior of ingot and powder metallurgy 6061Al, J. Alloy Compd. 440

[76] K.T. Park, F.A. Mohamed, Creep strengthening in a discontinuous SiC-Al composite, Metall. Mater. Trans. A 26

[77] S. Spigarelli, E. Evangelista, H.J. McQueen, Study of hot workability of a heat treated AA6082 aluminum alloy,

Scr. Mater. 49 (2003) 179-183.

[78] E.A. El-Danaf, A.A. Almajid, M.S. Soliman, High-temperature deformation and ductility of a modified 5083 Al alloy, J. Mater. Eng. Perform. 17

[79] E.A. El-Danaf, A.A. Almajid, M.S. Soliman, Hot deformation of

[60] A.G.T. Miserez, Fracture and Toughening of High Volume Fraction Ceramic Particle Reinforced Metals. PhD Thesis, École Polytechnique Federale de Lausanne 2003.

[61] H. Zhang, M.W. Chen, K.T. Ramesh, J. Yed , J.M. Schoenung, S.C. Chin, Tensile behavior and dynamic failure of aluminum 6092/B4C composites. Mater. Sci. Eng A, A433(2006):pp.70-82.

[62] X. Kai, Y. Zhao, A. Wang, Hot deformation behavior of in situ nano ZrB2 reinforced 2024Al matrix composites. Compos Sci Technol, 116

[63] M. Kouzeli and A. Mortensen A. Size dependent strengthening in particle reinforced aluminum. Acta Mater, 50

[64] M.C. Shankar, J. P. Ka, R. Shettya, Individual and combined effect of reinforcements on stir cast aluminum metal matrix composites-a review. Int J Curr Eng Technol, 3(2013): pp.922-934.

[65] X. Kai, Y. Zhao, A. Wang, Hot deformation behavior of in situ nano ZrB2 reinforced 2024Al matrix composites. Compos Sci Technol,

[66] M.E. Kassner, M-T Perez-Prado, Five-power-law creep in single phase metals and alloys. Prog Mater Sci, 45

[67] M. Abo-Elkhier M.S. Soliman, Superplastic characteristics of finegrained 7475 aluminum alloy. J Mater

Eng Perform 15(2006): 76-80.

[68] M.J. McQueen M.E. Kassner Elevated temperature deformation: Hot working amplifies creep. Mater Sci

Eng A, 58 (2005), pp. 410-411.

[69] Zhang H, Ramesh KT and

Sci Eng A, A384 (2004): 26-34.

Chin ESC. High strain rate response of aluminum 6092/B4C composites. Mater

(2015):pp. 1-8.

(2002): 39-51.

116(2015):pp. 1-8.

(2000):pp. 1-102.

**62**

[80] E.A. El-Danaf, M.S. Soliman, A.A. Almajid, Effect of solution heat treatment on the hot workability of Al-Mg-Si Alloy, Mater. Manuf. Process. 24 (2009) 637-643.

[81] S. Spigarelli, C. Paoletti, A new model for the description of creep behavior of aluminum-based composites reinforced with nanosized particles, Compos. Part A:Appl. Sci. Manuf. 112 (2018) 346-355.

**Chapter 4**

**Abstract**

*Fawzy H. Samuel*

grain refining, mechanical properties

**1. Introduction**

**65**

Applications of Rare Earth Metals

*Agnes-Marie Samuel, Victor Songmene, Herebert W. Doty and*

The present article reviews a large number of research publications on the effect of mischmetal (MM), rare earth metals (RE), La or Ce, and combinations of La + Ce on the performance of Al-Si cast alloys mainly 319, 356, 380, 413, and 390 alloys. Most of these articles focused on the use of rare earth metals as a substitute for strontium (Sr) as a eutectic silicon (Si) modifier if added in low percentage (< 1 wt.%) to avoid precipitation of a significant amount of insoluble intermetallics and hence poor mechanical properties. Other points that were considered were the affinity of RE to react with Sr., reducing its effectiveness as modifier, as well as the grain refining efficiency of the added RE in any form. None of these articles mentioned the exact composition of the RE used and percentage of tramp elements inherited from the parent ore. Using high purity La or Ce proved to have no effect on the Si shape, size or distribution, in particular at low solidification rates (thick sections). However, regardless the source of the RE, its addition to Sr-modified alloys reduced the modification effect. As for grain refining, apparently a high percentage of RE (> 1 wt.%) is required to achieve a noticeable reduction in grain size, however at the cost of alloy brittleness.

**Keywords:** aluminum alloys, mischmetal, rare earth metals, La, Ce, modification,

hardness (6.5-mohs scale) and low solubility in aluminum matrix which would enhance the alloy wear resistance. The eutectic temperature of Al-11.7% Si alloys is near 577 °C as shown in **Figure 1**, which makes Al-Si alloys easy to cast using different techniques [1–5]. In general, alloys containing Cu and Mg are hardened applying a suitable heat treatable cycle that depends on the thickness of the casting [6–10]. During solidification, the liquid moves along the liquidus line, thus increasing the amount of aluminum, **Figure 2(a)**. At the eutectic temperature, **Figure 2 (b)**, almost 50% of liquid has been solidified. As the temperature continues to decrease (577 °C), the rest of the liquid decomposes into solid Al mixed with solid Si

Soundness of the cast component depends mainly on the ability of the liquid metal to feed interdendritic regions [11]. In the event of low fluidity, shrinkage cavities may occur, as shown in **Figure 3** [12–14], and in the case of automotive

) coupled with its high

Silicon is characterized by its low density ((2.34 g cm<sup>3</sup>

but on a finer scale as presented in **Figure 2(c)**.

in Al-Si Cast Alloys

*Mohamed Gamal Mahmoud, Yasser Zedan,*

## **Chapter 4**

## Applications of Rare Earth Metals in Al-Si Cast Alloys

*Mohamed Gamal Mahmoud, Yasser Zedan, Agnes-Marie Samuel, Victor Songmene, Herebert W. Doty and Fawzy H. Samuel*

### **Abstract**

The present article reviews a large number of research publications on the effect of mischmetal (MM), rare earth metals (RE), La or Ce, and combinations of La + Ce on the performance of Al-Si cast alloys mainly 319, 356, 380, 413, and 390 alloys. Most of these articles focused on the use of rare earth metals as a substitute for strontium (Sr) as a eutectic silicon (Si) modifier if added in low percentage (< 1 wt.%) to avoid precipitation of a significant amount of insoluble intermetallics and hence poor mechanical properties. Other points that were considered were the affinity of RE to react with Sr., reducing its effectiveness as modifier, as well as the grain refining efficiency of the added RE in any form. None of these articles mentioned the exact composition of the RE used and percentage of tramp elements inherited from the parent ore. Using high purity La or Ce proved to have no effect on the Si shape, size or distribution, in particular at low solidification rates (thick sections). However, regardless the source of the RE, its addition to Sr-modified alloys reduced the modification effect. As for grain refining, apparently a high percentage of RE (> 1 wt.%) is required to achieve a noticeable reduction in grain size, however at the cost of alloy brittleness.

**Keywords:** aluminum alloys, mischmetal, rare earth metals, La, Ce, modification, grain refining, mechanical properties

#### **1. Introduction**

Silicon is characterized by its low density ((2.34 g cm<sup>3</sup> ) coupled with its high hardness (6.5-mohs scale) and low solubility in aluminum matrix which would enhance the alloy wear resistance. The eutectic temperature of Al-11.7% Si alloys is near 577 °C as shown in **Figure 1**, which makes Al-Si alloys easy to cast using different techniques [1–5]. In general, alloys containing Cu and Mg are hardened applying a suitable heat treatable cycle that depends on the thickness of the casting [6–10]. During solidification, the liquid moves along the liquidus line, thus increasing the amount of aluminum, **Figure 2(a)**. At the eutectic temperature, **Figure 2 (b)**, almost 50% of liquid has been solidified. As the temperature continues to decrease (577 °C), the rest of the liquid decomposes into solid Al mixed with solid Si but on a finer scale as presented in **Figure 2(c)**.

Soundness of the cast component depends mainly on the ability of the liquid metal to feed interdendritic regions [11]. In the event of low fluidity, shrinkage cavities may occur, as shown in **Figure 3** [12–14], and in the case of automotive

**Figure 1.** *Al-Si binary diagram [1].*

**Figure 2.**

*Progress in liquid to solid during solidification: Start, (b) at eutectic temperature, (c) room temperature [2].*

**Figure 3.**

*(a) Schematic diagram showing movement of liquid metal around the dendritic structure, (b) formation if shrinkage cavities [11].*

components, rendering the casting poor pressures tightness. Another point to consider is the introduction of tangled oxide films (bifilms) as shown in **Figure 4**.

> containing at least 4–5 elements (5–35 at.% concentrations) are considered as HEAs [19]. Research indicates that some HEAs have considerably better strength-toweight ratios, with a higher degree of fracture resistance, tensile strength, as well as corrosion and oxidation resistance than conventional alloys. The research investigations on the Al-Si alloys being reviewed here also focus on achieving the same

**O** Annealed – Applies to product which has been heated to produce the lowest strength condition

**T** Age hardened alloys - To produce stable other than F, O, or H. The "T" is always followed by

**H** Strain Hardened – Applies to products which are strengthened through cold-working.

In this review, the microstructures and mechanical properties of aluminum alloys containing different rare earth element additions are discussed, mainly 319,

characteristics.

**67**

**Figure 4.**

**Table 1.**

**Table 2.**

*Classification of aluminum cast alloys.*

**W** Solution Heat-Treated

*Basic heat treatment designations.*

one or more digits

**F** As fabricated – no heat treatment

**Letter Details**

*Examples of tangled oxide films in aluminum alloys [13].*

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

**Alloy Series Principal Alloying Element 1xx.x** 99.000% minimum Aluminum

**3xx.x** Silicon Plus Copper and/or Magnesium

**2xx.x** Copper

**4xx.x** Silicon **5xx.x** Magnesium **6xx.x** Unused Series

**7xx.x** Zinc **8xx.x** Tin

to improve ductility and dimensional stability

**Table 1** shows the classification of aluminum cast alloys [14] depending on the main alloying element. The number following the decimal point indicates if the alloys are in the form of castings (.0) or ingots (.1 or .2), whereas **Table 2** depicts the different heat treatment designations for these alloys [15, 16]. The composition of Al-Si alloys that are commonly used in aluminum automotive industries is listed in **Table 3** [17].

High entropy alloys (HEAs) are currently receiving much attention in materials engineering because they have potentially desirable properties [18]. Alloys

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

#### **Figure 4.**

*Examples of tangled oxide films in aluminum alloys [13].*


#### **Table 1.**

*Classification of aluminum cast alloys.*


#### **Table 2.**

components, rendering the casting poor pressures tightness. Another point to consider is the introduction of tangled oxide films (bifilms) as shown in **Figure 4**. **Table 1** shows the classification of aluminum cast alloys [14] depending on the main alloying element. The number following the decimal point indicates if the alloys are in the form of castings (.0) or ingots (.1 or .2), whereas **Table 2** depicts the different heat treatment designations for these alloys [15, 16]. The composition of Al-Si alloys that are commonly used in aluminum automotive industries is listed

*(a) Schematic diagram showing movement of liquid metal around the dendritic structure, (b) formation if*

*Progress in liquid to solid during solidification: Start, (b) at eutectic temperature, (c) room temperature [2].*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

High entropy alloys (HEAs) are currently receiving much attention in materials

engineering because they have potentially desirable properties [18]. Alloys

in **Table 3** [17].

**66**

**Figure 1.**

**Figure 2.**

**Figure 3.**

*shrinkage cavities [11].*

*Al-Si binary diagram [1].*

*Basic heat treatment designations.*

containing at least 4–5 elements (5–35 at.% concentrations) are considered as HEAs [19]. Research indicates that some HEAs have considerably better strength-toweight ratios, with a higher degree of fracture resistance, tensile strength, as well as corrosion and oxidation resistance than conventional alloys. The research investigations on the Al-Si alloys being reviewed here also focus on achieving the same characteristics.

In this review, the microstructures and mechanical properties of aluminum alloys containing different rare earth element additions are discussed, mainly 319,


#### **Table 3.**

*Composition of Al-Si based cast alloys commonly used in automotive components [17].*

356, 380, 411, and 390 alloys which constitute alloys widely used in automotive components. A number of scientific investigations are reported in the literature on the effects of rare earth elements and mischmetal (MM), which is a mixture of rare earth (RE) elements found in abundance in nature with Cerium (Ce) and Lanthanum (La) together comprising approximately 90% of mischmetal. The RE metals investigated include Cerium (Ce), Lanthanum (La), Yttrium (Y) [20–25], Erbium (Er) [26–30], Neodymium (Nd) [31–35], Ytterbium (Yb) [36–40], Samarium (Sm) [41–42], Scandium (Sc) [43–47], Europium (Eu) [48, 49], and Gadolinium (Gd) [50]. Among these, the effects of MM, Ce, La and mixed additions of Ce and La are reviewed.

#### **1.1 Effects of Mischmetal (MM) additions**

The effects of small amounts of mischmetal (MM) on the dendrite arm spacing, and the Brinell hardness of Al-1.0% Mg-0.5% Si alloy were investigated by Young et al. [51]. Their results showed that in the range of 0.5–4.0 wt.% MM the hardness increased by more than 30% and the dendrite arm spacing decreased from 50 μm to 18 μm. Ravi et al. [52] analyzed the mechanical properties and microstructure of cast Al-7Si-0.3 Mg (LM25/356) alloy and reported that addition of above 1.0 wt.% of MM lowered the YS, UTS and percent elongation, with an increase in the Brinell hardness. The mechanical properties decreased due to the formation of Ce and La hard intermetallic compounds in the matrix and consumption of a certain amount of Mg in their formation, which reduced the strengthening constituent Mg2Si formed, contributing to the observed decrease. The yield strength (YS), ultimate tensile strength (UTS), and pct elongation of the Al-7Si-0.3 Mg alloy (in the T6 condition) decreased with the increase in Fe content (from 0.2 to 0.6 pct), as shown in **Figure 5**.

revealed that addition of 3.0 wt.% of MM leads to depressions of 12-17 °C in the primary Si reaction temperature and 2–7 °C in the eutectic Si temperature. Increasing the level of MM additions to in situ Al-15Mg2Si composite alloy leads to: (i) a reduction in the size of Mg2Si particles, (ii) a change in the morphology of eutectic Mg2Si from fibrous to flake like, and (iii) formation of RE-containing compounds in

*Scanning electron micrograph of the deep etched Al-21% Si-3% RE alloy. The RE modified primary Si shows*

*Mechanical properties of the LM 25 alloys (T6 condition) containing Fe and Mischmetal [52].*

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

Ravi et al. [57] studied the effect of 1.0 wt.% MM additions on the microstructural characteristics and the room and elevated temperature tensile properties of Al-7Si-0.3 Mg (LM 25/356) alloy containing excess iron up to 0.6 wt.%. The results

i. Alloys with Fe contents ranging from 0.2 to 0.6 wt.%, exhibit grain refinement and partial modification of the eutectic silicon and the finer

the form of Al11RE3 [56].

*typical polyhedral shape [54].*

showed that:

**69**

**Figure 6.**

**Figure 5.**

The microstructure and thermal analysis of Al-21 wt.% Si alloys with MM addition were discussed by Chang et al. [53, 54]. According to the authors, addition of 2.0 wt.% MM leads to a morphological change in the primary Si crystals from starlike to polyhedral shape [55] as displayed in **Figure 6**. The thermal analysis results

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

**Figure 5.** *Mechanical properties of the LM 25 alloys (T6 condition) containing Fe and Mischmetal [52].*

#### **Figure 6.**

356, 380, 411, and 390 alloys which constitute alloys widely used in automotive components. A number of scientific investigations are reported in the literature on the effects of rare earth elements and mischmetal (MM), which is a mixture of rare earth (RE) elements found in abundance in nature with Cerium (Ce) and Lanthanum (La) together comprising approximately 90% of mischmetal. The RE metals investigated include Cerium (Ce), Lanthanum (La), Yttrium (Y) [20–25], Erbium (Er) [26–30], Neodymium (Nd) [31–35], Ytterbium (Yb) [36–40], Samarium (Sm) [41–42], Scandium (Sc) [43–47], Europium (Eu) [48, 49], and Gadolinium (Gd) [50]. Among these, the effects of MM, Ce, La and mixed additions of Ce and La are

**Method(b) Si Cu Mg Fe Zn Others**

**Alloy Elements (wt. %)**

319.0 S, P 6.0 3.5 <0.10 <1.0 <1.0 332.0 P 9.5 3.0 1.0 1.2 1.0 355.0 S, P 5.0 1.25 0.5 <0.06 <0.35 A356.0 S, P 7.0 <0.20 0.35 <0.2 <0.1

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

380.0 D 8.5 3.5 <0.1 <1.3 <3.0

413.0 D 12.0 <0.1 <0.10 <2.0 — 443.0 S, P - 5.25 <0.3 <0.05 <0.8 <0.5

*Composition of Al-Si based cast alloys commonly used in automotive components [17].*

A357.0 S, P 7.0 <0.20 0.55 <0.2 <0.1 0.05 Be

383.0 D 10.0 2.5 0.10 1.3 3.0 0.15 Sn 384.0 D 11.0 2.0 <0.3 <1.3 <3.0 0.35 Sn 390.0 D 17.0 4.5 0.55 <1.3 <0.1 <0.1 Mg

The effects of small amounts of mischmetal (MM) on the dendrite arm spacing, and the Brinell hardness of Al-1.0% Mg-0.5% Si alloy were investigated by Young et al. [51]. Their results showed that in the range of 0.5–4.0 wt.% MM the hardness increased by more than 30% and the dendrite arm spacing decreased from 50 μm to 18 μm. Ravi et al. [52] analyzed the mechanical properties and microstructure of cast Al-7Si-0.3 Mg (LM25/356) alloy and reported that addition of above 1.0 wt.% of MM lowered the YS, UTS and percent elongation, with an increase in the Brinell hardness. The mechanical properties decreased due to the formation of Ce and La hard intermetallic compounds in the matrix and consumption of a certain amount of Mg in their formation, which reduced the strengthening constituent Mg2Si formed, contributing to the observed decrease. The yield strength (YS), ultimate tensile strength (UTS), and pct elongation of the Al-7Si-0.3 Mg alloy (in the T6 condition) decreased with the increase in Fe content (from 0.2 to 0.6 pct), as shown

The microstructure and thermal analysis of Al-21 wt.% Si alloys with MM addition were discussed by Chang et al. [53, 54]. According to the authors, addition of 2.0 wt.% MM leads to a morphological change in the primary Si crystals from starlike to polyhedral shape [55] as displayed in **Figure 6**. The thermal analysis results

reviewed.

**Table 3.**

in **Figure 5**.

**68**

**1.1 Effects of Mischmetal (MM) additions**

*S = sand casting, P = permanent mold casting, D = die casting.*

*Scanning electron micrograph of the deep etched Al-21% Si-3% RE alloy. The RE modified primary Si shows typical polyhedral shape [54].*

revealed that addition of 3.0 wt.% of MM leads to depressions of 12-17 °C in the primary Si reaction temperature and 2–7 °C in the eutectic Si temperature. Increasing the level of MM additions to in situ Al-15Mg2Si composite alloy leads to: (i) a reduction in the size of Mg2Si particles, (ii) a change in the morphology of eutectic Mg2Si from fibrous to flake like, and (iii) formation of RE-containing compounds in the form of Al11RE3 [56].

Ravi et al. [57] studied the effect of 1.0 wt.% MM additions on the microstructural characteristics and the room and elevated temperature tensile properties of Al-7Si-0.3 Mg (LM 25/356) alloy containing excess iron up to 0.6 wt.%. The results showed that:

i. Alloys with Fe contents ranging from 0.2 to 0.6 wt.%, exhibit grain refinement and partial modification of the eutectic silicon and the finer intermetallic compounds formed with Ce, La, and other elements, thereby improving the strength as well as the ductility of the alloy relative to the same alloys without MM addition.

ii. In the case of non-modified alloys MM addition resulted in partially

iii. The effect of MM as a modifier is more effective at high cooling rate (corresponding to DAS 40 μm) than at the low cooling rate (DAS 120 μm) for all the as-cast non-modified alloys.

A413.1 and A319.1 alloys, compared to A356.2 alloy.

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

*1.1.1 MM-modified as-cast A356 alloys*

for the primary α-Al phase.

**Figure 8.**

**71**

modified eutectic silicon particles. This effect was more pronounced in the

iv. MM-containing intermetallic phases were observed at high and low cooling rates, each exhibiting a specific Ce/La ratio and morphology. Many of these MM-containing intermetallic phases were found to contain Sr., which confirmed the interaction of MM with Sr. – see **Figure 9** and Appendix A.

Combined addition of MM and Mn is an effective way to improve the strength of A390 alloy at elevated temperature by 25% [65]. Other studies by Zhu et al. [66, 67] reported the effects of 0.1–1.0 wt.% Ce-based MM additions on the microstructure, tensile properties and fracture behavior of as-cast and T6-treated A356 alloys. The main findings from their work are listed below for the specified conditions.

MM-containing intermetallic compounds cannot act as potential nucleate sites

i. The modification effect of MM depends on the addition level. Minor additions of MM (less than 0.2 wt.%) result in partial modification, while

ii. The fracture path goes through the interdendritic region composed of eutectic silicon and MM-containing intermetallic compounds.

*Effect of mischmetal additions and aging temperature on the hardness of A319.1 alloys solidified at high*

*solidification rates (DAS 40 μm of Sr-modified conditions [63].*

more than 0.3 wt.% MM leads to full modification.

ii. Alloys containing 0.6 wt.% Fe led to the formation of fine and fibrous shaped intermetallic compounds containing Fe and Si, which reduced the effective amount of Fe available for formation of β and π phases, thereby reducing the size and volume of Fe-containing intermetallics, which, in turn, reduced the deleterious effect of Fe and improved the alloy strength and ductility.

Wan et al. [58] found that up to1.0 wt%, MM addition refined the grain size of cast Al-Cu-Mg-Si alloy and changed the eutectic Si morphology from needle-like and laminar to a granular type. Also, with the increase of the MM level, the tensile strength and elongation of the alloy first increased and then began to decrease. Alloy with 0.7 wt.% MM exhibited the highest mechanical properties. In another study, Chong et al. [59] studied the combined effects of P and MM on the microstructure and mechanical properties of hypereutectic Al-20%Si alloy. It was observed that, in general, alloys with the addition of 0.08% P and 0.6% MM exhibited highest mechanical properties and had the optimal microstructure compared to the alloy with no addition; refinement of the primary Si particles from 66.4 μm to 23.3 μm, and the eutectic silicon from 8.3 μm to 5.2 μm, was also noted. With respect to the mechanical properties, the ultimate tensile strength improved from 256 MPa to 306 MPa, and the ductility increased from 0.35% to 0.48%. **Figure 7** shows the average primary Si size of the tested alloys with different P contents.

According to El Sebaie et al. [60–64] the presence of MM in unmodified and Sr-modified A319.1, A356.2 and A413.1 Al-Si casting alloys led to the following observations:

i. In general, the hardness values of the as-cast alloys were higher at high cooling rates than at low cooling rates. With MM, the hardness decreased at both solidification rates. **Figure 8** shows the hardness values obtained for these alloys after different aging treatments.

**Figure 7.** *Effect of P content on primary Si size of Al-20%Si alloys under same RE content condition [59].*

intermetallic compounds formed with Ce, La, and other elements, thereby improving the strength as well as the ductility of the alloy relative to the

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

ii. Alloys containing 0.6 wt.% Fe led to the formation of fine and fibrous shaped intermetallic compounds containing Fe and Si, which reduced the effective amount of Fe available for formation of β and π phases, thereby reducing the size and volume of Fe-containing intermetallics, which, in turn, reduced the deleterious effect of Fe and improved the alloy strength and ductility.

Wan et al. [58] found that up to1.0 wt%, MM addition refined the grain size of cast Al-Cu-Mg-Si alloy and changed the eutectic Si morphology from needle-like and laminar to a granular type. Also, with the increase of the MM level, the tensile strength and elongation of the alloy first increased and then began to decrease. Alloy with 0.7 wt.% MM exhibited the highest mechanical properties. In another study, Chong et al. [59] studied the combined effects of P and MM on the microstructure and mechanical properties of hypereutectic Al-20%Si alloy. It was observed that, in general, alloys with the addition of 0.08% P and 0.6% MM exhibited highest

mechanical properties and had the optimal microstructure compared to the alloy with no addition; refinement of the primary Si particles from 66.4 μm to 23.3 μm, and the eutectic silicon from 8.3 μm to 5.2 μm, was also noted. With respect to the mechanical properties, the ultimate tensile strength improved from 256 MPa to 306 MPa, and the ductility increased from 0.35% to 0.48%. **Figure 7** shows the average primary Si size

According to El Sebaie et al. [60–64] the presence of MM in unmodified and Sr-modified A319.1, A356.2 and A413.1 Al-Si casting alloys led to the following

i. In general, the hardness values of the as-cast alloys were higher at high cooling rates than at low cooling rates. With MM, the hardness decreased at both solidification rates. **Figure 8** shows the hardness values obtained for

same alloys without MM addition.

of the tested alloys with different P contents.

these alloys after different aging treatments.

*Effect of P content on primary Si size of Al-20%Si alloys under same RE content condition [59].*

observations:

**Figure 7.**

**70**


Combined addition of MM and Mn is an effective way to improve the strength of A390 alloy at elevated temperature by 25% [65]. Other studies by Zhu et al. [66, 67] reported the effects of 0.1–1.0 wt.% Ce-based MM additions on the microstructure, tensile properties and fracture behavior of as-cast and T6-treated A356 alloys. The main findings from their work are listed below for the specified conditions.

### *1.1.1 MM-modified as-cast A356 alloys*

MM-containing intermetallic compounds cannot act as potential nucleate sites for the primary α-Al phase.


### **Figure 8.**

*Effect of mischmetal additions and aging temperature on the hardness of A319.1 alloys solidified at high solidification rates (DAS 40 μm of Sr-modified conditions [63].*

**Figure 9.**

*Backscattered images of A319.1 alloy samples containing (a, b) 0 wt-% and (c, d) 6 wt-%MM depicting intermetallic phases observed under high cooling rate conditions (S: Sr. modified; T: Heat treated samples) [64].*

### *1.1.2 MM-modified T6-A356 alloys*

T6 heat treatment has great influence on the spheroidization of eutectic silicon particles in MM-modified A356 alloys than that in the unmodified alloys.


The effects of different addition levels (0.0–1.0 wt.%) of La-based MM and heat treatment on the microstructure and tensile properties of two different sections of Al-Si casting alloy A357 were studied by Mousavi et al. [68–70]. Optimum recommended levels of MM are 0.1 wt.% and 0.3 wt.% for thin and thick sections of the casting, respectively. Examination of the microstructure at high level of MM (0.5 wt.%) exhibited the precipitation of a new AlSiLa intermetallic phase as shown in **Figure 11**.

A356 alloy wheel refined by the AlTiB-MM were improved significantly. The tensile strength, yield strength, and elongation of the wheel spokes improved by approxi-

*(a) SEM photograph of Al–Si–La compound intermetallic in A357 alloy with 0.5 wt% mischmetal casting in thick section mold and (b) EDS spectrum showing the distribution of Al, Si and La in the intermetallic [70].*

*Fracture surface parallel to the tensile direction of the T6-A356 alloys modified by 0.5 wt.% MM showing*

*intercrystalline crack of RE-containing intermetallic compounds at the fracture surface [67].*

Dang et al. [74] investigated the effects of the use of rare earth (RE) in the form

of Al-10% RE master alloy (a mixture of Ce and La) and pouring temperature (1124 K through 1524 K in increments of 100 K) on the microstructure and mechanical properties of T6-treated Al-25%Si alloy. The authors observed that for the unmodified alloy, the primary Si morphology was transformed from platelets to fine polyhedral form, and the average size decreased with increase in pouring temperature (from 125 μm at 1124 K to 62 μm at 1524 K). With a 1.2 wt.% RE addition, the primary Si exhibited a small blocky morphology, with an average particle size of 47 μm. In addition, the study showed that T6 MM-Al-25%Si alloy

exhibited an improvement in the mechanical properties compared to the

unmodified alloy, where the maximum tensile strength and elongation (208.3 MPa and 1.01%) were obtained for the sample modified with 1.2 wt.% RE followed by

mately 11.3%, 10.8% and 44.1%, respectively.

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

**Figure 10.**

**Figure 11.**

**73**

The results of Jiang et al. [71] on the microstructure, tensile properties and fracture behavior of as-cast and T6 A357 alloy revealed that addition of MM reduced the size of the primary α-Al dendrites i.e., the SDAS value, and also improved the eutectic Si particle morphology. Accordingly, the MM-modified A357 alloy exhibited improvement in the tensile properties, particularly the elongation, in the T6-treated condition. The fracture surface of the tensile-tested sample of the unmodified alloy showed a clear brittle fracture, whereas that of the MM-modified A357 alloy exhibited dimple rupture and cracked eutectic Si particles, resulting in superior ductility. The results of Zhang et al. [72, 73] demonstrated that the AlTiB-MM addition to A356 alloy provided the most effective and synergetic grain size refinement compared to individual AlTiB or MM additions. Also, the properties of

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

#### **Figure 10.**

*1.1.2 MM-modified T6-A356 alloys*

**Figure 9.**

in **Figure 11**.

**72**

T6 heat treatment has great influence on the spheroidization of eutectic silicon

*Backscattered images of A319.1 alloy samples containing (a, b) 0 wt-% and (c, d) 6 wt-%MM depicting intermetallic phases observed under high cooling rate conditions (S: Sr. modified; T: Heat treated samples) [64].*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

i. The UTS, YS, and %EL of the T6-treated A356 alloys with and without modification are improved due to the spheroidization of eutectic silicon

ii. SEM images show that the MM-modified T6 A356 alloy undergoes ductile fracture. It is worth noting that the eutectic silicon and MM-containing intermetallic particles provide the weak locations for initiation of the

The effects of different addition levels (0.0–1.0 wt.%) of La-based MM and heat treatment on the microstructure and tensile properties of two different sections of

recommended levels of MM are 0.1 wt.% and 0.3 wt.% for thin and thick sections of the casting, respectively. Examination of the microstructure at high level of MM (0.5 wt.%) exhibited the precipitation of a new AlSiLa intermetallic phase as shown

The results of Jiang et al. [71] on the microstructure, tensile properties and fracture behavior of as-cast and T6 A357 alloy revealed that addition of MM reduced the size of the primary α-Al dendrites i.e., the SDAS value, and also improved the eutectic Si particle morphology. Accordingly, the MM-modified A357 alloy

exhibited improvement in the tensile properties, particularly the elongation, in the T6-treated condition. The fracture surface of the tensile-tested sample of the unmodified alloy showed a clear brittle fracture, whereas that of the MM-modified A357 alloy exhibited dimple rupture and cracked eutectic Si particles, resulting in superior ductility. The results of Zhang et al. [72, 73] demonstrated that the AlTiB-MM addition to A356 alloy provided the most effective and synergetic grain size refinement compared to individual AlTiB or MM additions. Also, the properties of

Al-Si casting alloy A357 were studied by Mousavi et al. [68–70]. Optimum

particles in MM-modified A356 alloys than that in the unmodified alloys.

particles and Mg2Si precipitation hardening.

fracture as displayed in **Figure 10**.

*Fracture surface parallel to the tensile direction of the T6-A356 alloys modified by 0.5 wt.% MM showing intercrystalline crack of RE-containing intermetallic compounds at the fracture surface [67].*

**Figure 11.**

*(a) SEM photograph of Al–Si–La compound intermetallic in A357 alloy with 0.5 wt% mischmetal casting in thick section mold and (b) EDS spectrum showing the distribution of Al, Si and La in the intermetallic [70].*

A356 alloy wheel refined by the AlTiB-MM were improved significantly. The tensile strength, yield strength, and elongation of the wheel spokes improved by approximately 11.3%, 10.8% and 44.1%, respectively.

Dang et al. [74] investigated the effects of the use of rare earth (RE) in the form of Al-10% RE master alloy (a mixture of Ce and La) and pouring temperature (1124 K through 1524 K in increments of 100 K) on the microstructure and mechanical properties of T6-treated Al-25%Si alloy. The authors observed that for the unmodified alloy, the primary Si morphology was transformed from platelets to fine polyhedral form, and the average size decreased with increase in pouring temperature (from 125 μm at 1124 K to 62 μm at 1524 K). With a 1.2 wt.% RE addition, the primary Si exhibited a small blocky morphology, with an average particle size of 47 μm. In addition, the study showed that T6 MM-Al-25%Si alloy exhibited an improvement in the mechanical properties compared to the unmodified alloy, where the maximum tensile strength and elongation (208.3 MPa and 1.01%) were obtained for the sample modified with 1.2 wt.% RE followed by

## *Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

the T6 heat treatment. Tensile fracture exhibited three stages: microcrack initiation, crack coalescence, and quick crack propagation as shown in **Figure 12**.

**1.2 Effects of cerium additions**

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

strengthening phase.

itation enthalpy.

**75**

The results arising from the investigations of Song et al. [77, 78] showed that individual addition of 0.3%Ce or 0.2%Ti to Al-Cu-Mg-Ag alloys can decrease the grain size of the as-cast alloy, increase the nucleation rate of the Ω (metastable Al2Cu) phase, inhibit the growth of the Ω phase during aging, and thereby increase the volume fraction and decrease the spacing of the Ω phase. Based on these microstructural observations, the yield strength and tensile strength of the alloy are increased. However, combined addition of Ce and Ti led to the formation of (Ce, Ti)-containing intermetallic compounds and increased the grain size during casting, with no influence on the nucleation and growth of the Ω phase during aging. The alloy containing both Ce and Ti had a relatively lower Vickers hardness and strength compared to the alloy containing individual additions of Ce or Ti. In another research, Song et al. [79] reported that Ce improved the thermal stability of the Ω phase by decreasing the diffusion velocity of the Cu atoms, and hence decreasing the coarsening speed of the phase, as well as through the aggregation of Ce atoms at the Ω phase/Al matrix interface, increasing the energy barrier for the thickening of the Ω phase plates which coarsen through a ledge nucleation mechanism. The

The results on the effects of different levels of Ce additions on the microstructure, thermal behavior and mechanical properties of hypereutectic AlSi17CuMg alloy illustrated that addition of Ce (up to 1.0 wt.% Ce) can achieve refinement of the primary and eutectic silicon morphology. In general, alloy containing 1.0 wt.% Ce exhibited the best results with respect to the microstructural and strength properties. It was also observed that with 1.0 wt.% Ce, the alloy produced the highest

The eutectic silicon modification in A356 alloy modified with 1.0 wt.% Ce was greatly improved [80]. However, the thermal analysis revealed that there is no direct relation between the eutectic growth temperature and silicon modification. The microstructural characterization showed that two kinds of Ce-containing intermetallic phases were found, including Ce-23Al-22Si and Al-17Ce-12Ti-2Si-2 Mg (in wt.%). While the ductility of the Ce-modified alloys was enhanced for Ce additions of 0.6 wt.% and above, there was no positive effect on the ultimate tensile strength; this was attributed to the formation of the Al-17Ce-12Ti-2Si-2 Mg phase which reduced the amount of free Mg available for precipitation of the Mg2Si

The effects of various concentrations of Ce (0.0, 0.01, 0.02, 0.05 and 0.1 wt.%) on the solidification and mechanical properties of AA A360 (Al-10%Si-0.5%Mg) alloy were reported by Voncina et al. [81]. The results showed that the solidus temperature decreased with increasing Ce addition. The eutectic (αAl + Mg2Si) temperature also decreased with Ce addition. It was found that the precipitation enthalpy decreased with the Ce addition, while precipitation occurred more rapidly and intensively, indicating increased reaction kinetics. Mechanical properties like tensile strength and hardness also increased with the Ce addition, where the hardness of the investigated alloy could be attributed to the phase composed of Al, Ce, Mg and Si. The precipitation enthalpy also decreased with increasing Ce addition and increased with increasing cooling rate as determined from simple DSC analysis, as shown in **Figure 14**. It was anticipated that Ce in A360 alloy decreases the activation energy for the precipitation of the Mg2Si phase and consequently precip-

It is worth noting that in another study Voncina et al. [82] reported that addition of Ce to A380 alloy led to a change in the morphology of eutectic Al2Cu phase from "crumbled" to fully formed (finer eutectic-like to blocky form) and caused the

strength of the Al-Cu-Mg-Ag alloy is improved, as a result.

reduction in the liquidus temperature from 686.6 °C to 591.9 °C. [79].

According to Mahmoud et al. [75, 76], depending upon the amount of added Ti, two RE-based intermetallics can be formed: (i) a white phase, mainly platelet-like (approximately 2.5 μm thick), that is rich in RE, Si, Cu, and Al, and (ii) a second phase made up of mainly gray sludge particles (star-like) branching in different directions. The gray phase is rich in Ti with some RE (almost 20% of that in the white phase) with traces of Si and Cu. There is a strong interaction between RE and Sr., leading to a reduction in the efficiency of Sr. as a eutectic Si modifier, causing particle demodification. **Figure 13(a)** shows the actual morphology of the white phase which is likely to be thin platelets about 2.5 μm thick. The morphology of the gray sludge is well illustrated in **Figure 13(b)** exhibiting the branching of the gray phase particles in different directions.

**Figure 12.**

*The fracture morphology of Al-25% Si alloy in: (a) and (b) 1.2% RE, (c) and (d) 1.2% RE + T6 heat treatment [74].*

#### **Figure 13.**

*RE-based intermetallic phases: (a) platelet-like phase; the two straight arrows pointing towards each other simply highlight the thickness of the platelet-like phase and the curved arrow that links the label 2.5 μm to the platelet indicates the actual thickness, (b) the gray sludge phase [76].*

## **1.2 Effects of cerium additions**

the T6 heat treatment. Tensile fracture exhibited three stages: microcrack initiation,

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

*The fracture morphology of Al-25% Si alloy in: (a) and (b) 1.2% RE, (c) and (d) 1.2% RE + T6 heat*

*RE-based intermetallic phases: (a) platelet-like phase; the two straight arrows pointing towards each other simply highlight the thickness of the platelet-like phase and the curved arrow that links the label 2.5 μm to the*

*platelet indicates the actual thickness, (b) the gray sludge phase [76].*

According to Mahmoud et al. [75, 76], depending upon the amount of added Ti, two RE-based intermetallics can be formed: (i) a white phase, mainly platelet-like (approximately 2.5 μm thick), that is rich in RE, Si, Cu, and Al, and (ii) a second phase made up of mainly gray sludge particles (star-like) branching in different directions. The gray phase is rich in Ti with some RE (almost 20% of that in the white phase) with traces of Si and Cu. There is a strong interaction between RE and Sr., leading to a reduction in the efficiency of Sr. as a eutectic Si modifier, causing particle demodification. **Figure 13(a)** shows the actual morphology of the white phase which is likely to be thin platelets about 2.5 μm thick. The morphology of the gray sludge is well illustrated in **Figure 13(b)** exhibiting the branching of the gray

crack coalescence, and quick crack propagation as shown in **Figure 12**.

phase particles in different directions.

**Figure 12.**

**Figure 13.**

**74**

*treatment [74].*

The results arising from the investigations of Song et al. [77, 78] showed that individual addition of 0.3%Ce or 0.2%Ti to Al-Cu-Mg-Ag alloys can decrease the grain size of the as-cast alloy, increase the nucleation rate of the Ω (metastable Al2Cu) phase, inhibit the growth of the Ω phase during aging, and thereby increase the volume fraction and decrease the spacing of the Ω phase. Based on these microstructural observations, the yield strength and tensile strength of the alloy are increased. However, combined addition of Ce and Ti led to the formation of (Ce, Ti)-containing intermetallic compounds and increased the grain size during casting, with no influence on the nucleation and growth of the Ω phase during aging. The alloy containing both Ce and Ti had a relatively lower Vickers hardness and strength compared to the alloy containing individual additions of Ce or Ti. In another research, Song et al. [79] reported that Ce improved the thermal stability of the Ω phase by decreasing the diffusion velocity of the Cu atoms, and hence decreasing the coarsening speed of the phase, as well as through the aggregation of Ce atoms at the Ω phase/Al matrix interface, increasing the energy barrier for the thickening of the Ω phase plates which coarsen through a ledge nucleation mechanism. The strength of the Al-Cu-Mg-Ag alloy is improved, as a result.

The results on the effects of different levels of Ce additions on the microstructure, thermal behavior and mechanical properties of hypereutectic AlSi17CuMg alloy illustrated that addition of Ce (up to 1.0 wt.% Ce) can achieve refinement of the primary and eutectic silicon morphology. In general, alloy containing 1.0 wt.% Ce exhibited the best results with respect to the microstructural and strength properties. It was also observed that with 1.0 wt.% Ce, the alloy produced the highest reduction in the liquidus temperature from 686.6 °C to 591.9 °C. [79].

The eutectic silicon modification in A356 alloy modified with 1.0 wt.% Ce was greatly improved [80]. However, the thermal analysis revealed that there is no direct relation between the eutectic growth temperature and silicon modification. The microstructural characterization showed that two kinds of Ce-containing intermetallic phases were found, including Ce-23Al-22Si and Al-17Ce-12Ti-2Si-2 Mg (in wt.%). While the ductility of the Ce-modified alloys was enhanced for Ce additions of 0.6 wt.% and above, there was no positive effect on the ultimate tensile strength; this was attributed to the formation of the Al-17Ce-12Ti-2Si-2 Mg phase which reduced the amount of free Mg available for precipitation of the Mg2Si strengthening phase.

The effects of various concentrations of Ce (0.0, 0.01, 0.02, 0.05 and 0.1 wt.%) on the solidification and mechanical properties of AA A360 (Al-10%Si-0.5%Mg) alloy were reported by Voncina et al. [81]. The results showed that the solidus temperature decreased with increasing Ce addition. The eutectic (αAl + Mg2Si) temperature also decreased with Ce addition. It was found that the precipitation enthalpy decreased with the Ce addition, while precipitation occurred more rapidly and intensively, indicating increased reaction kinetics. Mechanical properties like tensile strength and hardness also increased with the Ce addition, where the hardness of the investigated alloy could be attributed to the phase composed of Al, Ce, Mg and Si. The precipitation enthalpy also decreased with increasing Ce addition and increased with increasing cooling rate as determined from simple DSC analysis, as shown in **Figure 14**. It was anticipated that Ce in A360 alloy decreases the activation energy for the precipitation of the Mg2Si phase and consequently precipitation enthalpy.

It is worth noting that in another study Voncina et al. [82] reported that addition of Ce to A380 alloy led to a change in the morphology of eutectic Al2Cu phase from "crumbled" to fully formed (finer eutectic-like to blocky form) and caused the

structure [84]. Accordingly, the ultimate tensile strength (UTS) and elongation (%

Ye et al. [85] investigated the influence of Ce content (0.2% and 0.4 wt.%) on the impact properties and microstructures of 2519A aluminum alloy, a new version of 2519 alloy (with a higher Cu/Mg ratio) used for armor material. Based on the results of their research, it was found that 0.2 wt.% Ce addition leads to an increase in the volume fraction of the precipitation phase, in addition to a more dispersive and homogeneous distribution of finer θ' (Al2Cu) precipitates, which result in improving the ability of the alloy for absorbing impact energy. Formation of Al8Cu4Ce phase which is thermally stable at high temperature is also expected to enhance the high temperature mechanical properties of the alloy. Yii et al. [86] reported that the addition of Ce in the Al-20%Si alloys refined the Si primary phase as the Ce additions were increased. The results also showed that addition of Ce in the range of 0.46 to 2.24 wt.% led to the formation of fine cells dispersed in the Almatrix. These cells consisted of a mixture of eutectic Si particles and Ce-containing intermetallic phases (Al3Ce and CeAl1.2 Si10.8). The amount of the intermetallic

Promising scientific investigations were made by Ahmad and coworkers [87, 88]

i. The addition of Ce to ADC12 alloy leads to improvement in the Si particles

ii. Cooling rate has no significant effect in the 1.5 wt.% Ce-modified ADC12 alloy, compared to the base alloy and this may be attributed to the

iii. Investigation of the Al-Si eutectic phase using the thermal analysis technique showed that addition of Ce had a significant effect on the nucleation, growth, and minimum temperatures of Al-Si, and decreased as the Ce concentration increased; refinement of the Si structure was observed up to 1.0 wt.% Ce. In addition, the growth and nucleation temperatures of the Al–Cu phase, which is the last phase to solidify, also increased with increasing level of Ce. The formation of Ce-containing intermetallic compounds such as Al-Si-Ce and Al-Si-Cu-Ce affected the degree of Si

iv. Ce addition refined the secondary dendrite arm spacing (SDAS) by

approximately 36%. In addition, the tensile strength and quality index of Al-11%Si-Cu-Mg increased to 237.6 MPa and 265 MPa, respectively, after

Effect of solidification rate and rare earth metal addition on the microstructural characteristics and porosity formation in A356 alloy was investigated by Mahmoud et al. [89]. According to the atomic radius ratio, γLa/γSi is 1.604 and γCe/γSi is 1.559, theoretically, which shows that Ce is relatively more effective than La. These findings confirm that Sr. is the most dominating modification agent. Interaction between rare earth (RE) metals and Sr. would reduce the effectiveness of Sr. Although modification with Sr. causes the formation of shrinkage porosity, it also reacts with RE-rich intermetallics, resulting in their fragmentation. **Figure 17** reveals the distribution of La, Ce, and Sr. in RE-rich platelets, which explains the

on the influence of Ce on the microstructure of a commercial Al-11%Si-Cu-Mg eutectic cast alloy (ADC12). The main findings from their studies are summarized

modification and reduces the Si particle size by 62% [87].

El) increased by 68.2% and 53.1%, respectively, as a result of these effects.

phases increased with increasing Ce addition.

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

formation of intermetallic phases.

modification [88] - **Figure 16** and **Table 4**.

the addition of 0.1 wt.% Ce.

below.

**77**

**Figure 14.**

*Precipitation enthalpy regarding the Ce addition and cooling rate at heating DSC curves.*

**Figure 15.** *Al-corner of ternary Al-Si- Ce system [82].*

formation of small primary crystals of αAl which resulted in grain refining of the alloy. A needle-shaped AlCeCuSi phase (Al9Ce2Cu5Si3) was also detected. The solidification of hypoeutectic Al-Si alloys with Ce addition can be described by the Al-corner of ternary system Al-Si-Ce (**Figure 15**).

Effects of Ce-Sr interaction on the nucleation of primary α-Al phase dendrites in hypoeutectic Al-7%Si-Mg cast alloy were examined by Chen et al. [83]. It was found that with addition of Ce and Sr., the grain size of the dendritic α-Al phase becomes well refined, decreasing from 150 μm to 90 μm, and is attributed to the exponential increase in nucleation frequency (1024), compared to the unmodified alloy, and restricted growth. Increasing the Ce level (0, 0.3, 0.5, 0.8 and 1 wt.%) in Al-20%Si alloy would cause a significant refinement of the primary Si crystals with the change in the morphology of the eutectic Si phase from coarse platelet like to a fine fibrous

#### *Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

structure [84]. Accordingly, the ultimate tensile strength (UTS) and elongation (% El) increased by 68.2% and 53.1%, respectively, as a result of these effects.

Ye et al. [85] investigated the influence of Ce content (0.2% and 0.4 wt.%) on the impact properties and microstructures of 2519A aluminum alloy, a new version of 2519 alloy (with a higher Cu/Mg ratio) used for armor material. Based on the results of their research, it was found that 0.2 wt.% Ce addition leads to an increase in the volume fraction of the precipitation phase, in addition to a more dispersive and homogeneous distribution of finer θ' (Al2Cu) precipitates, which result in improving the ability of the alloy for absorbing impact energy. Formation of Al8Cu4Ce phase which is thermally stable at high temperature is also expected to enhance the high temperature mechanical properties of the alloy. Yii et al. [86] reported that the addition of Ce in the Al-20%Si alloys refined the Si primary phase as the Ce additions were increased. The results also showed that addition of Ce in the range of 0.46 to 2.24 wt.% led to the formation of fine cells dispersed in the Almatrix. These cells consisted of a mixture of eutectic Si particles and Ce-containing intermetallic phases (Al3Ce and CeAl1.2 Si10.8). The amount of the intermetallic phases increased with increasing Ce addition.

Promising scientific investigations were made by Ahmad and coworkers [87, 88] on the influence of Ce on the microstructure of a commercial Al-11%Si-Cu-Mg eutectic cast alloy (ADC12). The main findings from their studies are summarized below.


Effect of solidification rate and rare earth metal addition on the microstructural characteristics and porosity formation in A356 alloy was investigated by Mahmoud et al. [89]. According to the atomic radius ratio, γLa/γSi is 1.604 and γCe/γSi is 1.559, theoretically, which shows that Ce is relatively more effective than La. These findings confirm that Sr. is the most dominating modification agent. Interaction between rare earth (RE) metals and Sr. would reduce the effectiveness of Sr. Although modification with Sr. causes the formation of shrinkage porosity, it also reacts with RE-rich intermetallics, resulting in their fragmentation. **Figure 17** reveals the distribution of La, Ce, and Sr. in RE-rich platelets, which explains the

formation of small primary crystals of αAl which resulted in grain refining of the alloy. A needle-shaped AlCeCuSi phase (Al9Ce2Cu5Si3) was also detected. The solidification of hypoeutectic Al-Si alloys with Ce addition can be described by the

*Precipitation enthalpy regarding the Ce addition and cooling rate at heating DSC curves.*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

Effects of Ce-Sr interaction on the nucleation of primary α-Al phase dendrites in hypoeutectic Al-7%Si-Mg cast alloy were examined by Chen et al. [83]. It was found that with addition of Ce and Sr., the grain size of the dendritic α-Al phase becomes well refined, decreasing from 150 μm to 90 μm, and is attributed to the exponential increase in nucleation frequency (1024), compared to the unmodified alloy, and restricted growth. Increasing the Ce level (0, 0.3, 0.5, 0.8 and 1 wt.%) in Al-20%Si alloy would cause a significant refinement of the primary Si crystals with the change in the morphology of the eutectic Si phase from coarse platelet like to a fine fibrous

Al-corner of ternary system Al-Si-Ce (**Figure 15**).

*Al-corner of ternary Al-Si- Ce system [82].*

**Figure 14.**

**Figure 15.**

**76**

#### **Figure 16**.

*Cooling curve and the first and second derivative of the 0.1 wt% Ce-containing alloy with characteristic parameters and α-Al phase arrest regions, showing points of interest [88].*


#### **Table 4.**

*Solidification characteristic parameters identified during solidification of the α-A1 phase and at the end of solidification.*

partial modification of the surrounding Si particles as less Sr. is available for modification of the eutectic Si.

> concluded that the addition of La or Ce leads only to fragmentation of the Si platelets in the case of non-modified alloys and only partial modification in the case of Sr-modified alloys. The direct advantage of the addition of RE metals to nongrained alloys is the reduction in the grain size by about 50% at 3 wt.% RE addition. According to Song et al. [79], when RE was added to 356 alloy in the amount of

> *(a) Backscattered electron image showing Ce-rich platelets in Sr-modified A356 alloy containing 1.025% wt.% Ce, and elemental distribution of (b) Ce, (c) Sr., and (d) EDS spectrum corresponding to (a) [89].*

**Figure 17.**

**Figure 18.**

**79**

*Solidification curves of unmodified 356 alloy [90].*

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

0.6 wt.%, the mean grain size was reduced by about 50%. A similar effect was observed in the work of Ibrahim et al. [91, 92] on the effect of rare earth metals on the mechanical properties of 356 and 413 alloys, as shown in **Figure 19**, in particular, in alloys 3 L and 4 L, see **Figure 19** (b, e). Due to Ce-Ti interaction, the grain refining effectiveness was reduced in the 3C and 4C alloys, as shown in **Figure 19**

#### **1.3 Effects of lanthanum additions**

It is inferred from the work of Mahmoud et al. [90] using high purity (99.5%) lanthanum or cerium, that La and Ce have more or less the same effect on the microstructures, with the La- and Ce-containing intermetallics displaying similar morphologies. Regardless of the alloy composition, an addition of 150 ppm Sr. or 0.2% RE results in improving the UTS by 25–52% in the T6 condition, with a decrease in ductility from 3% to 2.1%. The addition or RE metals (La + Ce) up to 3 wt.% leads to an increase in the freezing range through an increase in the melting point of the non-modified alloys, with decrease in the Al-Si eutectic temperature, by 12 °C and 8 °C, respectively, at 3 wt.% addition, **Figure 18**. The authors

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

**Figure 17.**

*(a) Backscattered electron image showing Ce-rich platelets in Sr-modified A356 alloy containing 1.025% wt.% Ce, and elemental distribution of (b) Ce, (c) Sr., and (d) EDS spectrum corresponding to (a) [89].*

**Figure 18.** *Solidification curves of unmodified 356 alloy [90].*

concluded that the addition of La or Ce leads only to fragmentation of the Si platelets in the case of non-modified alloys and only partial modification in the case of Sr-modified alloys. The direct advantage of the addition of RE metals to nongrained alloys is the reduction in the grain size by about 50% at 3 wt.% RE addition.

According to Song et al. [79], when RE was added to 356 alloy in the amount of 0.6 wt.%, the mean grain size was reduced by about 50%. A similar effect was observed in the work of Ibrahim et al. [91, 92] on the effect of rare earth metals on the mechanical properties of 356 and 413 alloys, as shown in **Figure 19**, in particular, in alloys 3 L and 4 L, see **Figure 19** (b, e). Due to Ce-Ti interaction, the grain refining effectiveness was reduced in the 3C and 4C alloys, as shown in **Figure 19**

partial modification of the surrounding Si particles as less Sr. is available for modi-

*Solidification characteristic parameters identified during solidification of the α-A1 phase and at the end of*

<sup>α</sup>A1 <sup>N</sup> *<sup>t</sup>*<sup>s</sup>

*Cooling curve and the first and second derivative of the 0.1 wt% Ce-containing alloy with characteristic*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

*parameters and α-Al phase arrest regions, showing points of interest [88].*

*T*<sup>α</sup>A1 <sup>M</sup> *=*˚C Minimum temperature of α-A1

*T*<sup>α</sup>A1 <sup>G</sup> *=*˚C Maximum growth temperature of α-A1 *<sup>T</sup>*<sup>α</sup>A1 <sup>R</sup> *<sup>=</sup>*˚<sup>C</sup> Recalescence temperature = *<sup>T</sup>*<sup>α</sup>A1 <sup>G</sup> *<sup>T</sup>*<sup>α</sup>A1 <sup>M</sup>

*T*<sup>α</sup>A1 <sup>N</sup> *=*˚C Nucleation temperature of α-A1 (onset of a-A1 phase)

*<sup>T</sup>*<sup>α</sup>A1 <sup>N</sup> *<sup>=</sup>*˚<sup>C</sup> Nucleation undercooling temperature = *<sup>T</sup>*<sup>α</sup>A1 <sup>N</sup> *<sup>T</sup>*<sup>α</sup>A1 <sup>M</sup>

*T*A1Si <sup>N</sup> *=*˚C Nucleation temperature of Al-Si (end of α-A1 phase)

*T*s/°C Solidus temperature (end of solidification) *t*s/s Solidus time (end of solidification) <sup>Δ</sup>*T*s/°C Total temperature range = *<sup>T</sup>*<sup>α</sup>A1 <sup>N</sup> *<sup>T</sup>*<sup>s</sup>

Δ*t*s/s Total temperature range = *t*

<sup>Δ</sup>*T*<sup>α</sup>A1*=*˚<sup>C</sup> Solidification temperature range of <sup>α</sup>-A1 phase = *<sup>T</sup>*<sup>α</sup>A1 <sup>N</sup> *<sup>T</sup>*<sup>Α</sup><sup>1</sup>Si <sup>N</sup>

**Parameter Description**

It is inferred from the work of Mahmoud et al. [90] using high purity (99.5%) lanthanum or cerium, that La and Ce have more or less the same effect on the microstructures, with the La- and Ce-containing intermetallics displaying similar morphologies. Regardless of the alloy composition, an addition of 150 ppm Sr. or 0.2% RE results in improving the UTS by 25–52% in the T6 condition, with a decrease in ductility from 3% to 2.1%. The addition or RE metals (La + Ce) up to 3 wt.% leads to an increase in the freezing range through an increase in the melting point of the non-modified alloys, with decrease in the Al-Si eutectic temperature, by 12 °C and 8 °C, respectively, at 3 wt.% addition, **Figure 18**. The authors

fication of the eutectic Si.

**Table 4.**

**78**

*solidification.*

**Figure 16**.

**1.3 Effects of lanthanum additions**

**Figure 19.**

*Effect of La and Ce addition on the grain size in A356 alloys: (a) alloy 3, no RE addition, (b) alloy 3 L, addition of 1% La, (c) alloy 3C, addition of 1% Ce. Effect of La and Ce addition on the grain size in A413 alloys: (d) alloy 4, no RE addition, (e) alloy 4 L, addition of 1% La, (f) alloy 4C, addition of 1% Ce.*

(c, f). Microstructural characterization of Al-Si cast alloys containing rare earth additions was performed by Elgallad et al. [93]. The main findings of this study were that the addition of La and/or Ce resulted in the formation of a whitecoloured Al-Si-La/Ce/(La,Ce) phase in both A356 and A413 alloys. In addition, the presence of Ti in the A356 alloy allowed for the formation of a gray-colored Al-Ti-La/Ce phase besides the Al-Si-La/Ce/(La,Ce) phase. The formation of these phases significantly increased the phase volume fraction of intermetallics in the A356 and A413 alloys. In the presence of Sr., the white-colored Al-Si-La/Ce/(La,Ce) phase was found to also contain Sr. (1 at%). No specific Sr-La/Ce intermetallic phases were detected in the microstructures of the alloys investigated. **Figure 20** (i) shows the DSC cooling curves of the A413 alloy before and after the addition of La and Ce, individually or in combination, whereas **Figure 20** (ii) shows the DSC heating curves of the A413 alloy without and with La and/or Ce.

trace additions of La (0.05% - 0.1 wt.%) on the microstructures and tensile properties of B-refined and Sr-modified Al-11%Si-1.5%Cu-0.3%Mg casting alloys were investigated by Lu et al. [95] who found that introducing La/B in the weight ratio of 2:1 produced well refined α-Al grains and modified eutectic Si particles in the alloy, as well as strengthening intermetallic precipitates, which improved the ultimate tensile strength from 234 to 270 MPa, and elongation and from 4.0 to 5.8%,

*(i) DSC cooling curves of (a) A413 and (b) A356 alloys, respectively, without and with La and/or Ce.(ii) DSC heating curves of: (a) A413 (alloy 4) and (b) A356 (alloy 3) alloys without and with La and/or Ce [93].*

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

A study on the synergistic effect of Sr. and La on the microstructure and mechanical properties of A356.2 alloy was carried out by Qiu et al. [96]. It was found that, with the addition of 0.5 wt.% Al-6Sr-7La master alloy, the alloy exhibited optimal microstructure and mechanical properties, with the secondary dendrite arm spacing (SDAS) decreasing to 17.9 μm and the acicular eutectic silicon transforming to a fibrous form. With the improved microstructure, the ultimate tensile strength, yield strength and elongation of the alloy (with 0.5 wt.% Al-6Sr-7La) increased to 228.15 MPa, 108.13 MPa and 11.92%, respectively, which were much better than those of the "traditionally treated" A356.2 alloy (using 0.2 wt.%

Al-5Ti-1B for grain refining and 0.2 wt.% Al-10Sr for Si modification).

from 22 to 9.7 μm, i.e., by 84.5% and 55.8%, respectively.

Tang et al. [97] investigated the effect of Sr. and La addition on the microstructure and mechanical properties of a secondary Al-Si-Cu-Fe alloy. The quantitative metallographic results indicated that addition of different levels of Sr. and La modification agents, added in the form of Al-10%Sr. and Al-10%La, produced varied refinement effects on the mean length of needle-like phases and the SDAS value. The total dosage of Sr. and La varied from 0.04 to 0.2 wt.% (Sr/La = 1:1). The minimum mean length of needle-like phases (Sr/La = 1:1) and the SDAS value (Sr/La = 1:5) were obtained by setting the addition amount of the modification agent at 0.12 wt.%. The mean length of the needle-like phase dropped from 364 to 55.3 μm, while the SDAS decreased

respectively.

**81**

**Figure 20.**

The results on various La additions on the microstructure of as-cast ADC12 (Al-11%Si-Cu-Mg) alloy [94] indicated that the α-Al and eutectic Si crystals were modified with the addition of 0.3 wt.% La. The eutectic Si crystals showed a granular distribution. At the same time, the alloy possessed the best mechanical properties. However, as the La addition was increased beyond 0.3 wt.%, the microstructure coarsened gradually and the mechanical properties decreased as a result.

Song et al. [77–79] analyzed the impact of different additions of La (0.0, 0.3, 0.6, and 0.9 wt.%) on the microstructure and hot crack resistance of ADC12 alloy. The results showed that, as the La added increased from 0.0% to 0.6 wt%, the structure of the α-Al phase gradually varied from a well-developed dendritic crystal into fine dendritic crystal, equiaxed crystal and spheroidal crystal; the eutectic silicon morphology varied from needle-like or tabular shape into a fine rod-like shape; the hot cracking force of the alloy also gradually decreased. Optimum alloy modification, alloy refinement and hot cracking resistance were achieved at 0.6 wt.% La addition. However, when the addition of La reached 0.9 wt.%, the excessive amount of La segregated at the grain boundaries, forming intermetallic phases.

Similarly, Chen et al. [84] evaluated the effects of combined addition of lanthanum and boron (B) on the grain refinement of Al-Si casting alloys, and found that such additions can effectively refine the grains of Al-Si alloys compared to individual addition of boron. This work also reported that with addition of La, the tensile properties of the alloy, in particular, the elongation are enhanced. The response of

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

**Figure 20.**

(c, f). Microstructural characterization of Al-Si cast alloys containing rare earth additions was performed by Elgallad et al. [93]. The main findings of this study were that the addition of La and/or Ce resulted in the formation of a whitecoloured Al-Si-La/Ce/(La,Ce) phase in both A356 and A413 alloys. In addition, the presence of Ti in the A356 alloy allowed for the formation of a gray-colored Al-Ti-La/Ce phase besides the Al-Si-La/Ce/(La,Ce) phase. The formation of these phases significantly increased the phase volume fraction of intermetallics in the A356 and A413 alloys. In the presence of Sr., the white-colored Al-Si-La/Ce/(La,Ce) phase was found to also contain Sr. (1 at%). No specific Sr-La/Ce intermetallic phases were detected in the microstructures of the alloys investigated. **Figure 20** (i) shows the DSC cooling curves of the A413 alloy before and after the addition of La and Ce, individually or in combination, whereas **Figure 20** (ii) shows the DSC heating

*Effect of La and Ce addition on the grain size in A356 alloys: (a) alloy 3, no RE addition, (b) alloy 3 L, addition of 1% La, (c) alloy 3C, addition of 1% Ce. Effect of La and Ce addition on the grain size in A413 alloys: (d) alloy 4, no RE addition, (e) alloy 4 L, addition of 1% La, (f) alloy 4C, addition of 1% Ce.*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

The results on various La additions on the microstructure of as-cast ADC12 (Al-11%Si-Cu-Mg) alloy [94] indicated that the α-Al and eutectic Si crystals were modified with the addition of 0.3 wt.% La. The eutectic Si crystals showed a granular distribution. At the same time, the alloy possessed the best mechanical properties. However, as the La addition was increased beyond 0.3 wt.%, the microstructure coarsened gradually and the mechanical properties decreased

Song et al. [77–79] analyzed the impact of different additions of La (0.0, 0.3, 0.6, and 0.9 wt.%) on the microstructure and hot crack resistance of ADC12 alloy. The results showed that, as the La added increased from 0.0% to 0.6 wt%, the structure of the α-Al phase gradually varied from a well-developed dendritic crystal into fine dendritic crystal, equiaxed crystal and spheroidal crystal; the eutectic silicon morphology varied from needle-like or tabular shape into a fine rod-like shape; the hot cracking force of the alloy also gradually decreased. Optimum alloy modification, alloy refinement and hot cracking resistance were achieved at 0.6 wt.% La addition. However, when the addition of La reached 0.9 wt.%, the excessive amount of La

Similarly, Chen et al. [84] evaluated the effects of combined addition of lanthanum and boron (B) on the grain refinement of Al-Si casting alloys, and found that such additions can effectively refine the grains of Al-Si alloys compared to individual addition of boron. This work also reported that with addition of La, the tensile properties of the alloy, in particular, the elongation are enhanced. The response of

curves of the A413 alloy without and with La and/or Ce.

segregated at the grain boundaries, forming intermetallic phases.

as a result.

**80**

**Figure 19.**

*(i) DSC cooling curves of (a) A413 and (b) A356 alloys, respectively, without and with La and/or Ce.(ii) DSC heating curves of: (a) A413 (alloy 4) and (b) A356 (alloy 3) alloys without and with La and/or Ce [93].*

trace additions of La (0.05% - 0.1 wt.%) on the microstructures and tensile properties of B-refined and Sr-modified Al-11%Si-1.5%Cu-0.3%Mg casting alloys were investigated by Lu et al. [95] who found that introducing La/B in the weight ratio of 2:1 produced well refined α-Al grains and modified eutectic Si particles in the alloy, as well as strengthening intermetallic precipitates, which improved the ultimate tensile strength from 234 to 270 MPa, and elongation and from 4.0 to 5.8%, respectively.

A study on the synergistic effect of Sr. and La on the microstructure and mechanical properties of A356.2 alloy was carried out by Qiu et al. [96]. It was found that, with the addition of 0.5 wt.% Al-6Sr-7La master alloy, the alloy exhibited optimal microstructure and mechanical properties, with the secondary dendrite arm spacing (SDAS) decreasing to 17.9 μm and the acicular eutectic silicon transforming to a fibrous form. With the improved microstructure, the ultimate tensile strength, yield strength and elongation of the alloy (with 0.5 wt.% Al-6Sr-7La) increased to 228.15 MPa, 108.13 MPa and 11.92%, respectively, which were much better than those of the "traditionally treated" A356.2 alloy (using 0.2 wt.% Al-5Ti-1B for grain refining and 0.2 wt.% Al-10Sr for Si modification).

Tang et al. [97] investigated the effect of Sr. and La addition on the microstructure and mechanical properties of a secondary Al-Si-Cu-Fe alloy. The quantitative metallographic results indicated that addition of different levels of Sr. and La modification agents, added in the form of Al-10%Sr. and Al-10%La, produced varied refinement effects on the mean length of needle-like phases and the SDAS value. The total dosage of Sr. and La varied from 0.04 to 0.2 wt.% (Sr/La = 1:1). The minimum mean length of needle-like phases (Sr/La = 1:1) and the SDAS value (Sr/La = 1:5) were obtained by setting the addition amount of the modification agent at 0.12 wt.%. The mean length of the needle-like phase dropped from 364 to 55.3 μm, while the SDAS decreased from 22 to 9.7 μm, i.e., by 84.5% and 55.8%, respectively.

Effect of solution treatment on the microstructure and mechanical properties of A356.2 alloy treated with Al–Sr–La master alloy was examined by Ding et al. [98] who found that the optimal solution treatment parameters for A356.2 aluminum alloy treated by Al–6Sr–7La are: solution treatment at 540 °C for 3 h, followed by quenching in 60 °C water – see **Figure 21**. The alloy under this condition possesses the optimal comprehensive conditions/values of microstructure, eutectic silicon morphology, UTS, YS, and EL, which is beneficial to the subsequent aging process.

The effects of Ti - La interaction on the microstructure and mechanical properties of B-refined and Sr-modified Al-11%Si alloys showed that the addition of 0.05 wt.% B induces a transformation of the eutectic Si from finely fibrous to coarse plate-like morphology in the Al-11%Si alloy modified with 0.02 wt% Sr., owing to the poisoning of impurity induced twinning (IIT) mechanism [99]. Thus, the eutectic Si growth occurred only by the twin plane re-entrant (TPRE) mechanism. Both Ti and La can neutralize the poisoning effect of the interaction between Sr. and B in the Al–11%Si alloy; however, the neutralizing effect of La is dependent on the addition sequence. The combined addition of La and B elements promoted the effective refinement of α-Al grains, but an inhomogeneous modification of the eutectic Si phase was also observed, leading to a slight decrease in the elongation. The poisoning effect can also be proved by the reduction of multiple Si twins as shown in **Figure 22**. **Figure 23** display the affinity of RE metals to react with Sr. leading to the observed loss of modification in the present alloys [100].

(a) 20 °C; (b) 30 °C; (c) 40 °C; (d) 50 °C; (e) 60 °C; and (f) 70 °C [98].

**Figure 22.**

**Figure 23.**

**83**

*(a) TEM bright field image and corresponding selected area diffraction pattern of Si particle, tilted to [011]Si zone axis, in the S5 alloy (0.1216% La, 0.0526%B, 0.0218%Sr); (b) assembly of TEM bright field images of different Si particles taken from the S5 alloy; (c) TEM bright field image and corresponding selected area*

*diffraction pattern of Si plates, tilted to [011]Si zone axis, in the S6 alloy.*

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

*Ce-Sr interactions in A356 alloy modified with 1.0% Ce + 0.01% Sr. [100].*

#### **Figure 21.**

*Effect of quenching temperatures on the microstructure of A56.2-Sr-La alloy: (a) 20 °C; (b) 30 °C; (c) 40 °C; (d) 50 °C; (e) 60 °C; and (f) 70 °C [98].*

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

#### **Figure 22.**

Effect of solution treatment on the microstructure and mechanical properties of A356.2 alloy treated with Al–Sr–La master alloy was examined by Ding et al. [98] who found that the optimal solution treatment parameters for A356.2 aluminum alloy treated by Al–6Sr–7La are: solution treatment at 540 °C for 3 h, followed by quenching in 60 °C water – see **Figure 21**. The alloy under this condition possesses the optimal comprehensive conditions/values of microstructure, eutectic silicon morphology, UTS, YS, and EL, which is beneficial to the subsequent aging process. The effects of Ti - La interaction on the microstructure and mechanical proper-

ties of B-refined and Sr-modified Al-11%Si alloys showed that the addition of 0.05 wt.% B induces a transformation of the eutectic Si from finely fibrous to coarse plate-like morphology in the Al-11%Si alloy modified with 0.02 wt% Sr., owing to the poisoning of impurity induced twinning (IIT) mechanism [99]. Thus, the eutectic Si growth occurred only by the twin plane re-entrant (TPRE) mechanism. Both Ti and La can neutralize the poisoning effect of the interaction between Sr. and B in the Al–11%Si alloy; however, the neutralizing effect of La is dependent on the addition sequence. The combined addition of La and B elements promoted the effective refinement of α-Al grains, but an inhomogeneous modification of the eutectic Si phase was also observed, leading to a slight decrease in the elongation. The poisoning effect can also be proved by the reduction of multiple Si twins as shown in **Figure 22**. **Figure 23** display the affinity of RE metals to react with Sr.

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

leading to the observed loss of modification in the present alloys [100].

**Figure 21.**

**82**

*(d) 50 °C; (e) 60 °C; and (f) 70 °C [98].*

(a) 20 °C; (b) 30 °C; (c) 40 °C; (d) 50 °C; (e) 60 °C; and (f) 70 °C [98].

*Effect of quenching temperatures on the microstructure of A56.2-Sr-La alloy: (a) 20 °C; (b) 30 °C; (c) 40 °C;*

*(a) TEM bright field image and corresponding selected area diffraction pattern of Si particle, tilted to [011]Si zone axis, in the S5 alloy (0.1216% La, 0.0526%B, 0.0218%Sr); (b) assembly of TEM bright field images of different Si particles taken from the S5 alloy; (c) TEM bright field image and corresponding selected area diffraction pattern of Si plates, tilted to [011]Si zone axis, in the S6 alloy.*

**Figure 23.** *Ce-Sr interactions in A356 alloy modified with 1.0% Ce + 0.01% Sr. [100].*

particular, Ce and La, on the microstructure and mechanical properties of aluminum alloys. While a number of these investigations have been undertaken by Chinese researchers, due likely to the easily available natural source of rare earths in the form of mischmetal, studies by other researchers are also reported. Previous studies carried out by the TAMLA research group have investigated the influence of rare earth elements and mischmetal on the performance of A356, A413.1 and other Al-Si alloys. With the more recent focus on the development of new Al-Cu based alloys for high temperature performance of automotive components, it was considered worthwhile to also investigate the effects of Ce and La rare earth metal additions to these alloys, taking into consideration low and relatively high Si levels. Putting the importance of rare earth elements into proper perspective, recently, University of Kentucky researchers have reported producing nearly pure rare earth concentrates from Kentucky coal sources [103–105]. The patent pending process developed by Honaker and Zhang is a low cost and environmentally friendly recovery process. Interest in rare earth elements is currently at its peak in the U.S.A., with the Department of Energy investing millions in research, as REs constitute essential components of diverse

**Calculated formula Shape and**

**color**

medium gray

Chinese script, white

Chinese script, white

**Suggested formula**

Al15(MnFe)3Si2 + α. Iron

> Al2MM\* Si2†

Al2MM\*

Si2† + 0.48 wt-%Sr

technologies in the high-tech and renewable energy industries.

1 Al 59.76 70.74 Al16(MnFe)4Si3 Chinese script,

La = 1.57:1)

La = 1.48:1)

**at-% Av.**

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

2 Al 23.21 41.32 Al10(CeLaPrNd)4Si9 (Ce/

3 Al 22.7 41.06 Al10(CeLaPrNd)4Si9 (Ce/

**Appendix A- Mischmetals**

**Element wt-% Av.**

> Si 10.07 11.44 Fe 20.3 11.6 Mn 8.31 4.83 Total 98.44 98.61

> Si 21.99 38.77 Ce 26.93 9.47 La 17.05 6.03 Pr 6.83 2.31 Nd 2.66 0.9 Total 98.67 98.8

> Si 22.5 39.11 Ce 26.65 9.28 La 17.78 6.24 Pr 6.53 2.26 Nd 2.39 0.8 Sr 0.48 0.26 Total 99.03 99.01

**Phase no.**

**85**

**Figure 24.**

*Bright field TEM micrograph of the La/Ce, Si and Al phases (a) squares indicate where elemental analysis was carried out; (b, c) isolated numbers indicate the zone where diffraction patterns were taken.*

#### **1.4 Effects of mixed cerium and lanthanum (RE) additions**

The microstructure and mechanical properties of the automotive A356 aluminum alloy reinforced with 0.2 wt.% Al-6Ce-3La (coded ACL) were investigated by Aguirre-De la Torre et al. [101]. In this study, the ACL was added to the molten A356 alloy in the as-received condition and also processed by another route, employing mechanical milling and powder metallurgy techniques. Scanning electron microscopic observations indicated a homogeneous dispersion of La/Ce phases using both routes. In regard to the mechanical properties, however, the modified A356 alloy with the ACL added in the as-received condition, showed an improvement in the mechanical performance of the A356 alloy over that reinforced with the mechanically milled ACL. A bright field TEM micrograph from the FIB-milled TEM sample is presented in **Figure 24** revealing the presence of three phases in different shades of gray as observed in **Figure 24(a)**.

Also, Wang et al. [102] studied the effects of mixed La and Ce rare earth additions on the microstructure and properties of Al-0.75%Mg-0.6%Si alloy. The results showed that the mixed addition of La and Ce had a positive effect on the grain refinement of the investigated alloy. Accordingly, the tensile strength and elongation of Al-0.75%Mg-0.6%Si alloy gradually increased with the increase in the amount of La and Ce added.

Another study was carried out by Du et al. [103–105] on the influence of 0.25 wt. % and 0.50 wt.% mixed additions of Ce and La on the microstructure and mechanical properties of an Al-Cu-Mn-Mg-Fe alloy. With the mixed addition, two intermetallic phases, Al8Cu4Ce and Al6Cu6La, were formed. The results also showed that the 0.25 wt.% addition could promote the formation of a denser precipitation of Al20Cu2Mn3 and Al6(Mn,Fe) phases, which improved the mechanical properties of the alloy at room temperature. However, up to 0.50 wt.% Ce-La addition promoted the formation of coarse Al8Cu4Ce phase, in addition to the Al6Cu6La and Al6(Mn, Fe) phases, which resulted in weakened mechanical properties [106, 107].

#### **2. Summary**

The review of the literature presented in this chapter has highlighted the numerous studies that have been carried out on the effects of rare earth elements, in

#### *Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

particular, Ce and La, on the microstructure and mechanical properties of aluminum alloys. While a number of these investigations have been undertaken by Chinese researchers, due likely to the easily available natural source of rare earths in the form of mischmetal, studies by other researchers are also reported. Previous studies carried out by the TAMLA research group have investigated the influence of rare earth elements and mischmetal on the performance of A356, A413.1 and other Al-Si alloys. With the more recent focus on the development of new Al-Cu based alloys for high temperature performance of automotive components, it was considered worthwhile to also investigate the effects of Ce and La rare earth metal additions to these alloys, taking into consideration low and relatively high Si levels. Putting the importance of rare earth elements into proper perspective, recently, University of Kentucky researchers have reported producing nearly pure rare earth concentrates from Kentucky coal sources [103–105]. The patent pending process developed by Honaker and Zhang is a low cost and environmentally friendly recovery process. Interest in rare earth elements is currently at its peak in the U.S.A., with the Department of Energy investing millions in research, as REs constitute essential components of diverse technologies in the high-tech and renewable energy industries.


## **Appendix A- Mischmetals**

**1.4 Effects of mixed cerium and lanthanum (RE) additions**

*carried out; (b, c) isolated numbers indicate the zone where diffraction patterns were taken.*

shades of gray as observed in **Figure 24(a)**.

amount of La and Ce added.

**2. Summary**

**84**

**Figure 24.**

The microstructure and mechanical properties of the automotive A356 aluminum alloy reinforced with 0.2 wt.% Al-6Ce-3La (coded ACL) were investigated by Aguirre-De la Torre et al. [101]. In this study, the ACL was added to the molten A356 alloy in the as-received condition and also processed by another route, employing mechanical milling and powder metallurgy techniques. Scanning electron microscopic observations indicated a homogeneous dispersion of La/Ce phases using both routes. In regard to the mechanical properties, however, the modified A356 alloy with the ACL added in the as-received condition, showed an improvement in the mechanical performance of the A356 alloy over that reinforced with the mechanically milled ACL. A bright field TEM micrograph from the FIB-milled TEM sample is presented in **Figure 24** revealing the presence of three phases in different

*Bright field TEM micrograph of the La/Ce, Si and Al phases (a) squares indicate where elemental analysis was*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

Also, Wang et al. [102] studied the effects of mixed La and Ce rare earth additions on the microstructure and properties of Al-0.75%Mg-0.6%Si alloy. The results showed that the mixed addition of La and Ce had a positive effect on the grain refinement of the investigated alloy. Accordingly, the tensile strength and elongation of Al-0.75%Mg-0.6%Si alloy gradually increased with the increase in the

Fe) phases, which resulted in weakened mechanical properties [106, 107].

The review of the literature presented in this chapter has highlighted the numerous studies that have been carried out on the effects of rare earth elements, in

Another study was carried out by Du et al. [103–105] on the influence of 0.25 wt. % and 0.50 wt.% mixed additions of Ce and La on the microstructure and mechanical properties of an Al-Cu-Mn-Mg-Fe alloy. With the mixed addition, two intermetallic phases, Al8Cu4Ce and Al6Cu6La, were formed. The results also showed that the 0.25 wt.% addition could promote the formation of a denser precipitation of Al20Cu2Mn3 and Al6(Mn,Fe) phases, which improved the mechanical properties of the alloy at room temperature. However, up to 0.50 wt.% Ce-La addition promoted the formation of coarse Al8Cu4Ce phase, in addition to the Al6Cu6La and Al6(Mn,


*Al reading could be higher than the actual content due to the small size of the examined particles.*

#### **Table A.**

*Chemical compositions of intermetallic phases observed in as-cast A413.1 alloy containing 6 wt.% MM (WDS analysis, SDAS:120 mm) [62].*



**Author details**

**Table B-2.**

Victor Songmene<sup>2</sup>

(Qc), Canada

Québec, Canada

MI, USA

**87**

Mohamed Gamal Mahmoud<sup>1</sup>

Faculty of Engineering, Egypt

, Yasser Zedan<sup>2</sup>

**Phase Color Elements (At.%) Suggested phase Al Ti Fe Cu Si Ce**

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

**white-1** 51.26 0.000 0.000 0.073 25.64 21.82 Al2CeSi **white-2** 45.37 0.000 0.000 0.276 28.55 24.59 Al2CeSi **white-3** 45.49 0.000 0.017 0.292 28.42 24.56 Al2CeSi **white-4** 45.15 0.000 0.014 0.259 29.20 24.20 Al2CeSi **white-5** 51.86 0.000 0.009 0.077 25.33 21.46 Al2CeSi

*WDS analysis of RE intermetallic phases observed with 5.0 wt.% Ce [75].*

**gray-4** 83.99 6.394 0.008 2.558 1.601 4.187 Al21Ti2La (with trace of Cu and Si) **gray-5** 85.53 6.377 0.015 0.946 1.659 4.272 Al21Ti2La(with trace of Cu and Si)

, Herebert W. Doty<sup>4</sup> and Fawzy H. Samuel<sup>3</sup>

1 Department of Mechanical Design and Production (MDP), Cairo University,

2 Département de Génie Mécanique École de Technologie Supérieure, Montréal

3 Département des Sciences Appliquées, Université du Québec à Chicoutimi,

4 Materials Technology, General Motors Global Technology Center, Warren,

© 2021 The Author(s). Licensee IntechOpen. 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,

\*Address all correspondence to: fhsamuel@uqac.ca

provided the original work is properly cited.

, Agnes-Marie Samuel<sup>3</sup>

,

\*

#### **Table B-1.**

*WDS analysis of RE intermetallic phases observed with 1.0 wt.% Ce [70].*


*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*


**Table B-2.**

**Appendix – B Ce addition**

*analysis, SDAS:120 mm) [62].*

**Phase no.**

*\**

*†*

**Table A.**

*MM: mischmetal.*

**Element wt-% Av.**

> Si 10.92 18.58 Cu 14.02 9.03 Ni 6.05 5.40 Ce 21.78 7.05 La 10.41 3.36 Pr 5.12 1.57 Nd 1.91 0.57 Total 101.2 98.16

**at-% Av.**

4 Al 31.02 52.6 Al10(CeLaPrNd)2(CuNi)3Si3

**Table B-1.**

**86**

**Phase Color Elements (At.%) Suggested phase**

*Al reading could be higher than the actual content due to the small size of the examined particles.*

**White-1** 40.54 0.000 0.069 12.38 26.96 18.58 Al9Ce4Cu2Si4 **White-2** 40.44 0.000 0.069 11.80 27.73 18.56 Al9Ce4Cu2Si4 **White-3** 46.08 0.000 0.112 12.04 24.28 16.52 Al9Ce4Cu2Si4 **White-4** 40.21 0.000 0.122 12.97 27.04 18.45 Al9Ce4Cu2Si4 **White-5** 40.57 0.000 0.096 12.24 26.96 18.78 Al9Ce4Cu2Si4

*WDS analysis of RE intermetallic phases observed with 1.0 wt.% Ce [70].*

**Phase Color Elements (At.%) Suggested phase Al Ti Fe Cu Si Ce**

**gray-1** 85.78 6.314 0.014 1.080 1.388 4.156 Al21Ti2La(with trace of Cu and Si) **gray-2** 85.40 6.228 0.015 1.410 1.431 4.134 Al21Ti2La (with trace of Cu and Si) **gray-3** 84.52 6.255 0.019 1.777 1.869 4.265 Al21Ti2La (with trace of Cu and Si)

**Gray-1** 84.41 6.426 0.025 2.050 1.526 4.009 Al21Ti2Ce (with trace of Cu and Si) **Gray-2** 83.08 6.516 0.052 2.308 2.726 4.053 Al21Ti2Ce (with trace of Cu and Si) **Gray-3** 83.88 6.696 0.015 1.963 2.278 4.019 Al21Ti2Ce (with trace of Cu and Si) **Gray-4** 83.66 6.423 0.013 2.278 2.331 3.972 Al21Ti2Ce (with trace of Cu and Si) **Gray-5** 84.60 6.630 0.025 1.734 1.872 4.003 Al21Ti2Ce (with trace of Cu and Si)

*Chemical compositions of intermetallic phases observed in as-cast A413.1 alloy containing 6 wt.% MM (WDS*

**Calculated formula Shape and**

(Ce/La = 2.1:1)

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

**color**

Plate. like, medium light gray

**Suggested formula**

Al5MM\* (CuNi) Si† + excess of Al

**Al Ti Fe Cu Si Ce**

*WDS analysis of RE intermetallic phases observed with 5.0 wt.% Ce [75].*

## **Author details**

Mohamed Gamal Mahmoud<sup>1</sup> , Yasser Zedan<sup>2</sup> , Agnes-Marie Samuel<sup>3</sup> , Victor Songmene<sup>2</sup> , Herebert W. Doty<sup>4</sup> and Fawzy H. Samuel<sup>3</sup> \*

1 Department of Mechanical Design and Production (MDP), Cairo University, Faculty of Engineering, Egypt

2 Département de Génie Mécanique École de Technologie Supérieure, Montréal (Qc), Canada

3 Département des Sciences Appliquées, Université du Québec à Chicoutimi, Québec, Canada

4 Materials Technology, General Motors Global Technology Center, Warren, MI, USA

\*Address all correspondence to: fhsamuel@uqac.ca

© 2021 The Author(s). Licensee IntechOpen. 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.

## **References**

[1] J. L. Murray A. J. McAlister, The Al-Si (Aluminum-Silicon) system, Bulletin of Alloy Phase Diagrams volume, 1984, 74, https://doi.org/10.1007/BF02868729.

[2] M. F. Ashby and D. R.H. Jones, in Engineering Materials 2 (Fourth Edition), 2013, Eutectic structure

[3] M. D. Sabatino and L. Arnberg, Castability of aluminium alloys, Transactions of The Indian Institute of Metals, 2009, 62, pp. 321–325

[4] G. Timelli and F. Bonollo, "Microstructure, Defects and Properties in Aluminum Alloys Castings: A Review," Proc. Int. Conf. Aluminium Two Thousand, Firenze (2007).

[5] Y.W. Lee, E. Chang, C.F. Chieu, "Modeling of Feeding Behavior of Solidifying Al-7Si-0.3Mg Alloy Plate Casting," Metall. Trans. B, 1990,21, pp. 715–722.

[6] L.H. Shang, F. Paray, J.E.Gruzleski, S. Bergeron, C. Mercadante, C.A. Loong, "Prediction of Microporosity in Al-Si Castings in Low Pressure Permanent Mould Casting Using Criteria Functions," Int. J. Cast Metals Res., 2004,17, pp. 193–200.

[7] S. Fox, and J. Campbell, "Visualisation of Oxide Film Defects During Solidification of Aluminium Alloys," Scripta Mater., 2000, 43, pp. 881–886.

[8] J. Campbell, "An Overview of the Effects of Bifilms on the Structure and Properties of Cast Alloys," Metall. And Mater. Trans. B, 2006, 37, pp. 857–863.

[9] J. Campbell and R.A. Harding, "Casting Technology,"TALAT 2.0 CD-ROM, EAA, Brussels (2000).

[10] J. Campbell, "Castings," Elsevier Science Ltd., Oxford (2003).

[11] E. Fiorese, F. Bonollo, G. Timelli, L. Arnberg, E. Gariboldi, NEW CLASSIFICATION OF DEFECTS AND IMPERFECTIONS FOR ALUMINUM ALLOY CASTINGS, International Journal of Metalcasting, 2015, 9, pp. 55–66.

modified with Ba, Ca, Y and Yb. Journal of light metals, 2001. 1(4): p. 219–228.

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

> [31] Colombo, M., E. Gariboldi, and A. Morri, Er addition to Al-Si-Mg-based casting alloy: Effects on microstructure, room and high temperature mechanical properties. Journal of Alloys and Compounds, 2017. 708: p. 1234–1244.

[32] Colombo, M., E. Gariboldi, and A. Morri, Influences of different Zr additions on the microstructure, room and high temperature mechanical properties of an Al-7Si-0.4 Mg alloy modified with 0.25% Er. Materials Science and Engineering: A, 2017.

[33] Li, Q.L., et al. Effect of Rare Earth Er on Microstructure and Mechanical Properties of Cast Al-Si-Mg Alloy. in Materials Science Forum. 2017. Trans

[34] Xu, C., et al., Effect of Al–P–Ti– TiC–Nd 2 O 3 modifier on the microstructure and mechanical

properties of hypereutectic Al–20wt.%

Engineering: A, 2007. 452: p. 341–346.

[35] Shi, W.X., et al., Effect of Nd on microstructure and wear resistance of hypereutectic Al-20%Si alloy. Journal of Alloys and Compounds, 2010. 508(2): p.

Si alloy. Materials Science and

[36] Weixi, S., et al., Effect of neodymium on primary silicon and mechanical properties of hypereutectic Al-15% Si alloy. Journal of Rare Earths,

[37] Ren, X., et al., Effect of Nd on microstructure and properties of 2A70 alloy. Journal of Alloys and Compounds,

[38] Tang, Q., et al., The effects of neodymium addition on the intermetallic microstructure and

[39] Ahmad, R., M. Asmael, and M. Amzar, Effect of ytterbium addition on

mechanical properties of Al-7Si-0.3 Mg-0.3 Fe alloys. Journal of Alloys and

2010. 28: p. 367–370.

Compounds, 2018.

Tech Publ.

480–485.

2017.

[22] Li, B., F. Kong, and Y. Chen, Effect of yttrium addition on microstructures and room temperature tensile properties of Ti-47 Al alloy. Journal of rare earths,

[23] Li, H.Z., et al., Effect of Y content on microstructure and mechanical properties of 2519 aluminum alloy. Transactions of Nonferrous Metals Society of China, 2007. 17(6):

[24] Zheng, L. and H. Yongmei, Effect of yttrium on the microstructure of a semisolid A356 Al alloy. Rare Metals, 2008.

[25] Sheng, M., et al., Effects of Y and Y combined with Al-5Ti-1B on the microstructure and mechanical properties of hypoeutectic Al-Si alloy.

JOM, 2015. 67(2): p. 330–335.

Forum. 2002. Trans Tech Publ.

Materials Forum. 2004.

p. 598–603.

Tech Publ.

[26] Nie, Z.R., et al. Research on rare earth in aluminum. in Materials Science

[27] Nie, Z., et al. Advanced aluminum alloys containing rare-earth erbium. in

[28] Xu, G.F., et al., Effect of trace rare earth element Er on Al-Zn-Mg alloy. Transactions of Nonferrous Metals Society of China, 2006. 16(3):

[29] Li, Y.T., et al. Alloying behavior of rare-earth Er in Al-Cu-Mg-Ag alloy. in Materials science forum. 2007. Trans

[30] Wen, S.P., et al., The effect of erbium on the microstructure and mechanical properties of Al-Mg-Mn-Zr alloy. Materials Science and Engineering

a-Structural Materials Properties Microstructure and Processing, 2009.

516(1–2): p. 42–49.

**89**

2006. 24(3): p. 352–356.

p. 1194–1198.

27(5): p. 536–540.

[12] J. Campbell, "Materials Perspective, Entrainment Defects," Mater. Sci. Technol., 2006, 22, pp. 132–136.

[13] John Campbell, in Complete Casting Handbook (Second Edition), 2015.

[14] Aluminum: Properties and Physical Metallurgy Editor: John E. Hatch, ASM International, ISBN: 978–0–87170-176-3.

[15] K.Gupta, D.J.Lloyd, S.A.Court, Precipitation hardening in Al–Mg–Si alloys with and without excess Si, Materials Science and Engineering: A , 2001, 316, pp.11–17.

[16] H. Liao, Y. Wu, K. Ding, Materials Science and Engineering A, 2013, 560, pp.811–816.

[17] Aluminum Casting Technology, American Foundrymen<s Society, inc., Des Plains, Il, 1986.

[18] Y.F. Ye, Q. Wang, J. Lu, C.T. Liu and Y. Yang, High-entropy alloy: challenges and prospects, Materials Today, 2016, 19, pp. 349–362.

[19] E. J. Pickering and N. G. Jones, High-entropy alloys: a critical assessment of their founding principles and future prospects, International Materials Reviews, 2016, 61, pp. 183–202.

[20] Knuutinen, A., et al., Modification of Al–Si alloys with Ba, Ca, Y and Yb. Journal of Light Metals, 2001. 1(4): p. 229–240.

[21] Nogita, K., et al., Mechanisms of eutectic solidification in Al–Si alloys

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

modified with Ba, Ca, Y and Yb. Journal of light metals, 2001. 1(4): p. 219–228.

**References**

[1] J. L. Murray A. J. McAlister, The Al-Si (Aluminum-Silicon) system, Bulletin of Alloy Phase Diagrams volume, 1984, 74, https://doi.org/10.1007/BF02868729.

[11] E. Fiorese, F. Bonollo, G. Timelli, L.

CLASSIFICATION OF DEFECTS AND IMPERFECTIONS FOR ALUMINUM ALLOY CASTINGS, International Journal of Metalcasting, 2015, 9,

[12] J. Campbell, "Materials Perspective, Entrainment Defects," Mater. Sci. Technol., 2006, 22, pp. 132–136.

[13] John Campbell, in Complete Casting Handbook (Second Edition), 2015.

[14] Aluminum: Properties and Physical Metallurgy Editor: John E. Hatch, ASM International, ISBN: 978–0–87170-176-3.

[15] K.Gupta, D.J.Lloyd, S.A.Court, Precipitation hardening in Al–Mg–Si alloys with and without excess Si, Materials Science and Engineering: A ,

[16] H. Liao, Y. Wu, K. Ding, Materials Science and Engineering A, 2013, 560,

[17] Aluminum Casting Technology, American Foundrymen<s Society, inc.,

[18] Y.F. Ye, Q. Wang, J. Lu, C.T. Liu and Y. Yang, High-entropy alloy: challenges and prospects, Materials Today, 2016, 19, pp. 349–362.

[19] E. J. Pickering and N. G. Jones, High-entropy alloys: a critical

assessment of their founding principles and future prospects, International Materials Reviews, 2016, 61, pp. 183–202.

[20] Knuutinen, A., et al., Modification of Al–Si alloys with Ba, Ca, Y and Yb. Journal of Light Metals, 2001. 1(4):

[21] Nogita, K., et al., Mechanisms of eutectic solidification in Al–Si alloys

2001, 316, pp.11–17.

Des Plains, Il, 1986.

pp.811–816.

p. 229–240.

Arnberg, E. Gariboldi, NEW

pp. 55–66.

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

[2] M. F. Ashby and D. R.H. Jones, in Engineering Materials 2 (Fourth Edition), 2013, Eutectic structure

[3] M. D. Sabatino and L. Arnberg, Castability of aluminium alloys, Transactions of The Indian Institute of

Metals, 2009, 62, pp. 321–325

[4] G. Timelli and F. Bonollo,

in Aluminum Alloys Castings: A Review," Proc. Int. Conf. Aluminium Two Thousand, Firenze (2007).

[5] Y.W. Lee, E. Chang, C.F. Chieu, "Modeling of Feeding Behavior of Solidifying Al-7Si-0.3Mg Alloy Plate Casting," Metall. Trans. B, 1990,21, pp.

[6] L.H. Shang, F. Paray, J.E.Gruzleski, S. Bergeron, C. Mercadante, C.A. Loong, "Prediction of Microporosity in Al-Si Castings in Low Pressure Permanent Mould Casting Using Criteria Functions," Int. J. Cast Metals Res.,

715–722.

881–886.

**88**

2004,17, pp. 193–200.

[7] S. Fox, and J. Campbell,

"Visualisation of Oxide Film Defects During Solidification of Aluminium Alloys," Scripta Mater., 2000, 43, pp.

[8] J. Campbell, "An Overview of the Effects of Bifilms on the Structure and Properties of Cast Alloys," Metall. And Mater. Trans. B, 2006, 37, pp. 857–863.

[9] J. Campbell and R.A. Harding, "Casting Technology,"TALAT 2.0 CD-

[10] J. Campbell, "Castings," Elsevier

ROM, EAA, Brussels (2000).

Science Ltd., Oxford (2003).

"Microstructure, Defects and Properties

[22] Li, B., F. Kong, and Y. Chen, Effect of yttrium addition on microstructures and room temperature tensile properties of Ti-47 Al alloy. Journal of rare earths, 2006. 24(3): p. 352–356.

[23] Li, H.Z., et al., Effect of Y content on microstructure and mechanical properties of 2519 aluminum alloy. Transactions of Nonferrous Metals Society of China, 2007. 17(6): p. 1194–1198.

[24] Zheng, L. and H. Yongmei, Effect of yttrium on the microstructure of a semisolid A356 Al alloy. Rare Metals, 2008. 27(5): p. 536–540.

[25] Sheng, M., et al., Effects of Y and Y combined with Al-5Ti-1B on the microstructure and mechanical properties of hypoeutectic Al-Si alloy. JOM, 2015. 67(2): p. 330–335.

[26] Nie, Z.R., et al. Research on rare earth in aluminum. in Materials Science Forum. 2002. Trans Tech Publ.

[27] Nie, Z., et al. Advanced aluminum alloys containing rare-earth erbium. in Materials Forum. 2004.

[28] Xu, G.F., et al., Effect of trace rare earth element Er on Al-Zn-Mg alloy. Transactions of Nonferrous Metals Society of China, 2006. 16(3): p. 598–603.

[29] Li, Y.T., et al. Alloying behavior of rare-earth Er in Al-Cu-Mg-Ag alloy. in Materials science forum. 2007. Trans Tech Publ.

[30] Wen, S.P., et al., The effect of erbium on the microstructure and mechanical properties of Al-Mg-Mn-Zr alloy. Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 2009. 516(1–2): p. 42–49.

[31] Colombo, M., E. Gariboldi, and A. Morri, Er addition to Al-Si-Mg-based casting alloy: Effects on microstructure, room and high temperature mechanical properties. Journal of Alloys and Compounds, 2017. 708: p. 1234–1244.

[32] Colombo, M., E. Gariboldi, and A. Morri, Influences of different Zr additions on the microstructure, room and high temperature mechanical properties of an Al-7Si-0.4 Mg alloy modified with 0.25% Er. Materials Science and Engineering: A, 2017.

[33] Li, Q.L., et al. Effect of Rare Earth Er on Microstructure and Mechanical Properties of Cast Al-Si-Mg Alloy. in Materials Science Forum. 2017. Trans Tech Publ.

[34] Xu, C., et al., Effect of Al–P–Ti– TiC–Nd 2 O 3 modifier on the microstructure and mechanical properties of hypereutectic Al–20wt.% Si alloy. Materials Science and Engineering: A, 2007. 452: p. 341–346.

[35] Shi, W.X., et al., Effect of Nd on microstructure and wear resistance of hypereutectic Al-20%Si alloy. Journal of Alloys and Compounds, 2010. 508(2): p. 480–485.

[36] Weixi, S., et al., Effect of neodymium on primary silicon and mechanical properties of hypereutectic Al-15% Si alloy. Journal of Rare Earths, 2010. 28: p. 367–370.

[37] Ren, X., et al., Effect of Nd on microstructure and properties of 2A70 alloy. Journal of Alloys and Compounds, 2017.

[38] Tang, Q., et al., The effects of neodymium addition on the intermetallic microstructure and mechanical properties of Al-7Si-0.3 Mg-0.3 Fe alloys. Journal of Alloys and Compounds, 2018.

[39] Ahmad, R., M. Asmael, and M. Amzar, Effect of ytterbium addition on microstructure and hardness of Al-6.5Si-1Zn secondary cast alloy. 2006.

[40] Xiao, D., et al., Effect of rare earth Yb addition on mechanical properties of Al–5 3Cu–0 8Mg–0 6 Ag alloy. Materials science and technology, 2007. 23(10): p. 1156–1160.

[41] Zhang, X.M., et al., Effects of Yb addition on microstructures and mechanical properties of 2519A aluminum alloy plate. Transactions of Nonferrous Metals Society of China, 2010. 20(5): p. 727–731.

[42] Li, B., et al., Microstructure evolution and modification mechanism of the ytterbium modified Al-7.5%Si-0.45%Mg alloys. Journal of Alloys and Compounds, 2011. 509(7): p. 3387–3392.

[43] Li, J.H., et al., Refinement of Eutectic Si Phase in Al-5Si Alloys with Yb Additions. Metallurgical and Materials Transactions a-Physical Metallurgy and Materials Science, 2013. 44a(2): p. 669–681.

[44] Hu, Z., et al., Solidification behavior, microstructure and silicon twinning of Al-10Si alloys with Yb addition. Journal of Rare Earths, 2018.

[45] Li, Q., et al., Mechanical Properties and Microstructural Evolution of Yb-Modified Al-20% Si Alloy. Journal of Materials Engineering and Performance, 2018: p. 1–10.

[46] Nogita, K., S.D. McDonald, and A. K. Dahle, Eutectic modification of Al-Si alloys with rare earth metals. Materials Transactions, 2004. 45(2): p. 323–326.

[47] Chen, Z.W., P. Chen, and C.Y. Ma, Microstructures and mechanical properties of Al-Cu-Mn alloy with La and Sm addition. Rare Metals, 2012. 31 (4): p. 332–335.

[48] Qiu, H., H. Yan, and Z. Hu, Effect of samarium (Sm) addition on the

microstructures and mechanical properties of Al–7Si–0.7 Mg alloys. Journal of Alloys and Compounds, 2013. 567: p. 77–81.

alloys. Materials Characterization, 2016.

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

2015, American Foundry Society:

[64] El Sebaie, O., The Effect of Mischmetal, Cooling Rate and Heat Treatment on the Microstructure and

Hardness of 319, 356, and 413

and performance of A390 alloy. Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 2009.

modification treatment on the

527(1–2): p. 146–149.

p. 2685–2694.

Aluminum-silicon Alloys. 2006, UQAC.

[65] Li, Y.G., et al., Effect of co-addition of RE, Fe and Mn on the microstructure

[66] Zhu, M., et al., Effect of mischmetal

microstructure, tensile properties, and fracture behavior of Al-7.0%Si-0.3%Mg foundry aluminum alloys. Journal of Materials Science, 2011. 46(8):

[67] Zhu, M., et al., Effects of T6 heat treatment on the microstructure, tensile properties, and fracture behavior of the modified A356 alloys. Materials & Design, 2012. 36: p. 243–249.

[68] Mousavi, G.S., M. Emamy, and J. Rassizadehghani, The effect of

mischmetal and heat treatment on the microstructure and tensile properties of A357 Al-Si casting alloy. Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 2012. 556: p. 573–581.

[69] Mahmoud, M.G., et al., Effect of Solidification Rate and Rare Earth Metal

[70] Mahmoud, M.G., et al., Effect of Rare Earth Metals, Sr, and Ti Addition on the Microstructural Characterization of A413.1 Alloy. Advances in Materials Science and Engineering, 2017. 2017.

Addition on the Microstructural Characteristics and Porosity Formation in A356 Alloy. Advances in Materials Science and Engineering, 2017: p. 1–15.

Columbus, Ohio.

[57] Ravi, M., et al., The effect of mischmetal addition on the structure and mechanical properties of a cast Al-7Si-0.3 Mg alloy containing excess iron (up to 0.6 Pct). Metallurgical and Materials Transactions A, 2002. 33(2):

[58] Weiwei, W.A.N., et al., Study of

properties of an Al-Cu-Mg-Si cast alloy. Rare Metals, 2006. 25(6): p. 129–132.

[59] Chong, C., et al., Influences of complex modification of P and RE on microstructure and mechanical properties of hypereutectic Al-20Si alloy. Transactions of Nonferrous Metals Society of China, 2007. 17(2): p.

[60] El Sebaie, O., et al., The effects of mischmetal, cooling rate and heat treatment on the hardness of A319. 1, A356. 2 and A413. 1 Al–Si casting alloys. Materials Science and Engineering: A,

[61] El Sebaie , O., et al., The effects of mischmetal, cooling rate and heat treatment on the eutectic Si particle characteristics of A319.1, A356.2 and A413.1 Al–Si casting alloys. Materials Science and Engineering: A, 2008. 480

[62] Elsebaie, O., F.H. Samuel, and S. Al Kahtani, Intermetallic phases observed in non-modified and Sr modified Al-Si cast alloys containing mischmetal. International Journal of Cast Metals Research, 2013. 26(1): p. 1–15.

[63] Doty, H.W., S.A. Alkahtani, O. Elsebaie and F.H. Samuel, Influence of Metallurgical Parameters on the Impact Toughness of Near Eutectic Al-Si Alloys, in 119th Metalcasting Congress, AFS 2015, Columbus, OH, April 21–23, 2015.

2008. 486(1): p. 241–252.

(1–2): p. 342–355.

**91**

rare earth element effect on microstructures and mechanical

120: p. 129–142.

p. 391–400.

301–306.

[49] Zhang, W., et al., Effects of Sc content on the microstructure of As-Cast Al-7wt.% Si alloys. Materials Characterization, 2012. 66: p. 104–110.

[50] Patakham, U., J. Kajornchaiyakul, and C. Limmaneevichitr, Modification mechanism of eutectic silicon in Al–6Si– 0.3 Mg alloy with scandium. Journal of Alloys and Compounds, 2013. 575: p. 273–284.

[51] C. Y. Young, L. Qingchaun, and J. Zhuling, "Influence of cerium and mischmetal on the hardness and brightness of Al-Mg-Si alloys," Journal of the Less Common Metals, vol. 110, 1985, pp. 175–178.

[52] M. Ravi, U. T. S. Pillai, B. C. Pai, A. D. Damodaran, and E. S. Dwarakadasa, "A study of the influence of mischmetal additions to ai-7si-0.3 mg (lm 25/356) alloy," Metallurgical and Materials Transactions A, vol. 27, 1996, pp. 1283–1292.

[53] Chang, J.Y., et al., Rare earth concentration in the primary Si crystal in rare earth added Al-21wt.%Si alloy. Scripta Materialia, 1998. 39(3): p. 307–314.

[54] Chang, J.Y., I.G. Moon, and C.S. Choi, Refinement of cast microstructure of hypereutectic Al-Si alloys through the addition of rare earth metals. Journal of Materials Science, 1998. 33(20): p. 5015– 5023.

[55] Li, J.H., et al., Modification of eutectic Si in Al-Si alloys with Eu addition. Acta Materialia, 2015. 84: p. 153–163.

[56] Mao, F., et al., The interaction between Eu and P in high purity Al-7Si *Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

alloys. Materials Characterization, 2016. 120: p. 129–142.

microstructure and hardness of Al-6.5Si-1Zn secondary cast alloy. 2006.

Al–5 3Cu–0 8Mg–0 6 Ag alloy. Materials science and technology, 2007.

23(10): p. 1156–1160.

2010. 20(5): p. 727–731.

44a(2): p. 669–681.

2018: p. 1–10.

(4): p. 332–335.

**90**

[42] Li, B., et al., Microstructure evolution and modification mechanism of the ytterbium modified Al-7.5%Si-0.45%Mg alloys. Journal of Alloys and Compounds, 2011. 509(7): p. 3387–3392.

[43] Li, J.H., et al., Refinement of Eutectic Si Phase in Al-5Si Alloys with Yb Additions. Metallurgical and Materials Transactions a-Physical Metallurgy and Materials Science, 2013.

[44] Hu, Z., et al., Solidification behavior, microstructure and silicon twinning of Al-10Si alloys with Yb addition. Journal of Rare Earths, 2018.

[45] Li, Q., et al., Mechanical Properties and Microstructural Evolution of Yb-Modified Al-20% Si Alloy. Journal of Materials Engineering and Performance,

[46] Nogita, K., S.D. McDonald, and A. K. Dahle, Eutectic modification of Al-Si alloys with rare earth metals. Materials Transactions, 2004. 45(2): p. 323–326.

[47] Chen, Z.W., P. Chen, and C.Y. Ma, Microstructures and mechanical properties of Al-Cu-Mn alloy with La and Sm addition. Rare Metals, 2012. 31

[48] Qiu, H., H. Yan, and Z. Hu, Effect of samarium (Sm) addition on the

[40] Xiao, D., et al., Effect of rare earth Yb addition on mechanical properties of microstructures and mechanical properties of Al–7Si–0.7 Mg alloys. Journal of Alloys and Compounds, 2013.

[49] Zhang, W., et al., Effects of Sc content on the microstructure of As-Cast Al-7wt.% Si alloys. Materials Characterization, 2012. 66: p. 104–110.

[50] Patakham, U., J. Kajornchaiyakul, and C. Limmaneevichitr, Modification mechanism of eutectic silicon in Al–6Si– 0.3 Mg alloy with scandium. Journal of Alloys and Compounds, 2013. 575:

[51] C. Y. Young, L. Qingchaun, and J. Zhuling, "Influence of cerium and mischmetal on the hardness and brightness of Al-Mg-Si alloys," Journal of the Less Common Metals, vol. 110,

[52] M. Ravi, U. T. S. Pillai, B. C. Pai, A. D. Damodaran, and E. S. Dwarakadasa, "A study of the influence of mischmetal additions to ai-7si-0.3 mg (lm 25/356) alloy," Metallurgical and Materials Transactions A, vol. 27, 1996,

[53] Chang, J.Y., et al., Rare earth concentration in the primary Si crystal in rare earth added Al-21wt.%Si alloy. Scripta Materialia, 1998. 39(3):

[54] Chang, J.Y., I.G. Moon, and C.S. Choi, Refinement of cast microstructure of hypereutectic Al-Si alloys through the addition of rare earth metals. Journal of Materials Science, 1998. 33(20): p. 5015–

[55] Li, J.H., et al., Modification of eutectic Si in Al-Si alloys with Eu addition. Acta Materialia, 2015. 84:

[56] Mao, F., et al., The interaction between Eu and P in high purity Al-7Si

567: p. 77–81.

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

p. 273–284.

1985, pp. 175–178.

pp. 1283–1292.

p. 307–314.

5023.

p. 153–163.

[41] Zhang, X.M., et al., Effects of Yb addition on microstructures and mechanical properties of 2519A aluminum alloy plate. Transactions of Nonferrous Metals Society of China,

[57] Ravi, M., et al., The effect of mischmetal addition on the structure and mechanical properties of a cast Al-7Si-0.3 Mg alloy containing excess iron (up to 0.6 Pct). Metallurgical and Materials Transactions A, 2002. 33(2): p. 391–400.

[58] Weiwei, W.A.N., et al., Study of rare earth element effect on microstructures and mechanical properties of an Al-Cu-Mg-Si cast alloy. Rare Metals, 2006. 25(6): p. 129–132.

[59] Chong, C., et al., Influences of complex modification of P and RE on microstructure and mechanical properties of hypereutectic Al-20Si alloy. Transactions of Nonferrous Metals Society of China, 2007. 17(2): p. 301–306.

[60] El Sebaie, O., et al., The effects of mischmetal, cooling rate and heat treatment on the hardness of A319. 1, A356. 2 and A413. 1 Al–Si casting alloys. Materials Science and Engineering: A, 2008. 486(1): p. 241–252.

[61] El Sebaie , O., et al., The effects of mischmetal, cooling rate and heat treatment on the eutectic Si particle characteristics of A319.1, A356.2 and A413.1 Al–Si casting alloys. Materials Science and Engineering: A, 2008. 480 (1–2): p. 342–355.

[62] Elsebaie, O., F.H. Samuel, and S. Al Kahtani, Intermetallic phases observed in non-modified and Sr modified Al-Si cast alloys containing mischmetal. International Journal of Cast Metals Research, 2013. 26(1): p. 1–15.

[63] Doty, H.W., S.A. Alkahtani, O. Elsebaie and F.H. Samuel, Influence of Metallurgical Parameters on the Impact Toughness of Near Eutectic Al-Si Alloys, in 119th Metalcasting Congress, AFS 2015, Columbus, OH, April 21–23, 2015.

2015, American Foundry Society: Columbus, Ohio.

[64] El Sebaie, O., The Effect of Mischmetal, Cooling Rate and Heat Treatment on the Microstructure and Hardness of 319, 356, and 413 Aluminum-silicon Alloys. 2006, UQAC.

[65] Li, Y.G., et al., Effect of co-addition of RE, Fe and Mn on the microstructure and performance of A390 alloy. Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 2009. 527(1–2): p. 146–149.

[66] Zhu, M., et al., Effect of mischmetal modification treatment on the microstructure, tensile properties, and fracture behavior of Al-7.0%Si-0.3%Mg foundry aluminum alloys. Journal of Materials Science, 2011. 46(8): p. 2685–2694.

[67] Zhu, M., et al., Effects of T6 heat treatment on the microstructure, tensile properties, and fracture behavior of the modified A356 alloys. Materials & Design, 2012. 36: p. 243–249.

[68] Mousavi, G.S., M. Emamy, and J. Rassizadehghani, The effect of mischmetal and heat treatment on the microstructure and tensile properties of A357 Al-Si casting alloy. Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 2012. 556: p. 573–581.

[69] Mahmoud, M.G., et al., Effect of Solidification Rate and Rare Earth Metal Addition on the Microstructural Characteristics and Porosity Formation in A356 Alloy. Advances in Materials Science and Engineering, 2017: p. 1–15.

[70] Mahmoud, M.G., et al., Effect of Rare Earth Metals, Sr, and Ti Addition on the Microstructural Characterization of A413.1 Alloy. Advances in Materials Science and Engineering, 2017. 2017.

[71] Jiang, W.M., et al., Effects of rare earth elements addition on microstructures, tensile properties and fractography of A357 alloy. Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 2014. 597: p. 237–244.

[72] Zhang, J., et al., Microstructural development of Al–15wt.% Mg 2 Si in situ composite with mischmetal addition. Materials Science and Engineering: A, 2000. 281(1): p. 104– 112.

[73] Zhang, H.R., et al., Cooling Rate Sensitivity of RE-Containing Grain Refiner and Its Impact on the Microstructure and Mechanical Properties of A356 Alloy. Acta Metallurgica Sinica-English Letters, 2016. 29(5): p. 414–421.

[74] Dang, B., Z.Y. Jian, and J.F. Xu, Effects of rare-earth element addition and heat treatment on the microstructures and mechanical properties of Al-25% Si alloy. International Journal of Materials Research, 2017. 108(4): p. 269–274.

[75] Mahmoud, M.G., A.M. Samuel, H. W. Doty and F.H. Samuel, *Role of Heat Treatment on the Tensile Properties and Fractography of Al–1.2Si–2.4Cu and Al– 8.0Si–2.4Cu Cast Alloys Modified with Ce, La and Sr Addition.* International Journal of Metalcasting, 2020. **14**(1): p. 218–242.

[76] Mahmoud, M.G., A.M. Samuel, H. W. Doty and F.H. Samuel, *Effect of the Addition of La and Ce on the Solidification Behavior of Al–Cu and Al– Si–Cu Cast Alloys.* International Journal of Metalcasting, 2020. **14**(1): p. 191–206.

[77] Song, M., K.H. Chen, and L.P. Huang, Effects of Ce and Ti on the microstructures and mechanical properties of an Al-Cu-Mg-Ag alloy. Rare Metals, 2007. 26(1): p. 28–32.

[78] Song, M., D.H. Xiao, and F.Q. Zhang, Effect of Ce on the thermal stability of the Omega phase in an Al-Cu-Mg-Ag alloy. Rare Metals, 2009. 28 (2): p. 156–159.

(2016). Microstructural and Mechanical Properties of Al-20%Si Containing Cerium. Procedia Chemistry. 19. 304– 310. 10.1016/j.proche.2016.03.015.

*Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

*containing rare earth additions.*

p. 1337–1359.

p. 405–410.

Materials, 2018.

Materials, 2017. **9**(1).

Philosophical Magazine, 2018. **98**(15):

[94] Huang, X. and H. Yan, Effect of trace La addition on the microstructure and mechanical property of as-cast ADC12 Al-Alloy. Journal of Wuhan University of Technology-Materials Science Edition, 2013. 28(1): p. 202–205.

[95] Lu, T., et al., Effects of La addition on the microstructure and tensile properties of Al-Si-Cu-Mg casting alloys. International Journal of Minerals Metallurgy and Materials, 2015. 22(4):

[96] Qiu, C., et al., Synergistic effect of Sr and La on the microstructure and mechanical properties of A356. 2 alloy.

strontium and lanthanum simultaneous

mechanical properties of the secondary Al-Si-Cu-Fe alloy. Journal of Rare Earths, 2017. 35(5): p. 485–493.

[98] Ding, J., et al., Effect of Solution Treatment on Microstructure and Mechanical Properties of A356. 2 Aluminum Alloy Treated With Al–Sr– La Master Alloy. Advanced Engineering

[99] Li, C., et al., Effects of Ti and La Additions on the Microstructures and Mechanical Properties of B-Refined and Sr-Modified Al–11Si Alloys. Metals and Materials International, 2018: p. 1–10.

[100] Alkahtani, S.A., E.M. Elgallad, M. M. Tash, A.M. Samuel and F.H. Samuel, Effect of rare earth metals on the microstructure ao AL-SI based alloys.

[101] Aguirre-De la Torre, E., et al., Mechanical properties of the A356 aluminum alloy modified with La/Ce.

Materials & Design, 2016.

[97] Tang, P., et al., Influence of

addition on microstructure and

[87] Ahmad, R. and M.B.A. Asmael, Influence of Cerium on Microstructure and Solidification of Eutectic Al–Si

Manufacturing Processes, 2015. 31(15):

[88] Ahmad, R., et al., Reduction in secondary dendrite arm spacing in cast eutectic Al–Si piston alloys by cerium addition. International Journal of Minerals, Metallurgy, and Materials,

[89] Mahmoud, M.G., A.M. Samuel, H. W. Doty and F.H. Samuel, *Formation of Rare Earth Intermetallics in Al–Cu Cast Alloys*, in *Light Metals Symposium held at the 149th Annual Meeting and Exhibition, TMS 2020, 23 February 2020 through 27 February 2020*. 2020: San Diego; United

[90] Mahmoud, M.G., E.M. Elgallad, M. F. Ibrahim and F.H. Samuel, *Effect of Rare Earth Metals on Porosity Formation in A356 Alloy.* International Journal of Metalcasting, 2018. **12**(2): p. 251–265.

[91] Ibrahim, M.F., M.H. Abdelaziz, A. M. Samuel, H.W. Doty and F.H. Samuel,

[92] Ibrahim, A.I., E.M. Elgallad, A.M. Samuel, H.W. Doty and F.H. Samuel, *Effects of heat treatment and testing temperature on the tensile properties of Al–*

[93] Elgallad, E.M., M.F. Ibrahim, H.W. Doty and F.H. Samuel, *Microstructural characterisation of Al–Si cast alloys*

*Cu and Al–Cu–Si based alloys.* International Journal of Materials Research, 2018. **109**(4): p. 314–331.

*Effect of Rare Earth Metals on the Mechanical Properties and Fractography of Al–Si-Based Alloys.* International Journal of Metalcasting, 2020. **14**(1): p.

Piston Alloy. Materials and

2017. 24(1): p. 91–101.

States. p. 241–246.

108–124.

**93**

p. 1948–1957.

[79] Song, X.C., H. Yan, and F.H. Chen, Impact of Rare Earth Element La on Microstructure and Hot Crack Resistance of ADC12 Alloy. Journal of Wuhan University of Technology-Materials Science Edition, 2018. 33(1): p. 193–197.

[80] Xiao, D.H., J.N. Wang, and D.Y. Ding, Effect of minor cerium additions on microstructure and mechanical properties of cast Al – Cu – Mg – Ag alloy. Materials Science and Technology, 2004. 20(10): p. 1237–1240.

[81] Voncina, M., et al., Effect of Ce on solidification and mechanical properties of A360 alloy. Journal of Alloys and Compounds, 2011. 509(27): p. 7349–7355.

[82] Voncina, M., et al., Microstructure and grain refining performance of Ce on A380 alloy. Journal of Mining and Metallurgy, Section B: Metallurgy, 2012. 48(2): p. 265–272.

[83] Chen, Z.W., et al., Kinetic nucleation of primary alpha (Al) dendrites in Al-7%Si-Mg cast alloys with Ce and Sr additions. Transactions of Nonferrous Metals Society of China, 2013. 23(12): p. 3561–3567.

[84] Chen, Y., et al., Effects of combinative addition of lanthanum and boron on grain refinement of Al-Si casting alloys. Materials & Design, 2014. 64: p. 423–426.

[85] Ye, L.Y., et al., Influence of Ce addition on impact properties and microstructures of 2519A aluminum alloy. Materials Science and Engineering: A, 2013. 582: p. 84–90.

[86] Yii, S.L. & Norazman, Anas & Nasir, Ramdziah & Anasyida, A.S.. *Applications of Rare Earth Metals in Al-Si Cast Alloys DOI: http://dx.doi.org/10.5772/intechopen.96011*

(2016). Microstructural and Mechanical Properties of Al-20%Si Containing Cerium. Procedia Chemistry. 19. 304– 310. 10.1016/j.proche.2016.03.015.

[71] Jiang, W.M., et al., Effects of rare

[78] Song, M., D.H. Xiao, and F.Q. Zhang, Effect of Ce on the thermal stability of the Omega phase in an Al-Cu-Mg-Ag alloy. Rare Metals, 2009. 28

[79] Song, X.C., H. Yan, and F.H. Chen, Impact of Rare Earth Element La on Microstructure and Hot Crack

Resistance of ADC12 Alloy. Journal of Wuhan University of Technology-Materials Science Edition, 2018. 33(1):

[80] Xiao, D.H., J.N. Wang, and D.Y. Ding, Effect of minor cerium additions on microstructure and mechanical properties of cast Al – Cu – Mg – Ag alloy. Materials Science and Technology,

[81] Voncina, M., et al., Effect of Ce on solidification and mechanical properties of A360 alloy. Journal of Alloys and

[82] Voncina, M., et al., Microstructure and grain refining performance of Ce on A380 alloy. Journal of Mining and Metallurgy, Section B: Metallurgy, 2012.

dendrites in Al-7%Si-Mg cast alloys with Ce and Sr additions. Transactions of Nonferrous Metals Society of China,

combinative addition of lanthanum and boron on grain refinement of Al-Si casting alloys. Materials & Design, 2014.

[85] Ye, L.Y., et al., Influence of Ce addition on impact properties and microstructures of 2519A aluminum

Engineering: A, 2013. 582: p. 84–90.

[86] Yii, S.L. & Norazman, Anas & Nasir, Ramdziah & Anasyida, A.S..

alloy. Materials Science and

2004. 20(10): p. 1237–1240.

Compounds, 2011. 509(27):

[83] Chen, Z.W., et al., Kinetic nucleation of primary alpha (Al)

2013. 23(12): p. 3561–3567.

[84] Chen, Y., et al., Effects of

p. 7349–7355.

48(2): p. 265–272.

64: p. 423–426.

(2): p. 156–159.

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

p. 193–197.

microstructures, tensile properties and fractography of A357 alloy. Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 2014. 597: p. 237–244.

[72] Zhang, J., et al., Microstructural development of Al–15wt.% Mg 2 Si in situ composite with mischmetal addition. Materials Science and Engineering: A, 2000. 281(1): p. 104–

[73] Zhang, H.R., et al., Cooling Rate Sensitivity of RE-Containing Grain Refiner and Its Impact on the Microstructure and Mechanical Properties of A356 Alloy. Acta Metallurgica Sinica-English Letters,

[74] Dang, B., Z.Y. Jian, and J.F. Xu, Effects of rare-earth element addition

[75] Mahmoud, M.G., A.M. Samuel, H. W. Doty and F.H. Samuel, *Role of Heat Treatment on the Tensile Properties and Fractography of Al–1.2Si–2.4Cu and Al– 8.0Si–2.4Cu Cast Alloys Modified with Ce, La and Sr Addition.* International Journal of Metalcasting, 2020. **14**(1):

[76] Mahmoud, M.G., A.M. Samuel, H. W. Doty and F.H. Samuel, *Effect of the*

*Solidification Behavior of Al–Cu and Al– Si–Cu Cast Alloys.* International Journal

*Addition of La and Ce on the*

of Metalcasting, 2020. **14**(1):

[77] Song, M., K.H. Chen, and L.P. Huang, Effects of Ce and Ti on the microstructures and mechanical properties of an Al-Cu-Mg-Ag alloy. Rare Metals, 2007. 26(1): p. 28–32.

2016. 29(5): p. 414–421.

and heat treatment on the microstructures and mechanical properties of Al-25% Si alloy. International Journal of Materials Research, 2017. 108(4): p. 269–274.

p. 218–242.

p. 191–206.

**92**

earth elements addition on

112.

[87] Ahmad, R. and M.B.A. Asmael, Influence of Cerium on Microstructure and Solidification of Eutectic Al–Si Piston Alloy. Materials and Manufacturing Processes, 2015. 31(15): p. 1948–1957.

[88] Ahmad, R., et al., Reduction in secondary dendrite arm spacing in cast eutectic Al–Si piston alloys by cerium addition. International Journal of Minerals, Metallurgy, and Materials, 2017. 24(1): p. 91–101.

[89] Mahmoud, M.G., A.M. Samuel, H. W. Doty and F.H. Samuel, *Formation of Rare Earth Intermetallics in Al–Cu Cast Alloys*, in *Light Metals Symposium held at the 149th Annual Meeting and Exhibition, TMS 2020, 23 February 2020 through 27 February 2020*. 2020: San Diego; United States. p. 241–246.

[90] Mahmoud, M.G., E.M. Elgallad, M. F. Ibrahim and F.H. Samuel, *Effect of Rare Earth Metals on Porosity Formation in A356 Alloy.* International Journal of Metalcasting, 2018. **12**(2): p. 251–265.

[91] Ibrahim, M.F., M.H. Abdelaziz, A. M. Samuel, H.W. Doty and F.H. Samuel, *Effect of Rare Earth Metals on the Mechanical Properties and Fractography of Al–Si-Based Alloys.* International Journal of Metalcasting, 2020. **14**(1): p. 108–124.

[92] Ibrahim, A.I., E.M. Elgallad, A.M. Samuel, H.W. Doty and F.H. Samuel, *Effects of heat treatment and testing temperature on the tensile properties of Al– Cu and Al–Cu–Si based alloys.* International Journal of Materials Research, 2018. **109**(4): p. 314–331.

[93] Elgallad, E.M., M.F. Ibrahim, H.W. Doty and F.H. Samuel, *Microstructural characterisation of Al–Si cast alloys*

*containing rare earth additions.* Philosophical Magazine, 2018. **98**(15): p. 1337–1359.

[94] Huang, X. and H. Yan, Effect of trace La addition on the microstructure and mechanical property of as-cast ADC12 Al-Alloy. Journal of Wuhan University of Technology-Materials Science Edition, 2013. 28(1): p. 202–205.

[95] Lu, T., et al., Effects of La addition on the microstructure and tensile properties of Al-Si-Cu-Mg casting alloys. International Journal of Minerals Metallurgy and Materials, 2015. 22(4): p. 405–410.

[96] Qiu, C., et al., Synergistic effect of Sr and La on the microstructure and mechanical properties of A356. 2 alloy. Materials & Design, 2016.

[97] Tang, P., et al., Influence of strontium and lanthanum simultaneous addition on microstructure and mechanical properties of the secondary Al-Si-Cu-Fe alloy. Journal of Rare Earths, 2017. 35(5): p. 485–493.

[98] Ding, J., et al., Effect of Solution Treatment on Microstructure and Mechanical Properties of A356. 2 Aluminum Alloy Treated With Al–Sr– La Master Alloy. Advanced Engineering Materials, 2018.

[99] Li, C., et al., Effects of Ti and La Additions on the Microstructures and Mechanical Properties of B-Refined and Sr-Modified Al–11Si Alloys. Metals and Materials International, 2018: p. 1–10.

[100] Alkahtani, S.A., E.M. Elgallad, M. M. Tash, A.M. Samuel and F.H. Samuel, Effect of rare earth metals on the microstructure ao AL-SI based alloys. Materials, 2017. **9**(1).

[101] Aguirre-De la Torre, E., et al., Mechanical properties of the A356 aluminum alloy modified with La/Ce. Journal of Rare Earths, 2013. 31(8): p. 811–816.

[102] Wang, S.C., et al. Effects of La and Ce Mixed Rare Earth on Microstructure and Properties of Al-Mg-Si Aluminum Alloy. in Materials Science Forum. 2017. Trans Tech Publ.

[103] Du, J.D., et al., Effect of CeLa addition on the microstructures and mechanical properties of Al-Cu-Mn-Mg-Fe alloy. Materials Characterization, 2017. 123: p. 42–50.

[104] UK Researchers First to Produce High Grade Rare Earths From Coal: @universityofky; 2017 [updated 2017- 11-20. Available from: https://uknow. uky.edu/research/uk-researchers-firstproduce-high-grade-rare-earths-coal

[105] Mao, F., et al., Effect of Eu addition on the microstructures and mechanical properties of A356 aluminum alloys. Journal of Alloys and Compounds, 2015. 650: p. 896–906.

[106] Tzeng, Y.-C., et al., Effects of scandium addition on iron-bearing phases and tensile properties of Al–7Si– 0.6 Mg alloys. Materials Science and Engineering: A, 2014. 593: p. 103–110.

[107] Xu, C., et al., The synergic effects of Sc and Zr on the microstructure and mechanical properties of Al-Si-Mg alloy. Materials & Design, 2015. 88: p. 485–492.

**95**

**Chapter 5**

**Abstract**

**1. Introduction**

Recent Advances of High Entropy

Alloys: High Entropy Superalloys

*Modupeola Dada, Patricia Popoola, Ntombizodwa Mathe,* 

This study reviews the recent technological advancements in manufacturing technique; laser surface modification and material; High Entropy Superalloys. High Entropy Superalloys are current potential alternatives to nickel superalloys for gas turbine applications and these superalloys are presented as the most promising

Energy transformation comprises the turbine, which is an inner combustion device and a spinning engine that utilizes water, wind steam, helium and air to produce work [1]. Kaygusuz [2] stated that dams use turbines as an electrical generator producing electricity for residential and industrial consumption. Nonetheless, in 1939, the first jet engine that powered an aircraft was built consisting of the combustion chamber, the turbine and the compressor [3]. This turbine used air as its working fluid in an internal combustion engine and this engine, in turn, removes enough chemical energy to convert it to mechanical energy from the fuel source while using the working fluid to drive the propeller and the engine [4]. Bell and Partridge [5] anticipated that the Joule cycle is a theoretical cycle for gas turbine applications, where both expansion and compression routes take place in a rotating machine [6]. This comprises some reversible processes such as the turbine using the expansion process and fluid friction for an increase in entropy which causes a spontaneous reaction using the compression method in the Brayton cycle [7]. The gas turbine is characterized by extended overhaul intervals, an increased operating speed, less moving parts, availability, low maintenance, reliability, long life span and rugged design [8]. The design of a turbine engine dictates its performance and the performance requirement are determined by the shaft house power developed in certain temperature conditions which may be extreme. Therefore, the need for high-performance materials becomes necessary because one factor which affects the efficiency of the engine; the turbine inlet temperature is made up of materials which are designed to reduce flow losses and must withstand erosion, corrosion and stress at elevated temperatures [9]. According to Reed [10],

*Nicholus Malatji, Thabo Lengopeng and Afolabi Ayodeji*

**Keywords:** high entropy alloys, high entropy superalloys, nickel superalloys,

*Samson Adeosun, Sisa Pityana, Olufemi Aramide,* 

material for gas turbine engine applications.

turbine engine, laser surface modification

## **Chapter 5**

Journal of Rare Earths, 2013. 31(8):

[103] Du, J.D., et al., Effect of CeLa addition on the microstructures and mechanical properties of Al-Cu-Mn-Mg-Fe alloy. Materials Characterization,

[104] UK Researchers First to Produce High Grade Rare Earths From Coal: @universityofky; 2017 [updated 2017- 11-20. Available from: https://uknow. uky.edu/research/uk-researchers-firstproduce-high-grade-rare-earths-coal

[105] Mao, F., et al., Effect of Eu addition on the microstructures and mechanical properties of A356 aluminum alloys. Journal of Alloys and Compounds, 2015.

[106] Tzeng, Y.-C., et al., Effects of scandium addition on iron-bearing phases and tensile properties of Al–7Si– 0.6 Mg alloys. Materials Science and Engineering: A, 2014. 593: p. 103–110.

[107] Xu, C., et al., The synergic effects of Sc and Zr on the microstructure and mechanical properties of Al-Si-Mg alloy.

Materials & Design, 2015. 88:

[102] Wang, S.C., et al. Effects of La and Ce Mixed Rare Earth on Microstructure and Properties of Al-Mg-Si Aluminum Alloy. in Materials Science Forum. 2017.

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

p. 811–816.

Trans Tech Publ.

2017. 123: p. 42–50.

650: p. 896–906.

p. 485–492.

**94**

## Recent Advances of High Entropy Alloys: High Entropy Superalloys

*Modupeola Dada, Patricia Popoola, Ntombizodwa Mathe, Samson Adeosun, Sisa Pityana, Olufemi Aramide, Nicholus Malatji, Thabo Lengopeng and Afolabi Ayodeji*

## **Abstract**

This study reviews the recent technological advancements in manufacturing technique; laser surface modification and material; High Entropy Superalloys. High Entropy Superalloys are current potential alternatives to nickel superalloys for gas turbine applications and these superalloys are presented as the most promising material for gas turbine engine applications.

**Keywords:** high entropy alloys, high entropy superalloys, nickel superalloys, turbine engine, laser surface modification

#### **1. Introduction**

Energy transformation comprises the turbine, which is an inner combustion device and a spinning engine that utilizes water, wind steam, helium and air to produce work [1]. Kaygusuz [2] stated that dams use turbines as an electrical generator producing electricity for residential and industrial consumption. Nonetheless, in 1939, the first jet engine that powered an aircraft was built consisting of the combustion chamber, the turbine and the compressor [3]. This turbine used air as its working fluid in an internal combustion engine and this engine, in turn, removes enough chemical energy to convert it to mechanical energy from the fuel source while using the working fluid to drive the propeller and the engine [4]. Bell and Partridge [5] anticipated that the Joule cycle is a theoretical cycle for gas turbine applications, where both expansion and compression routes take place in a rotating machine [6]. This comprises some reversible processes such as the turbine using the expansion process and fluid friction for an increase in entropy which causes a spontaneous reaction using the compression method in the Brayton cycle [7]. The gas turbine is characterized by extended overhaul intervals, an increased operating speed, less moving parts, availability, low maintenance, reliability, long life span and rugged design [8]. The design of a turbine engine dictates its performance and the performance requirement are determined by the shaft house power developed in certain temperature conditions which may be extreme. Therefore, the need for high-performance materials becomes necessary because one factor which affects the efficiency of the engine; the turbine inlet temperature is made up of materials which are designed to reduce flow losses and must withstand erosion, corrosion and stress at elevated temperatures [9]. According to Reed [10],

#### **Figure 1.**

*Comparison graph between nickel super alloy and high entropy superalloy.*

superalloys especially Nickel Superalloys are materials generally used at elevated temperatures for these gas turbine applications attributed to their elevated temperature strength, corrosion resistance, excellent formability, cost and low density [11]. However, the nickel-based superalloy has a maximum service temperature, not over 650 °*C* attributed to the conversion of γ´ precipitate strengthening matrix to the δ phase over time [12]. More so, the nucleation and growth of some cavities along the transverse grain boundaries of these materials are the gas turbine airfoil's failure mechanisms [13]. Therefore; a need to develop new materials with improved properties was necessary and this was achieved by transforming conventional material into new ones via advanced industrial reproduction [14]. Miracle, Tsai [15] proposed High Entropy Superalloys (HESAs) as a new class of amalgam with superior properties compared to traditional superalloys as shown in **Figure 1**.

Their elemental composition, lower densities, high configurational entropy and core effects alongside possessing the γ´ precipitate reinforcement phase makes this superalloy a preferred alternative material for turbine engine applications [16]. In a previous study, additive manufacturing was presented as a potential advance manufacturing technique as opposed to conventional arc melting and casting fabrication processes. This study attempts to present HESAs as a promising material for gas turbine engine applications, as opposed to traditional Nickel-based superalloys [17].

### **2. Advances in material development**

#### **2.1 Super alloys**

Superalloys are stable materials; they do not oxidize or fall apart in very harsh environments and at high temperatures. These amalgams are used for power generation, industrial, marine and aerospace applications [10]. They are characterized by their excellent heat and oxidation resistance at elevated temperatures, high melting temperature, and high-temperature mechanical strength, good fracture toughness, and stress-rupture, creep resistance [18]. In general, superalloys contain more of Co, Ni, Cr or Fe but less of Ta, Hf, W, Cr, B, Mo, Nb, Al, Zr, C, Ti because these elements adversely affect the properties of the blend. Superalloys have a typical face-centred cubic structure and are characterized by a γ´ precipitate with operating temperatures above 600 °*C* [19]. This phase gives the superalloy a

**97**

at a density of 7 g/cm3

*Recent Advances of High Entropy Alloys: High Entropy Superalloys*

lizes the FCC structure and strengthens the superalloy [23].

principal yield strength which increases with a temperature rise. They may have equiaxed or columnar grain structures without exhibiting high-angle grain boundaries, which at high temperatures are sites for damage accumulation [20]. According to Graybill, Li [21], Superalloy's strengthening mechanism includes dispersion strengthening, solid solution and precipitation strengthening [22]. The dispersion of chemically inert carbides and nitride enhances the strength of the superalloy. Precipitation of all intermetallic phases, namely; carbides and FCC matrix γ´ precipitate enhances the strength of the superalloy through Ti, Cb, Ta and Al, which promote the formation of the γ´ precipitate. Finally, solid solution strengthening with tungsten, columbium, rhenium, molybdenum, rhenium and tantalum stabi-

Gessinger and Bomford [24] suggested that Superalloys for gas turbine applications are widely fabricated using powder metallurgy. However, Bewlay, Gigliotti [25] fabricated the turbine disks using hot die forging and roll forming. Lavella, Berruti [26] studied the residual stresses in Inconel 718 turbine disks fabricated by milling while Groh, Gabb [27] developed a turbine disk using the casting technique. Compared with these conventional techniques, the powder metallurgy process produces turbine disks which are extremely difficult to forge; die life is relatively poor and die fill is extremely difficult but not with additive manufacturing.

Gas turbine engines require higher temperatures for efficiency. This hightemperature application, therefore, requires excellent emission control with an advance in the combustion hardware of the engine. Nickel superalloys materials were developed for this purpose and they make up about half of the weight of materials used in turbine engines [28, 29]. They have an FCC nickel matrix which is stable enough for the alloy to be used for combustion liners, blades, vanes, thermal barrier coatings, burners and are also applied to bear loads of over 75% of their emergent melting temperature. This is attributed to their characteristic hightemperature rupture and creep resistance, lifetime expectancy, low operating costs and excellent thermal efficiency [30]. Nickel superalloys are also used in space vehicles, submarines, petrochemical equipment and nuclear reactors. Nickel-based superalloy 718 (IN718) is widely used in wrought or cast at 540°*C* for rotors in gas turbine applications [14]. Nickel superalloy 925 and 725 having good corrosion resistance are applied in the oil and gas industry where carbon dioxide, hydrogen sulphide, free Sulfur and chloride levels are significantly high. Nickel superalloy 706 (IN706) is used for power generation for its large diameter and lower concentrations of other alloying elements. Alloy 685 (Waspaloy) with high-temperature strength and age hardening is widely used for gas turbine engine applications [31]. The superalloy is resistant to corrosion and oxidation whilst withstanding extreme atmospheric conditions while in service. Other compositions are; Rene'

Jet engines, Inconel Alloy 600 used for stills, condensers, heaters and evaporator tubes. Alloy 601 is used for pollution control, power and aerospace applications [32]. Nimonic 90 is used for turbine disc, blades, hot-working tools and forging. However, Nickel-based superalloy is difficult to machine attributed to its hardness, toughness, and they possess high heat resistance at elevated temperatures alongside low thermal conductivity [33]. Machining at high pressure causes work hardening rapidly, which invariably causes the alloy component to warp. Furthermore, Nickel superalloys can be easily replaced with alloys that have high creep strength and niobium silicide was an appropriate system to replace nickel superalloy having 170 MPa creep strength and

to the nickel superalloy led to the enhancement of the superalloy's creep strength,

until recently. Research showed that the addition of Ru and Re

N6 used in

*DOI: http://dx.doi.org/10.5772/intechopen.96661*

*2.1.1 Nickel superalloys*

#### *Recent Advances of High Entropy Alloys: High Entropy Superalloys DOI: http://dx.doi.org/10.5772/intechopen.96661*

principal yield strength which increases with a temperature rise. They may have equiaxed or columnar grain structures without exhibiting high-angle grain boundaries, which at high temperatures are sites for damage accumulation [20]. According to Graybill, Li [21], Superalloy's strengthening mechanism includes dispersion strengthening, solid solution and precipitation strengthening [22]. The dispersion of chemically inert carbides and nitride enhances the strength of the superalloy. Precipitation of all intermetallic phases, namely; carbides and FCC matrix γ´ precipitate enhances the strength of the superalloy through Ti, Cb, Ta and Al, which promote the formation of the γ´ precipitate. Finally, solid solution strengthening with tungsten, columbium, rhenium, molybdenum, rhenium and tantalum stabilizes the FCC structure and strengthens the superalloy [23].

Gessinger and Bomford [24] suggested that Superalloys for gas turbine applications are widely fabricated using powder metallurgy. However, Bewlay, Gigliotti [25] fabricated the turbine disks using hot die forging and roll forming. Lavella, Berruti [26] studied the residual stresses in Inconel 718 turbine disks fabricated by milling while Groh, Gabb [27] developed a turbine disk using the casting technique. Compared with these conventional techniques, the powder metallurgy process produces turbine disks which are extremely difficult to forge; die life is relatively poor and die fill is extremely difficult but not with additive manufacturing.

### *2.1.1 Nickel superalloys*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

superalloys especially Nickel Superalloys are materials generally used at elevated temperatures for these gas turbine applications attributed to their elevated temperature strength, corrosion resistance, excellent formability, cost and low density [11]. However, the nickel-based superalloy has a maximum service temperature, not over 650 °*C* attributed to the conversion of γ´ precipitate strengthening matrix to

*Comparison graph between nickel super alloy and high entropy superalloy.*

 phase over time [12]. More so, the nucleation and growth of some cavities along the transverse grain boundaries of these materials are the gas turbine airfoil's failure mechanisms [13]. Therefore; a need to develop new materials with improved properties was necessary and this was achieved by transforming conventional material into new ones via advanced industrial reproduction [14]. Miracle, Tsai [15] proposed High Entropy Superalloys (HESAs) as a new class of amalgam with superior properties compared to traditional superalloys as shown in **Figure 1**.

Their elemental composition, lower densities, high configurational entropy and core effects alongside possessing the γ´ precipitate reinforcement phase makes this superalloy a preferred alternative material for turbine engine applications [16]. In a previous study, additive manufacturing was presented as a potential advance manufacturing technique as opposed to conventional arc melting and casting fabrication processes. This study attempts to present HESAs as a promising material for gas turbine engine applications, as opposed to traditional Nickel-based superalloys [17].

Superalloys are stable materials; they do not oxidize or fall apart in very harsh

environments and at high temperatures. These amalgams are used for power generation, industrial, marine and aerospace applications [10]. They are characterized by their excellent heat and oxidation resistance at elevated temperatures, high melting temperature, and high-temperature mechanical strength, good fracture toughness, and stress-rupture, creep resistance [18]. In general, superalloys contain more of Co, Ni, Cr or Fe but less of Ta, Hf, W, Cr, B, Mo, Nb, Al, Zr, C, Ti because these elements adversely affect the properties of the blend. Superalloys have a typical face-centred cubic structure and are characterized by a γ´ precipitate with operating temperatures above 600 °*C* [19]. This phase gives the superalloy a

**2. Advances in material development**

**2.1 Super alloys**

**96**

the δ

**Figure 1.**

Gas turbine engines require higher temperatures for efficiency. This hightemperature application, therefore, requires excellent emission control with an advance in the combustion hardware of the engine. Nickel superalloys materials were developed for this purpose and they make up about half of the weight of materials used in turbine engines [28, 29]. They have an FCC nickel matrix which is stable enough for the alloy to be used for combustion liners, blades, vanes, thermal barrier coatings, burners and are also applied to bear loads of over 75% of their emergent melting temperature. This is attributed to their characteristic hightemperature rupture and creep resistance, lifetime expectancy, low operating costs and excellent thermal efficiency [30]. Nickel superalloys are also used in space vehicles, submarines, petrochemical equipment and nuclear reactors. Nickel-based superalloy 718 (IN718) is widely used in wrought or cast at 540°*C* for rotors in gas turbine applications [14]. Nickel superalloy 925 and 725 having good corrosion resistance are applied in the oil and gas industry where carbon dioxide, hydrogen sulphide, free Sulfur and chloride levels are significantly high. Nickel superalloy 706 (IN706) is used for power generation for its large diameter and lower concentrations of other alloying elements. Alloy 685 (Waspaloy) with high-temperature strength and age hardening is widely used for gas turbine engine applications [31]. The superalloy is resistant to corrosion and oxidation whilst withstanding extreme atmospheric conditions while in service. Other compositions are; Rene' N6 used in Jet engines, Inconel Alloy 600 used for stills, condensers, heaters and evaporator tubes. Alloy 601 is used for pollution control, power and aerospace applications [32]. Nimonic 90 is used for turbine disc, blades, hot-working tools and forging. However, Nickel-based superalloy is difficult to machine attributed to its hardness, toughness, and they possess high heat resistance at elevated temperatures alongside low thermal conductivity [33]. Machining at high pressure causes work hardening rapidly, which invariably causes the alloy component to warp. Furthermore, Nickel superalloys can be easily replaced with alloys that have high creep strength and niobium silicide was an appropriate system to replace nickel superalloy having 170 MPa creep strength and at a density of 7 g/cm3 until recently. Research showed that the addition of Ru and Re to the nickel superalloy led to the enhancement of the superalloy's creep strength,

however; these additions are expensive and cause density inversion which results in the defect [34]. This superalloy's stability is limited at very high temperatures [35].

According to Durand-Charre [36], Nickel-based Superalloys are majorly FCC phase structured. However, in aluminum-nickel superalloy systems, a second precipitate phase is formed, which is usually Ni3Al in a composition containing an ordered intermetallic structure [37]. The γ´ phase relies on the cooling rate of the superalloy through the solvus temperature of 894 °*C* [38]. A fast cooling rate promotes a unimodal distribution of the γ´ precipitate, therefore, an increasing the volume of the γ´ phase through rapid solidification is essential to the strengthening properties of the superalloy [28]. Although the precipitate morphologies can be modified through heat treatment and other secondary phases observed in Nickelbased superalloys are ordered FCC γ´, FCC carbides, ordered body-centred tetragonal γ´´ and ordered orthorhombic intermetallic phases [39].

Pollock and Tin [28] did an intensive review on nickel superalloys and the authors stated that the commercial superalloys comprise Co, Cr, W, Mo, Ta, W, Nb, Re, Ti, Al, C, Hf, Y, B and Zr. The yield strength of the nickel superalloy is between 900–1300 MPa at room temperature and the fatigue life at 593 °*C* is 600 MPa at 106 cycles and 109 cycles. Additions of Re, Nb, W and Mo can be added for the solid solution strengthening of the superalloy. Y, Ta and Cr contribute to enhance the corrosion and oxidation properties of the superalloy. While Zr, C, Hf and B are carbides or borides forming agents that help enhances the mechanical properties of the superalloy as they are situated at the grain boundaries. The creep rupture life attains about 1100 °*C* at 137 MPa stress level after 1000 h which is about 90% fraction of the melting point signifying the need for innovative advanced materials having higher melting points for the hottest regions of the turbine engine.

#### *2.1.2 High entropy superalloys (HESAs)*

Throughout the years, alloys utilized for commercial reasons were structured by choosing an element which framed the network of the whole component with the addition of essential solutes to the base component [40, 41]. The blends of these combinations were reduced as could reasonably be expected for the immense development of mass intermetallic mixes existing within the molar atomic proportions of these alloys, hence, attaining a 40% mark or more. Along these lines, the intermetallic phases reduce the quality of the alloys while in service [42, 43]. Therefore, a need arose to search for alloys with atomic percentages lesser than 35% and the possibilities of combining many metallic principal elements in several atomic compositions were further investigated [44]. According to Ye, Wang [45], an innovative class of alloys with these attributes was discovered more than a decade ago by mixing multiple principal elements in equimolar or near-equimolar compositions. Yeh, Chen [46] named the alloys 'High Entropy Alloys' (HEAs). The authors defined HEAs as amalgams having compositions with at least five principal metallic elements, with these components having a molar atomic proportion between 5 and 35% [47]. Studies on HEAs have concluded that most HEAs comprise simple FCC, BCC or HCP solid solutions phase attributed to their thigh-entropy effect [48]. Wang, Li [49] suggested that these solid solution phases with little or no intermetallic matrix enable HEAs to have outstanding properties such as strength, extraordinary mechanical and physical properties at cryogenic temperatures, plastic strain, fracture strength and good ductility; they possess elevated-temperature oxidation resistance and excellent work hardenability and have been reported to possess distinctive magnetic and tribological properties [50, 51]. Furthermore, Senkov, Wilks [52] reported that HEAs are exceptional refractory materials and their fatigueresistance were reported to exceed conventional alloys by Hemphill [53].

**99**

*Recent Advances of High Entropy Alloys: High Entropy Superalloys*

In the literature, the development in the solid solution strengthening of High Entropy Alloys (HEAs) and the precipitation hardening properties of the alloys at temperatures above 1100 °*C* , led to the discovery of High Entropy Superalloys (HESAs). Yeh, Tsao [17] stated that these superalloys are simply HEAs with the bulk of γ´ precipitates and they are described by their high elongation at room temperature, compressive strength, lower densities, creep resistance and ultimate tensile strengths at elevated temperatures. Tsao, Yeh [54] suggested that High Entropy Superalloy (HESA) is made up of a first elemental content containing at least 35 at.% and each principal reinforcement elemental combination will have a second elemental content of more than 5 at%, for example, Ni40.7Al7.8Cr12.2Fe11.58Co20.6Ti7.2 HESA. Senkov, Isheim [55] developed a refractory high entropy superalloy and the authors anticipated that the first and second elemental composition content is derived by the mixing entropy of more than 1.5 R (R is the ideal gas constant) alongside the principal strengthening elemental composition, respectively.

Chen, Chang [56] studied the hierarchical microstructural strengthening of HESAs and the composition of strengthening elements can consist of Cu, Fe, Ti, Zr, Co, V, Al, Nb, Cr and Mn. While the overall structure can comprise Mn, Ni, Fe, Ti, Co, Cr and V while for the grain boundary strengthening; C, B and Hf are added but must not be over 15% of the superalloy's total compositional weight. Refractory elements like Ru, Ta, Re, Mo and W can be added but must also contain less than 15% of the total superalloy's weight [55]. Tsao, Yeh [57] in 2013 recommended the development of superalloys using HEAs microstructure with single phases and an additional second phase for elevated temperatures applications. Yeh, Tsao [17] then fabricated Ni40.7Al7.8Co20.6Cr12.2 Fe11.5Ti7.2 high entropy superalloy (HESA) via casting method. The authors reported that the microstructure of the composition was stable at elevated temperatures and the superalloy was made up of γ´ nanosized

.

Daoud, Manzoni [58] developed Al8Co17Cr17Cu8Fe17Ni33 HESA using thermoscalc, the authors compared the results with Alloy 800H and IN617. They reported that the HESA had higher tensile strength, this was attributed to two phases; one spherical γ´ precipitate which was less than 20 nm after the aging temperature at 700 °*C* and another less than 350 nm with an elongated morphology. He, Wang [59] fabricated Fe94Co94Ni94Cr94Ti2Al4 HESA with γ´ nanosized precipitates to manipulate the thermomechanical properties of HESAs, and they argued that the superalloy had γ´ nanosized precipitate with an outstanding yield strength and elongation [60]. According to Xiao, Gregoire [61], scanning alternating current calorimetry

Wang, Zhou [62] investigated Al0.2CrFeCoNi2Cu0.2 HESA and discovered the γ´ nanosized precipitate with 30% elongation. Tsao, Yeh [57] developed seven HESAs using the elements Ni, Fe, Al, Cr, Co, Ti by vacuum arc melting. They stated that the development of the γ precipitates in the superalloy was due to Fe, Cr elements and the γ´ matrix, they stated that substituting Ni with Ti enhances the thermal stability of HESAs thus encouraging the γ´ matrix and by controlling the elemental compositional partitioning in the middle of the γ-γ´ phase, the thermal properties of the high entropy γ matrix can be improved. More so, at elevated temperatures after long term exposures L12 γ´ nanosized precipitates were formed without topological closed packed phases. Gwalani, Soni [63] examined Al0.3CoCrFeNi2 and

Al0.3CoCrFeNi HESAs and the authors observed γ´ precipitate in the Al0.3CoCrFeNi super alloy until 550 °*C* but were replaced with a B2 phase at 700 °*C* after annealing

and yield strength superior to nickel superalloys at 1200°*C* . Kai, Cheng [66]

Senkov, Isheim [55] and Li, Lee [65] tested AlMo0.5NbTa0.5TiZr HESA by powder metallurgy and they all observed that the superalloy possessed high thermal stability

can be used to quantify the thermomechanical properties of superalloy.

*DOI: http://dx.doi.org/10.5772/intechopen.96661*

precipitates with a density lower than 8 g/cm3

attributed to the increase in aluminum content [64].

#### *Recent Advances of High Entropy Alloys: High Entropy Superalloys DOI: http://dx.doi.org/10.5772/intechopen.96661*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

nal γ´´ and ordered orthorhombic intermetallic phases [39].

however; these additions are expensive and cause density inversion which results in the defect [34]. This superalloy's stability is limited at very high temperatures [35]. According to Durand-Charre [36], Nickel-based Superalloys are majorly FCC phase structured. However, in aluminum-nickel superalloy systems, a second precipitate phase is formed, which is usually Ni3Al in a composition containing an ordered intermetallic structure [37]. The γ´ phase relies on the cooling rate of the superalloy through the solvus temperature of 894 °*C* [38]. A fast cooling rate promotes a unimodal distribution of the γ´ precipitate, therefore, an increasing the volume of the γ´ phase through rapid solidification is essential to the strengthening properties of the superalloy [28]. Although the precipitate morphologies can be modified through heat treatment and other secondary phases observed in Nickelbased superalloys are ordered FCC γ´, FCC carbides, ordered body-centred tetrago-

Pollock and Tin [28] did an intensive review on nickel superalloys and the authors stated that the commercial superalloys comprise Co, Cr, W, Mo, Ta, W, Nb, Re, Ti, Al, C, Hf, Y, B and Zr. The yield strength of the nickel superalloy is between 900–1300 MPa at room temperature and the fatigue life at 593 °*C* is 600 MPa at

solution strengthening of the superalloy. Y, Ta and Cr contribute to enhance the corrosion and oxidation properties of the superalloy. While Zr, C, Hf and B are carbides or borides forming agents that help enhances the mechanical properties of the superalloy as they are situated at the grain boundaries. The creep rupture life attains about 1100 °*C* at 137 MPa stress level after 1000 h which is about 90% fraction of the melting point signifying the need for innovative advanced materials

having higher melting points for the hottest regions of the turbine engine.

Throughout the years, alloys utilized for commercial reasons were structured by choosing an element which framed the network of the whole component with the addition of essential solutes to the base component [40, 41]. The blends of these combinations were reduced as could reasonably be expected for the immense development of mass intermetallic mixes existing within the molar atomic proportions of these alloys, hence, attaining a 40% mark or more. Along these lines, the intermetallic phases reduce the quality of the alloys while in service [42, 43]. Therefore, a need arose to search for alloys with atomic percentages lesser than 35% and the possibilities of combining many metallic principal elements in several atomic compositions were further investigated [44]. According to Ye, Wang [45], an innovative class of alloys with these attributes was discovered more than a decade ago by mixing multiple principal elements in equimolar or near-equimolar compositions. Yeh, Chen [46] named the alloys 'High Entropy Alloys' (HEAs). The authors defined HEAs as amalgams having compositions with at least five principal metallic elements, with these components having a molar atomic proportion between 5 and 35% [47]. Studies on HEAs have concluded that most HEAs comprise simple FCC, BCC or HCP solid solutions phase attributed to their thigh-entropy effect [48]. Wang, Li [49] suggested that these solid solution phases with little or no intermetallic matrix enable HEAs to have outstanding properties such as strength, extraordinary mechanical and physical properties at cryogenic temperatures, plastic strain, fracture strength and good ductility; they possess elevated-temperature oxidation resistance and excellent work hardenability and have been reported to possess distinctive magnetic and tribological properties [50, 51]. Furthermore, Senkov, Wilks [52] reported that HEAs are exceptional refractory materials and their fatigueresistance were reported to exceed conventional alloys by Hemphill [53].

cycles. Additions of Re, Nb, W and Mo can be added for the solid

**98**

106

cycles and 109

*2.1.2 High entropy superalloys (HESAs)*

In the literature, the development in the solid solution strengthening of High Entropy Alloys (HEAs) and the precipitation hardening properties of the alloys at temperatures above 1100 °*C* , led to the discovery of High Entropy Superalloys (HESAs). Yeh, Tsao [17] stated that these superalloys are simply HEAs with the bulk of γ´ precipitates and they are described by their high elongation at room temperature, compressive strength, lower densities, creep resistance and ultimate tensile strengths at elevated temperatures. Tsao, Yeh [54] suggested that High Entropy Superalloy (HESA) is made up of a first elemental content containing at least 35 at.% and each principal reinforcement elemental combination will have a second elemental content of more than 5 at%, for example, Ni40.7Al7.8Cr12.2Fe11.58Co20.6Ti7.2 HESA. Senkov, Isheim [55] developed a refractory high entropy superalloy and the authors anticipated that the first and second elemental composition content is derived by the mixing entropy of more than 1.5 R (R is the ideal gas constant) alongside the principal strengthening elemental composition, respectively.

Chen, Chang [56] studied the hierarchical microstructural strengthening of HESAs and the composition of strengthening elements can consist of Cu, Fe, Ti, Zr, Co, V, Al, Nb, Cr and Mn. While the overall structure can comprise Mn, Ni, Fe, Ti, Co, Cr and V while for the grain boundary strengthening; C, B and Hf are added but must not be over 15% of the superalloy's total compositional weight. Refractory elements like Ru, Ta, Re, Mo and W can be added but must also contain less than 15% of the total superalloy's weight [55]. Tsao, Yeh [57] in 2013 recommended the development of superalloys using HEAs microstructure with single phases and an additional second phase for elevated temperatures applications. Yeh, Tsao [17] then fabricated Ni40.7Al7.8Co20.6Cr12.2 Fe11.5Ti7.2 high entropy superalloy (HESA) via casting method. The authors reported that the microstructure of the composition was stable at elevated temperatures and the superalloy was made up of γ´ nanosized precipitates with a density lower than 8 g/cm3 .

Daoud, Manzoni [58] developed Al8Co17Cr17Cu8Fe17Ni33 HESA using thermoscalc, the authors compared the results with Alloy 800H and IN617. They reported that the HESA had higher tensile strength, this was attributed to two phases; one spherical γ´ precipitate which was less than 20 nm after the aging temperature at 700 °*C* and another less than 350 nm with an elongated morphology. He, Wang [59] fabricated Fe94Co94Ni94Cr94Ti2Al4 HESA with γ´ nanosized precipitates to manipulate the thermomechanical properties of HESAs, and they argued that the superalloy had γ´ nanosized precipitate with an outstanding yield strength and elongation [60]. According to Xiao, Gregoire [61], scanning alternating current calorimetry can be used to quantify the thermomechanical properties of superalloy.

Wang, Zhou [62] investigated Al0.2CrFeCoNi2Cu0.2 HESA and discovered the γ´ nanosized precipitate with 30% elongation. Tsao, Yeh [57] developed seven HESAs using the elements Ni, Fe, Al, Cr, Co, Ti by vacuum arc melting. They stated that the development of the γ precipitates in the superalloy was due to Fe, Cr elements and the γ´ matrix, they stated that substituting Ni with Ti enhances the thermal stability of HESAs thus encouraging the γ´ matrix and by controlling the elemental compositional partitioning in the middle of the γ-γ´ phase, the thermal properties of the high entropy γ matrix can be improved. More so, at elevated temperatures after long term exposures L12 γ´ nanosized precipitates were formed without topological closed packed phases. Gwalani, Soni [63] examined Al0.3CoCrFeNi2 and Al0.3CoCrFeNi HESAs and the authors observed γ´ precipitate in the Al0.3CoCrFeNi super alloy until 550 °*C* but were replaced with a B2 phase at 700 °*C* after annealing attributed to the increase in aluminum content [64].

Senkov, Isheim [55] and Li, Lee [65] tested AlMo0.5NbTa0.5TiZr HESA by powder metallurgy and they all observed that the superalloy possessed high thermal stability and yield strength superior to nickel superalloys at 1200°*C* . Kai, Cheng [66]

examined the oxidation behavior of a HESA in O2 environments. The Ni2FeCoCrAl0.5 HESA oxidation kinetics at 900 °*C* followed a parabolic-rate law forming scales which was dependent on the oxygen pressure. The results showed that the oxidation rates increased with an increase in oxygen pressure however, the kinetics of massloss was observed. Shafiee, Nili-Ahmadabadi [67] designed a wrought HESA using Phacomp and CALPHAD technique. The reports showed that the superalloy comprised of γ´ nanosized precipitates with lower densities, excellent workability and high thermal stability than Inconel 718 alloy and Waspaloy. Saito, Chen [68] discussed the influence of heat treatments on HESA microstructural evolution and results showed the cast HESA had coarsened γ´ precipitates attributed to microsegregation which decreased the solidus making the γ´ solvus unclear. Finally, Zhang, Huo [69] prepared cast Ni48-xCo18 Fe9.3Al9.7Cr10.5Ti4.5Mox HESA to investigate the mechanical and microstructural properties of the superalloy and they concluded that HESAs exhibits good compressive strength at elevated temperature, elongation and tensile strength at room temperatures than nickel superalloy.

HESAs are stable at elevated temperatures compared with commercial Rene' N6, Udimet 700 and Hastelloy X superalloys, this attributed to their sluggish diffusion and high entropy effect. At high temperatures, Nickel-based superalloys form intermetallic topological closed-packed (TCP) phases rich in Fe-Cr because of the high iron content in less than 100 h at 900 °*C* and this TCP phases formed is detrimental to the stability of superalloys at high temperatures [70]. However, at 900 °*C* in more than 200 h, there was no TCP phase observed in Ni40.7Al7.8Cr12.2 Fe11.58Co20.6Ti7.2 HESA and the γ-γ´ microstructure of the superalloy remained stable after isothermal aging for 500 h at 1050 °*C* [57]. The elevated temperature strength of HESAs has been reported to be higher than that of IN 617. This can be attributed to enhancing the APB energy, increasing the lattice distortion and/or adding refractory Ta and W elements in high concentration to the compositional system.

Yeh and Tsao [71] did a thorough analysis of HESAs with the elemental composition of Fe, C, Al, Mo, Cr, Ti, Ni, Co, Ta, W and Nb. The siderophile element was Nickel while the strengthening element was Nb and C. The authors reported that the HESA's microstructure comprised an FCC, L12 crystal structure and γ´ phase while the superalloys had hardness values of 400–470 HV at room temperature. At elevated temperatures, the HESAs hardness values recorded were between 300–350 HV. These values are greater than IN718 under high temperature. The yield strength of HESA at 1000 °*C* was about 500 MPa. At a strain of 150 MPa under a temperature of about 980 °*C* , the HESAs showed excellent elevated temperature creep strength when compared with first generational superalloys. The creep strength and fatigue resistance of HESAs is due to the positive lattice misfit of the superalloy [72]. The raft which directionally coarsens the γ´ precipitate is corresponding with the stress axis which results in a sluggish motion of dislocation in the γ-γ´ precipitate interface, thus hindering the propagation of cracks initiated by fatigue and perpendicular to the same stress axis [73]. The HESAs showed promising thermal stability with a compact protection layer of Cr2O3, Al2O3 observed on the surface of the HESAs at elevated temperatures, while the densities of the HESAs ranged from 7.78–7.94 g/ cm3 as opposed traditional superalloys that range between 7.8–9.4 g/cm3 attributed to the high concentration of Cr, Fe, Ti and Al elements [42]. Other superalloys used for turbine engine applications are presented in **Table 1**.

#### **2.2 Protection of superalloys in gas turbine applications**

Wee, Do [81] described in a review of the mechanical thermal properties of superalloys and the authors stated that superalloys are required to perform excellently under severe thermal and mechanical stresses. The turbine engine may experience failure attributed to linear and cyclic movements of the pistons, connecting rods, rotors and

**101**

*Recent Advances of High Entropy Alloys: High Entropy Superalloys*

**Structure**

γ γ and ´ Room-

γ γ, ´andTCP High strength

γ, FCC L12 High Strength-

temperature strength, hightemperature strength, Creep, Wear and oxidation resistance

at elevated temperatures, corrosion -resistant, thermal shock resistant, easy to machine

toughness and fatigue strength, corrosionresistant,

**Advantages Disadvantages Ref**

Difficult to machine, poor service performance, susceptible to defects, hot corrosion degradation

Low strength compared to other superalloys,

Low adhesive, high friction coefficient, low ductility

[74–76]

[77, 78]

[79, 80]

shafts majorly affecting the cascade fluids on the surface of the superalloy [82]. For turbine applications, superalloys comprise elements which are meant for elevated temperature strength required for efficiency [83]. However, these alloying elements may also adversely impact the superalloy's resistance when in this severe environmental conditions over some time. Therefore, there may be a need for additional protection of the superalloy through surface treatments [84]. There are several laser surface modification treatments, namely; laser surface hardening, laser surface heat treatment, Laser alloying, laser shot peening, laser surface dispersing and laser coatings and cladding [85]. Laser coatings enable the superalloy to be resistant to its environment, have microstructural stability and enhance its thermal, physical and mechanical properties [28, 86]. The coatings available can be classified as; overlay coatings, diffusion coatings and ceramic barriers [87, 88]. The deposition of Al from a different external source and diffusing it into the base superalloy to for an external layer is called aluminide or diffusion coating. Bonding an oxidation-resistant alloy which is weak but highly effective on a superalloy to enable surface protection and stability is called overlay cladding, while ceramic barriers are ceramic coatings attached to the surface of a superalloy [89].

Technological advancements in surface engineering have replaced conventional methods of surface treatments with laser surface modification (LSM) techniques. The use of lasers in LSM has been reported to produce wear, corrosion, fracture and fatigue resistant HEAs coatings. This is attributed to the energy absorption and rapid solidification of the deposition process, which promotes fine microstructures

According to Wu et al. [90] used laser surface alloying to study the phase evolution and cavitation erosion-corrosion behavior of a HEA coating in distilled water and NaCl solution. The study showed that the alloy's cavitation erosion resistance

**3. Advances in manufacturing technology**

necessary for surface modification.

*DOI: http://dx.doi.org/10.5772/intechopen.96661*

**SuperAlloy Composition Phase** 

IN800H (32Ni-21Cr-1.5Mn-01Si-0.3Ti-0.3Al-01Cbal Fe, wt.%)

2Ti-5.5 W-2.5Ta-0.1B, at%)

(Ti-48Al-2Cr-2Nb)

*Superalloys used for turbine engine applications.*

Cobalt-Based (Co-30Ni-11Al-

Iron-Based (Incoloy 800H, Type A-286 alloy, IN903)

Titaniumbased (TiAl, Ti6Al4V)

**Table 1.**


*Recent Advances of High Entropy Alloys: High Entropy Superalloys DOI: http://dx.doi.org/10.5772/intechopen.96661*

#### **Table 1.**

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

and tensile strength at room temperatures than nickel superalloy.

as opposed traditional superalloys that range between 7.8–9.4 g/cm3

for turbine engine applications are presented in **Table 1**.

**2.2 Protection of superalloys in gas turbine applications**

to the high concentration of Cr, Fe, Ti and Al elements [42]. Other superalloys used

Wee, Do [81] described in a review of the mechanical thermal properties of superalloys and the authors stated that superalloys are required to perform excellently under severe thermal and mechanical stresses. The turbine engine may experience failure attributed to linear and cyclic movements of the pistons, connecting rods, rotors and

attributed

examined the oxidation behavior of a HESA in O2 environments. The Ni2FeCoCrAl0.5 HESA oxidation kinetics at 900 °*C* followed a parabolic-rate law forming scales which was dependent on the oxygen pressure. The results showed that the oxidation rates increased with an increase in oxygen pressure however, the kinetics of massloss was observed. Shafiee, Nili-Ahmadabadi [67] designed a wrought HESA using Phacomp and CALPHAD technique. The reports showed that the superalloy comprised of γ´ nanosized precipitates with lower densities, excellent workability and high thermal stability than Inconel 718 alloy and Waspaloy. Saito, Chen [68] discussed the influence of heat treatments on HESA microstructural evolution and results showed the cast HESA had coarsened γ´ precipitates attributed to microsegregation which decreased the solidus making the γ´ solvus unclear. Finally, Zhang, Huo [69] prepared cast Ni48-xCo18 Fe9.3Al9.7Cr10.5Ti4.5Mox HESA to investigate the mechanical and microstructural properties of the superalloy and they concluded that HESAs exhibits good compressive strength at elevated temperature, elongation

HESAs are stable at elevated temperatures compared with commercial Rene' N6, Udimet 700 and Hastelloy X superalloys, this attributed to their sluggish diffusion and high entropy effect. At high temperatures, Nickel-based superalloys form intermetallic topological closed-packed (TCP) phases rich in Fe-Cr because of the high iron content in less than 100 h at 900 °*C* and this TCP phases formed is detrimental to the stability of superalloys at high temperatures [70]. However, at 900 °*C* in more than 200 h, there was no TCP phase observed in Ni40.7Al7.8Cr12.2 Fe11.58Co20.6Ti7.2 HESA and the γ-γ´ microstructure of the superalloy remained stable after isothermal aging for 500 h at 1050 °*C* [57]. The elevated temperature strength of HESAs has been reported to be higher than that of IN 617. This can be attributed to enhancing the APB energy, increasing the lattice distortion and/or adding refractory Ta and W elements in high concentration to the compositional system. Yeh and Tsao [71] did a thorough analysis of HESAs with the elemental composition of Fe, C, Al, Mo, Cr, Ti, Ni, Co, Ta, W and Nb. The siderophile element was Nickel while the strengthening element was Nb and C. The authors reported that the HESA's microstructure comprised an FCC, L12 crystal structure and γ´ phase while the superalloys had hardness values of 400–470 HV at room temperature. At elevated temperatures, the HESAs hardness values recorded were between 300–350 HV. These values are greater than IN718 under high temperature. The yield strength of HESA at 1000 °*C* was about 500 MPa. At a strain of 150 MPa under a temperature of about 980 °*C* , the HESAs showed excellent elevated temperature creep strength when compared with first generational superalloys. The creep strength and fatigue resistance of HESAs is due to the positive lattice misfit of the superalloy [72]. The raft which directionally coarsens the γ´ precipitate is corresponding with the stress axis which results in a sluggish motion of dislocation in the γ-γ´ precipitate interface, thus hindering the propagation of cracks initiated by fatigue and perpendicular to the same stress axis [73]. The HESAs showed promising thermal stability with a compact protection layer of Cr2O3, Al2O3 observed on the surface of the HESAs at elevated temperatures, while the densities of the HESAs ranged from 7.78–7.94 g/

**100**

cm3

*Superalloys used for turbine engine applications.*

shafts majorly affecting the cascade fluids on the surface of the superalloy [82]. For turbine applications, superalloys comprise elements which are meant for elevated temperature strength required for efficiency [83]. However, these alloying elements may also adversely impact the superalloy's resistance when in this severe environmental conditions over some time. Therefore, there may be a need for additional protection of the superalloy through surface treatments [84]. There are several laser surface modification treatments, namely; laser surface hardening, laser surface heat treatment, Laser alloying, laser shot peening, laser surface dispersing and laser coatings and cladding [85]. Laser coatings enable the superalloy to be resistant to its environment, have microstructural stability and enhance its thermal, physical and mechanical properties [28, 86]. The coatings available can be classified as; overlay coatings, diffusion coatings and ceramic barriers [87, 88]. The deposition of Al from a different external source and diffusing it into the base superalloy to for an external layer is called aluminide or diffusion coating. Bonding an oxidation-resistant alloy which is weak but highly effective on a superalloy to enable surface protection and stability is called overlay cladding, while ceramic barriers are ceramic coatings attached to the surface of a superalloy [89].

#### **3. Advances in manufacturing technology**

Technological advancements in surface engineering have replaced conventional methods of surface treatments with laser surface modification (LSM) techniques. The use of lasers in LSM has been reported to produce wear, corrosion, fracture and fatigue resistant HEAs coatings. This is attributed to the energy absorption and rapid solidification of the deposition process, which promotes fine microstructures necessary for surface modification.

According to Wu et al. [90] used laser surface alloying to study the phase evolution and cavitation erosion-corrosion behavior of a HEA coating in distilled water and NaCl solution. The study showed that the alloy's cavitation erosion resistance

was enhanced in distilled water but not in NaCl solution due to the corrosion. Zhang et al. [91] fabricated HEA by laser surface alloying to examine the properties of the alloy and they reported that the microhardness property of the coating was thrice the number of the substrate and there were improvements in the wear resistance of the alloy. Huang et al. [92] investigated an equimolar HEA on a titanium alloy substrate using LSM and the results also showed enhancements in the wear resistance of the alloy attributed to the manufacturing route which contributed to the formation of the phases observed in the BCC matrix. Nahmany et al. [93] used an electron beam surface remelting technique to modify two-five component HEAs, and the authors inspected the influence of these surface modification processes on the properties of the alloys. The authors observed a significant increase in the microhardness due to the rapid solidification and cooling process associated with the fabrication technique. From literature, it can be deduced that LSM classified into laser surface remelting, surface amorphisation, laser transformation hardening, shock hardening, laser cladding, laser surface alloying and laser shock peening using different types of lasers can be used to enhance the properties of HEAs [94].

## **4. Laser surface modification**

Laser application in surface modification techniques can be dated back to Albert Einstein who was the first scientist to conceive a stimulated emission in 1917 which today makes lasers applicable [95].

A laser is an abbreviation for "light amplification by stimulated emission of radiation". It is classified into CO2 and Excimer gaseous lasers, Nd:YAG Solid-state Lasers, Liquid Dye lasers and Yb-doped Fiber. These lasers consist of an optical resonator, a pumping energy outlet and a gain medium. The gain medium is located inside the optical resonator which amplifies a light beam using external energy supplied by the pumping energy outlet. They are classified into dyes, semiconductors or fibers, solid and gaseous states.

Lasers are generally characterized by the ability to avoid divergence in a longdistance, possession of an increased level of energy and monochromaticity [96].


**103**

*Recent Advances of High Entropy Alloys: High Entropy Superalloys*

are low beam quality, the severity of maintenance and high running cost [100]. Sharma et al. [101] reported using an excimer laser with a wavelength of 248 nm for target ablation during the creation of epitaxial single crystal high entropy ABO3 perovskite thin films. The authors described how this process was significant in understanding different bonding environments to develop macroscopic responses

c.Nd:Yag which is an acronym for neodymium-doped yttrium aluminum garnet laser is a 1064 nm solid-state laser made up of an active ion and a host from either glass or solid crystalline. It is one of the widely used for the surface modification of HEAs attributed to the ability of its light beam to be transported by flexible optical fibers, consequently increasing its delivery efficiency and compactness [102, 103]. It is also not limited by its mode of transport, which can occur both in pulse and continuous modes. Recently, diode lasers have been substituted for Xenon flash lamps as the pump source to improve the quality of the beam. More so, Nd: YVO4 is a recent substitute for the Nd:Yag laser due to its wider band absorption, high efficiency and lower operating threshold [104].

d.A Fiber Laser is about 848 nm in wavelength with a rare earth doped fiber used for high power generation due to its increased level of efficiency. The Yb-Doped fiber lasers have excellent electrical-to-optical efficiency with system compactness and high-quality beam. Neodymium, holmium, thulium, dysprosium, erbium and praseodymium are other rare earth elements used as a gain medium in fiber lasers. Fiber lasers are usually pumped with laser diodes; however, they are limited by their light propagation through the optical fiber which greatly influences the guiding medium compared with when the propagation occurs through the air inside the fiber. More so, other factors like the Kerr lens and Raman effects limit the performance of the laser, therefore, optical fibers with polarization maintenance are strongly recommended as the gain medium [105]. Fan et al. [106] examined the influence of fiber laser welding on the mechanical and microstructural proprieties in addition to the solute segregation of a high entropy alloy. The authors reported that the alloy showed dendritic structures with those fabricated using Nd:Yag laser and they observed copper's segregation to the interdendritic region were also attributed to its smaller bonding energies with other elements in the HEA composition, conversely; the alloy

showed better hardness and strength compared with the Nd:Yag.

grain refinement and corrosion resistance observed.

in research, hence, should be further explored.

e.Organic liquid dye lasers use organic dyes as the gain medium. These liquid dye with about 50–100 nm compared to solids have a higher density of atoms and they are evenly distributed. These lasers with wide bandwidth are replaceable and are transferred from very intricate regions which are sometimes used as solutes in considerable solvents to develop gain mediums [107]. Coumarin, pyrromethene, exalite, pyridine, styryl and fluorescein are dyes used for pulsed or tunable lasers. Nevertheless, these lasers are limited in applications because they require a large volume of organic solvents for efficiency. Xu et al. [108] used a laser stimulated fluorescence equipment consisting of an organic liquid dye to fabricate a HEA and study the performance of the coatings then the influence of aluminum on the properties of the alloy. The authors stated that the laser technology and the aluminum content enabled the phase transitions,

f. Other types of lasers are; semiconductor lasers, hybrid laser arc welding and free-electron lasers and the fabrications of HEAs using these lasers are limited

driven by complex exchange interactions and electron–phonon channels.

*DOI: http://dx.doi.org/10.5772/intechopen.96661*

#### *Recent Advances of High Entropy Alloys: High Entropy Superalloys DOI: http://dx.doi.org/10.5772/intechopen.96661*

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

was enhanced in distilled water but not in NaCl solution due to the corrosion. Zhang et al. [91] fabricated HEA by laser surface alloying to examine the properties of the alloy and they reported that the microhardness property of the coating was thrice the number of the substrate and there were improvements in the wear resistance of the alloy. Huang et al. [92] investigated an equimolar HEA on a titanium alloy substrate using LSM and the results also showed enhancements in the wear resistance of the alloy attributed to the manufacturing route which contributed to the formation of the phases observed in the BCC matrix. Nahmany et al. [93] used an electron beam surface remelting technique to modify two-five component HEAs, and the authors inspected the influence of these surface modification processes on the properties of the alloys. The authors observed a significant increase in the microhardness due to the rapid solidification and cooling process associated with the fabrication technique. From literature, it can be deduced that LSM classified into laser surface remelting, surface amorphisation, laser transformation hardening, shock hardening, laser cladding, laser surface alloying and laser shock peening using different types of lasers can be used to enhance the properties of HEAs [94].

Laser application in surface modification techniques can be dated back to Albert Einstein who was the first scientist to conceive a stimulated emission in 1917 which

A laser is an abbreviation for "light amplification by stimulated emission of radiation". It is classified into CO2 and Excimer gaseous lasers, Nd:YAG Solid-state Lasers, Liquid Dye lasers and Yb-doped Fiber. These lasers consist of an optical resonator, a pumping energy outlet and a gain medium. The gain medium is located inside the optical resonator which amplifies a light beam using external energy supplied by the pumping energy outlet. They are classified into dyes, semiconductors

Lasers are generally characterized by the ability to avoid divergence in a longdistance, possession of an increased level of energy and monochromaticity [96].

a.The CO2 laser comprises an electric pump, discharge tube, CO2 gas for the gain medium and optics such as silver or gold mirrors, zinc selenide lens and finally a window as the optical resonator. Although the Helium-Neon laser was the first gas laser developed in a Bell telephone laboratory, still, the CO2 is the most widely used gas laser for its high emission wavelength between 9–11 μm which offers very high power for surface modification. The process experiences low light absorption in the infrared regions, reduced optical fiber delivery, instability in the output power attributed to the contraction of the laser

structure and thermal expansion when pumping the gas by an AC or DC which sometimes limits its application. Zhang et al. [97] reported fabricating HEAs with CO2 laser, and the alloy had fine microstructural morphologies and higher mechanical properties. While Zheng et al. [98] mentioned that the HEAs coating fabricated using gas lasers had cellular crystals with dispersion precipitates

b.Excimer lasers, on the other hand, is a mixture of noble gases like helium buffer gas, xenon, argon and a chloride or fluoride halogen. Excimer which is about 248 nm is also known as excited dimers which are pumped using a pulsed electrical discharge for the production of nanosecond pulses in an ultraviolet region, for that reason; it can only be operated in a pulsed mode. Other limitations of this laser

although the hardness values were reported to be high [99].

**4. Laser surface modification**

today makes lasers applicable [95].

or fibers, solid and gaseous states.

**102**

are low beam quality, the severity of maintenance and high running cost [100]. Sharma et al. [101] reported using an excimer laser with a wavelength of 248 nm for target ablation during the creation of epitaxial single crystal high entropy ABO3 perovskite thin films. The authors described how this process was significant in understanding different bonding environments to develop macroscopic responses driven by complex exchange interactions and electron–phonon channels.


### **4.1 Laser surface melting (LSM)**

This type of surface modification is used for material hardening, electrochemical and tribological resistance and reduction in porosity. An increased rate of heat transfer occurs during the interaction between the substrate and the melted HEAs coating surface, especially during solidification. The rapid solidification and cooling rates invariably produce fine microstructures which also enhances the surface properties of the alloys. Chen et al. [109] used LSM on HEAs and they mentioned that the surface modification process increased the electrochemical and mechanical properties of the alloys. Ochelik et al. [105] found that the solidification rate influences the phases formed using LSM. The fast solidification rates promoted the BCC phase observed which was also responsible for the improved hardness properties of the alloys. Cai et al. [110] also reported observing a BCC solid solution phase and improved microhardness properties after using LSM. The as-remelted HEAs coatings had low wear mass loss showing an improvement in the wear resistance.

#### **4.2 Laser transformation hardening (LTH)**

The LTH heats the HEAs coating or films at a very high temperature with an unfocused beam, and then rapid cooling occurs immediately without letting equilibrium phases to form by quenching, as a result, generating very low thermal distortion. This method uses a diode laser or CO2 to increase the surface properties of the HEAs [111].

#### **4.3 Laser surface alloying (LSA)**

This involves the direct injection or pre-placement of additional elements unto the surface of the substrate by a laser source. Rapid solidification occurs with the substrate maintaining its temperature while acting as a heat sink, still the composition of the surface changes [112]. Therefore, re-solidification and rapid quenching follow due to the temperature difference between the surface of the substrate and the treated surface zone. Zhang et al. [113] fabricated HEA coatings by LSA, and the HEA coating had a BCC solid solution phase with improved mechanical and corrosion properties. Jiang et al. [114] fabricated HEAs on a 304 stainless steel substrate and they stated that although the alloy had FCC and BCC phases, the BCC phase was more predominant. The authors also recorded a substantial increase in the hardness with good wear-resistant properties.

#### **4.4 Laser glazing**

This method produces a nanocrystalline layer or thin amorphous layer on the surface of the substrate, energy is absorbed into the surface which melts the HEAs coating/films to a certain depth with a laser beam and rapid solidification occurs. This process is achieved using a high power density at a short period enough to create the amorphous structure needed for surface modification [115].

#### **5. Conclusion**

High-temperature properties of materials used for turbine engine applications are important for the reduction of fuel consumption, operating costs and pollution. Nickel-based superalloys are widely used due to its strength, resistance to degradation in oxidizing environments, toughness and density. However, Nickel superalloy

**105**

**Author details**

**Acknowledgements**

Modupeola Dada1

Afolabi Ayodeji<sup>2</sup>

Sisa Pityana<sup>3</sup>

\*, Patricia Popoola<sup>2</sup>

2 Tshwane University of Technology, Pretoria, South Africa

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

, Olufemi Aramide2

Technology, Pretoria, South Africa

Technology, Pretoria, South Africa

4 University of Lagos, Akoka, Lagos

provided the original work is properly cited.

, Ntombizodwa Mathe3

, Nicholus Malatji2

1 Chemical, Metallurgical and Materials Engineering, Tshwane University of

3 Council for Scientific and Industrial Research and Tshwane University of

© 2021 The Author(s). Licensee IntechOpen. 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,

, Samson Adeosun4

and

, Thabo Lengopeng2

, a positive lattice misfit and high yield

,

*Recent Advances of High Entropy Alloys: High Entropy Superalloys*

is not stable at elevated temperatures having a maximum service temperature of 649 °*C* , the superalloy at room temperature has a negative lattice misfit, poor thermal conductivity and difficult to machine. High Entropy Superalloys, with similar γ and γ´ phases as the Nickel-based superalloys, shows high tensile strength than Inconel 617 and Alloy 800 H. The superalloy exhibits good oxidation resis-

strength compared to traditional nickel superalloys. Controlling the elemental compositional partitioning between the γ-γ´ in high entropy superalloys makes the thermal stability higher than conventional nickel superalloys and equimolar or near equimolar high entropy alloys. Therefore γ´ precipitate strengthening of solid solution high entropy alloys to form High Entropy Superalloys is currently the most promising material for turbine engine applications. Laser surface modification

The authors will like to appreciate the National Laser Center (Laser Enabled Manufacturing Resource Group); Council for Scientific and Research (CSIR) and the Surface Engineering Research Laboratory; Tshwane University of Technology, Pretoria, South Africa for their scientific and technical support during this research.

treatments can be used as a protective mechanism for Superalloys.

*DOI: http://dx.doi.org/10.5772/intechopen.96661*

tance; have lower densities below 8 g/cm3

*Recent Advances of High Entropy Alloys: High Entropy Superalloys DOI: http://dx.doi.org/10.5772/intechopen.96661*

is not stable at elevated temperatures having a maximum service temperature of 649 °*C* , the superalloy at room temperature has a negative lattice misfit, poor

thermal conductivity and difficult to machine. High Entropy Superalloys, with similar γ and γ´ phases as the Nickel-based superalloys, shows high tensile strength than Inconel 617 and Alloy 800 H. The superalloy exhibits good oxidation resistance; have lower densities below 8 g/cm3 , a positive lattice misfit and high yield strength compared to traditional nickel superalloys. Controlling the elemental compositional partitioning between the γ-γ´ in high entropy superalloys makes the thermal stability higher than conventional nickel superalloys and equimolar or near equimolar high entropy alloys. Therefore γ´ precipitate strengthening of solid solution high entropy alloys to form High Entropy Superalloys is currently the most promising material for turbine engine applications. Laser surface modification treatments can be used as a protective mechanism for Superalloys.

## **Acknowledgements**

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

This type of surface modification is used for material hardening, electrochemical and tribological resistance and reduction in porosity. An increased rate of heat transfer occurs during the interaction between the substrate and the melted HEAs coating surface, especially during solidification. The rapid solidification and cooling rates invariably produce fine microstructures which also enhances the surface properties of the alloys. Chen et al. [109] used LSM on HEAs and they mentioned that the surface modification process increased the electrochemical and mechanical properties of the alloys. Ochelik et al. [105] found that the solidification rate influences the phases formed using LSM. The fast solidification rates promoted the BCC phase observed which was also responsible for the improved hardness properties of the alloys. Cai et al. [110] also reported observing a BCC solid solution phase and improved microhardness properties after using LSM. The as-remelted HEAs coatings had low wear mass loss showing an improvement in the wear resistance.

The LTH heats the HEAs coating or films at a very high temperature with an unfocused beam, and then rapid cooling occurs immediately without letting equilibrium phases to form by quenching, as a result, generating very low thermal distortion. This method uses a diode laser or CO2 to increase the surface properties

This involves the direct injection or pre-placement of additional elements unto the surface of the substrate by a laser source. Rapid solidification occurs with the substrate maintaining its temperature while acting as a heat sink, still the composition of the surface changes [112]. Therefore, re-solidification and rapid quenching follow due to the temperature difference between the surface of the substrate and the treated surface zone. Zhang et al. [113] fabricated HEA coatings by LSA, and the HEA coating had a BCC solid solution phase with improved mechanical and corrosion properties. Jiang et al. [114] fabricated HEAs on a 304 stainless steel substrate and they stated that although the alloy had FCC and BCC phases, the BCC phase was more predominant. The authors also recorded a substantial increase in the hardness

This method produces a nanocrystalline layer or thin amorphous layer on the surface of the substrate, energy is absorbed into the surface which melts the HEAs coating/films to a certain depth with a laser beam and rapid solidification occurs. This process is achieved using a high power density at a short period enough to cre-

High-temperature properties of materials used for turbine engine applications are important for the reduction of fuel consumption, operating costs and pollution. Nickel-based superalloys are widely used due to its strength, resistance to degradation in oxidizing environments, toughness and density. However, Nickel superalloy

ate the amorphous structure needed for surface modification [115].

**4.1 Laser surface melting (LSM)**

**4.2 Laser transformation hardening (LTH)**

of the HEAs [111].

**4.4 Laser glazing**

**5. Conclusion**

**4.3 Laser surface alloying (LSA)**

with good wear-resistant properties.

**104**

The authors will like to appreciate the National Laser Center (Laser Enabled Manufacturing Resource Group); Council for Scientific and Research (CSIR) and the Surface Engineering Research Laboratory; Tshwane University of Technology, Pretoria, South Africa for their scientific and technical support during this research.

## **Author details**

Modupeola Dada1 \*, Patricia Popoola<sup>2</sup> , Ntombizodwa Mathe3 , Samson Adeosun4 , Sisa Pityana<sup>3</sup> , Olufemi Aramide2 , Nicholus Malatji2 , Thabo Lengopeng2 and Afolabi Ayodeji<sup>2</sup>

1 Chemical, Metallurgical and Materials Engineering, Tshwane University of Technology, Pretoria, South Africa

2 Tshwane University of Technology, Pretoria, South Africa

3 Council for Scientific and Industrial Research and Tshwane University of Technology, Pretoria, South Africa

4 University of Lagos, Akoka, Lagos

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

© 2021 The Author(s). Licensee IntechOpen. 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.

## **References**

[1] Boyce, M.P., *Gas turbine engineering handbook*. 2011: Elsevier.

[2] Kaygusuz, K., *Sustainable development of hydroelectric power.* Energy sources, 2002. **24**(9): p. 803-815.

[3] Langston, L.S., G. Opdyke, and E. Dykewood, *Introduction to gas turbines for non-engineers.* Global Gas Turbine News, 1997. **37**(2): p. 1-9.

[4] Heywood, J.B., *Internal combustion engine fundamentals*. 2018: McGraw-Hill Education.

[5] Bell, M. and T. Partridge, *Thermodynamic design of a reciprocating Joule cycle engine.* Proceedings of the Institution of Mechanical Engineers, Part A: Journal of Power and Energy, 2003. **217**(3): p. 239-246.

[6] Cheng, C.-Y. and C.O.-K. Chen, *Power optimization of an irreversible Brayton heat engine.* Energy sources, 1997. **19**(5): p. 461-474.

[7] Viteri, F. and R.E. Anderson, *Semiclosed brayton cycle gas turbine power systems*. 2003, Google Patents.

[8] Walsh, P.P. and P. Fletcher, *Gas turbine engine*. 2005, Google Patents.

[9] Li, Y. and P. Nilkitsaranont, *Gas turbine performance prognostic for condition-based maintenance.* Applied energy, 2009. **86**(10): p. 2152-2161.

[10] Reed, R.C., *The superalloys: fundamentals and applications*. 2008: Cambridge university press.

[11] Liu, C., et al., *Improved castability of directionally solidified, Ni-based superalloy by the liquid metal cooling process.* Metallurgical and Materials Transactions A, 2012. **43**(2): p. 405-409.

[12] Wang, R.-Z., et al., *Creep-fatigue life prediction and interaction diagram in nickel-based GH4169 superalloy at 650 C based on cycle-by-cycle concept.* International Journal of Fatigue, 2017. **97**: p. 114-123.

[13] Ganji, D.K. and G. Rajyalakshmi, *Influence of Alloying Compositions on the Properties of Nickel-Based Superalloys: A Review*, in *Recent Advances in Mechanical Engineering*. 2020, Springer. p. 537-555.

[14] Akca, E. and A. Gürsel, *A review on superalloys and IN718 nickel-based INCONEL superalloy.* Periodicals of engineering and natural sciences, 2015. **3**(1).

[15] Miracle, D.B., et al., *Refractory high entropy superalloys (RSAs).* Scripta Materialia, 2020. **187**: p. 445-452.

[16] Tsao, T.-K., et al., *High temperature oxidation and corrosion properties of high entropy superalloys.* Entropy, 2016. **18**(2): p. 62.

[17] Yeh, A., et al., *Developing new type of high temperature alloys–high entropy superalloys.* International Journal of Metallurgical & Materials Engineering, 2015. **2015**.

[18] Pint, B.A., K. Unocic, and S. Dryepondt. *Oxidation of superalloys in extreme environments*. in *7th International Symposium on Superalloy*. 2010.

[19] Liu, X., et al., *Effects of Nb and W additions on the microstructures and mechanical properties of novel γ/γ'Co-V-Ti-Based superalloys.* Metals, 2018. **8**(7): p. 563.

[20] Locq, D., et al., *Development of new PM superalloys for high temperature applications.* Intermetallics and superalloys, 2000. **10**: p. 52-57.

**107**

*Recent Advances of High Entropy Alloys: High Entropy Superalloys*

[30] Reed, R., T. Tao, and N. Warnken, *Alloys-by-design: application to nickelbased single crystal superalloys.* Acta Materialia, 2009. **57**(19): p. 5898-5913.

[31] Kennedy, R., *ALLVAC® 718PLUS™,* 

*superalloy for the next forty years.* Superalloys, 2005. **718**(706): p. 1-14.

Superalloys, 2012. **911919**.

Science & Business Media.

a, 1998. **29**(12): p. 3069-3079.

[36] Durand-Charre, M., *The microstructure of superalloys*. 2017:

[37] Sikka, V., et al., *Advances in processing of Ni3Al-based intermetallics and applications.* Intermetallics, 2000.

[38] Frank, R.B., C.G. Roberts, and J. Zhang. *Effect of nickel content on delta solvus temperature and mechanical properties of alloy 718*. in *7th international symposium on* 

[39] Choudhury, I. and M. El-Baradie, *Machinability of nickel-base super alloys: a general review.* Journal of Materials Processing Technology, 1998. **77**(1-3):

**8**(9-11): p. 1329-1337.

*superalloy*. 2010.

p. 278-284.

3712-3721.

Routledge.

[32] Devaux, A., et al., *AD730TM-A new nickel-based superalloy for high temperature engine rotative parts.* TMS

[33] Thellaputta, G.R., P.S. Chandra, and C. Rao, *Machinability of nickel based superalloys: a review.* Materials Today: Proceedings, 2017. **4**(2): p.

[34] Flower, H.M., *High performance materials in aerospace*. 2012: Springer

[35] Lehockey, E., G. Palumbo, and P. Lin, *Improving the weldability and service performance of nickel-and iron-based superalloys by grain boundary engineering.* Metallurgical and materials transactions

*DOI: http://dx.doi.org/10.5772/intechopen.96661*

*Manufacturing Science and Engineering Conference*. 2018. American Society of

[22] Walston, S., et al., *Joint development of a fourth generation single crystal* 

[23] Perrut, M., et al., *High temperature materials for aerospace applications: Ni-based superalloys and γ-TiAl alloys.* Comptes Rendus Physique, 2018. **19**(8):

[24] Gessinger, G.H. and M. Bomford, *Powder metallurgy of superalloys.* International Metallurgical Reviews,

*manufacturing of aircraft engine disks by roll forming and hot die forging.* Journal of Materials Processing Technology, 2003.

[25] Bewlay, B., et al., *Net-shape* 

[26] Lavella, M., T. Berruti, and E. Bosco, *Residual stress analysis in Inconel 718 milled turbine disk.* International Journal of Machining and Machinability of Materials, 2008. **4**(2-3): p. 181-194.

[27] Groh, J., et al. *Development of a new cast and wrought alloy (René 65) for high temperature disk applications*. in *Proceedings of the 8th International Symposium on Superalloy 718 and Derivatives*. 2014. John Wiley & Sons.

[28] Pollock, T.M. and S. Tin, *Nickelbased superalloys for advanced turbine engines: chemistry, microstructure and properties.* Journal of propulsion and power, 2006. **22**(2): p. 361-374.

[29] Ezugwu, E., J. Bonney, and Y. Yamane, *An overview of the machinability of aeroengine alloys.* Journal of materials processing technology, 2003. **134**(2): p. 233-253.

[21] Graybill, B., et al. *Additive Manufacturing of nickel-based superalloys*. in *International* 

Mechanical Engineers.

*superalloy.* 2004.

p. 657-671.

1974. **19**(1): p. 51-76.

**135**(2-3): p. 324-329.

*Recent Advances of High Entropy Alloys: High Entropy Superalloys DOI: http://dx.doi.org/10.5772/intechopen.96661*

[21] Graybill, B., et al. *Additive Manufacturing of nickel-based superalloys*. in *International Manufacturing Science and Engineering Conference*. 2018. American Society of Mechanical Engineers.

[22] Walston, S., et al., *Joint development of a fourth generation single crystal superalloy.* 2004.

[23] Perrut, M., et al., *High temperature materials for aerospace applications: Ni-based superalloys and γ-TiAl alloys.* Comptes Rendus Physique, 2018. **19**(8): p. 657-671.

[24] Gessinger, G.H. and M. Bomford, *Powder metallurgy of superalloys.* International Metallurgical Reviews, 1974. **19**(1): p. 51-76.

[25] Bewlay, B., et al., *Net-shape manufacturing of aircraft engine disks by roll forming and hot die forging.* Journal of Materials Processing Technology, 2003. **135**(2-3): p. 324-329.

[26] Lavella, M., T. Berruti, and E. Bosco, *Residual stress analysis in Inconel 718 milled turbine disk.* International Journal of Machining and Machinability of Materials, 2008. **4**(2-3): p. 181-194.

[27] Groh, J., et al. *Development of a new cast and wrought alloy (René 65) for high temperature disk applications*. in *Proceedings of the 8th International Symposium on Superalloy 718 and Derivatives*. 2014. John Wiley & Sons.

[28] Pollock, T.M. and S. Tin, *Nickelbased superalloys for advanced turbine engines: chemistry, microstructure and properties.* Journal of propulsion and power, 2006. **22**(2): p. 361-374.

[29] Ezugwu, E., J. Bonney, and Y. Yamane, *An overview of the machinability of aeroengine alloys.* Journal of materials processing technology, 2003. **134**(2): p. 233-253. [30] Reed, R., T. Tao, and N. Warnken, *Alloys-by-design: application to nickelbased single crystal superalloys.* Acta Materialia, 2009. **57**(19): p. 5898-5913.

[31] Kennedy, R., *ALLVAC® 718PLUS™, superalloy for the next forty years.* Superalloys, 2005. **718**(706): p. 1-14.

[32] Devaux, A., et al., *AD730TM-A new nickel-based superalloy for high temperature engine rotative parts.* TMS Superalloys, 2012. **911919**.

[33] Thellaputta, G.R., P.S. Chandra, and C. Rao, *Machinability of nickel based superalloys: a review.* Materials Today: Proceedings, 2017. **4**(2): p. 3712-3721.

[34] Flower, H.M., *High performance materials in aerospace*. 2012: Springer Science & Business Media.

[35] Lehockey, E., G. Palumbo, and P. Lin, *Improving the weldability and service performance of nickel-and iron-based superalloys by grain boundary engineering.* Metallurgical and materials transactions a, 1998. **29**(12): p. 3069-3079.

[36] Durand-Charre, M., *The microstructure of superalloys*. 2017: Routledge.

[37] Sikka, V., et al., *Advances in processing of Ni3Al-based intermetallics and applications.* Intermetallics, 2000. **8**(9-11): p. 1329-1337.

[38] Frank, R.B., C.G. Roberts, and J. Zhang. *Effect of nickel content on delta solvus temperature and mechanical properties of alloy 718*. in *7th international symposium on superalloy*. 2010.

[39] Choudhury, I. and M. El-Baradie, *Machinability of nickel-base super alloys: a general review.* Journal of Materials Processing Technology, 1998. **77**(1-3): p. 278-284.

**106**

405-409.

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

[12] Wang, R.-Z., et al., *Creep-fatigue life prediction and interaction diagram in nickel-based GH4169 superalloy at 650 C based on cycle-by-cycle concept.* International Journal of Fatigue, 2017.

[13] Ganji, D.K. and G. Rajyalakshmi, *Influence of Alloying Compositions on the Properties of Nickel-Based Superalloys: A Review*, in *Recent Advances in Mechanical Engineering*. 2020, Springer. p. 537-555.

[14] Akca, E. and A. Gürsel, *A review on superalloys and IN718 nickel-based INCONEL superalloy.* Periodicals of engineering and natural sciences,

[15] Miracle, D.B., et al., *Refractory high entropy superalloys (RSAs).* Scripta Materialia, 2020. **187**: p. 445-452.

[16] Tsao, T.-K., et al., *High temperature oxidation and corrosion properties of high entropy superalloys.* Entropy, 2016.

[17] Yeh, A., et al., *Developing new type of high temperature alloys–high entropy superalloys.* International Journal of Metallurgical & Materials Engineering,

[18] Pint, B.A., K. Unocic, and S. Dryepondt. *Oxidation of* 

*superalloys in extreme environments*. in *7th International Symposium on* 

[19] Liu, X., et al., *Effects of Nb and W additions on the microstructures and mechanical properties of novel γ/γ'Co-V-Ti-Based superalloys.* Metals, 2018.

[20] Locq, D., et al., *Development of new PM superalloys for high temperature* 

*applications.* Intermetallics and superalloys, 2000. **10**: p. 52-57.

**97**: p. 114-123.

2015. **3**(1).

**18**(2): p. 62.

2015. **2015**.

*Superalloy*. 2010.

**8**(7): p. 563.

[1] Boyce, M.P., *Gas turbine engineering* 

[3] Langston, L.S., G. Opdyke, and E. Dykewood, *Introduction to gas turbines for non-engineers.* Global Gas Turbine

[4] Heywood, J.B., *Internal combustion engine fundamentals*. 2018: McGraw-Hill

*Thermodynamic design of a reciprocating Joule cycle engine.* Proceedings of the Institution of Mechanical Engineers, Part A: Journal of Power and Energy,

[6] Cheng, C.-Y. and C.O.-K. Chen, *Power optimization of an irreversible Brayton heat engine.* Energy sources,

[7] Viteri, F. and R.E. Anderson, *Semiclosed brayton cycle gas turbine power systems*. 2003, Google Patents.

[8] Walsh, P.P. and P. Fletcher, *Gas turbine engine*. 2005, Google Patents.

[9] Li, Y. and P. Nilkitsaranont, *Gas turbine performance prognostic for condition-based maintenance.* Applied energy, 2009. **86**(10): p. 2152-2161.

[10] Reed, R.C., *The superalloys: fundamentals and applications*. 2008:

[11] Liu, C., et al., *Improved castability of directionally solidified, Ni-based superalloy by the liquid metal cooling process.* Metallurgical and Materials Transactions A, 2012. **43**(2): p.

Cambridge university press.

*handbook*. 2011: Elsevier.

**References**

News, 1997. **37**(2): p. 1-9.

[5] Bell, M. and T. Partridge,

2003. **217**(3): p. 239-246.

1997. **19**(5): p. 461-474.

Education.

[2] Kaygusuz, K., *Sustainable development of hydroelectric power.* Energy sources, 2002. **24**(9): p. 803-815. [40] Gao, M.C. and D.E. Alman, *Searching for next single-phase highentropy alloy compositions.* Entropy, 2013. **15**(10): p. 4504-4519.

[41] Chatterjee, P., V.M. Athawale, and S. Chakraborty, *Selection of materials using compromise ranking and outranking methods.* Materials & Design, 2009. **30**(10): p. 4043-4053.

[42] Chen, J., et al., *A review on fundamental of high entropy alloys with promising high–temperature properties.* Journal of Alloys and Compounds, 2018. **760**: p. 15-30.

[43] Hummel, R.E., *Understanding materials science: history, properties, applications*. 2004: Springer Science & Business Media.

[44] Cantor, B., et al., *Microstructural development in equiatomic multicomponent alloys.* Materials Science and Engineering: A, 2004. **375**: p. 213-218.

[45] Ye, Y., et al., *High-entropy alloy: challenges and prospects.* Materials Today, 2016. **19**(6): p. 349-362.

[46] Yeh, J.W., et al., *Nanostructured highentropy alloys with multiple principal elements: novel alloy design concepts and outcomes.* Advanced Engineering Materials, 2004. **6**(5): p. 299-303.

[47] Joseph, J., *Study of direct laser fabricated high entropy alloys*. 2016, Deakin University.

[48] Kuehl, R.O. and R. Kuehl, *Design of experiments: statistical principles of research design and analysis.* 2000.

[49] Wang, Y.P., B.S. Li, and H.Z. Fu, *Solid solution or intermetallics in a highentropy alloy.* Advanced engineering materials, 2009. **11**(8): p. 641-644.

[50] Gludovatz, B., et al., *A fractureresistant high-entropy alloy for cryogenic*  *applications.* Science, 2014. **345**(6201): p. 1153-1158.

[51] Mishra, R.K. and R.R. Shahi, *Magnetic characteristics of high entropy alloys.* Magnetism and magnetic materials. IntechOpen, Rijeka, 2018: p. 67-80.

[52] Senkov, O., et al., *Refractory highentropy alloys.* Intermetallics, 2010. **18**(9): p. 1758-1765.

[53] Hemphill, e.a., *Fatigue behavior of Al0. 5CoCrCuFeNi high entropy alloys.* Acta Materialia, 2012. **60**(16): p. 5723-5734.

[54] Tsao, T.-K., et al., *The high temperature tensile and creep behaviors of high entropy superalloy.* Scientific reports, 2017. **7**(1): p. 1-9.

[55] Senkov, O.N., et al., *Development of a refractory high entropy superalloy.* Entropy, 2016. **18**(3): p. 102.

[56] Chen, Y.-T., et al., *Hierarchical microstructure strengthening in a single crystal high entropy superalloy.* Scientific reports, 2020. **10**(1): p. 1-11.

[57] Tsao, T.-K., A.-C. Yeh, and H. Murakami, *The microstructure stability of precipitation strengthened medium to high entropy superalloys.* Metallurgical and Materials Transactions A, 2017. **48**(5): p. 2435-2442.

[58] Daoud, H., et al., *Microstructure and tensile behavior of Al 8 Co 17 Cr 17 Cu 8 Fe 17 Ni 33 (at.%) high-entropy alloy.* Jom, 2013. **65**(12): p. 1805-1814.

[59] He, J., et al., *A precipitation-hardened high-entropy alloy with outstanding tensile properties.* Acta Materialia, 2016. **102**: p. 187-196.

[60] Zhao, Y., et al., *Thermal stability and coarsening of coherent particles in a precipitation-hardened (NiCoFeCr)* 

**109**

p. 1-12.

p. 1600.

*Recent Advances of High Entropy Alloys: High Entropy Superalloys*

[69] Zhang, L., et al., *Microstructure and mechanical properties of precipitationhardened cast high-entropy superalloys.* Materials Science and Technology,

[70] Belan, J., *GCP and TCP phases presented in nickel-base superalloys.* Materials Today: Proceedings, 2016.

[71] Yeh, A.-c. and T.-K. Tsao, *Highentropy superalloy*. 2019, Google Patents.

[72] Zhang, J., et al., *The effect of lattice misfit on the dislocation motion in superalloys during high-temperature low-stress creep.* Acta Materialia, 2005.

[73] Mughrabi, H., *The importance of sign and magnitude of γ/γ' lattice misfit in superalloys—with special reference to the new γ'-hardened cobalt-base superalloys.* Acta materialia, 2014. **81**:

[74] Chawla, V., et al., *Corrosion Behavior of Nanostructured TiAlN and AlCrN Hard Coatings on Superfer 800H Superalloy in Simulated Marine Environment.* Journal of Minerals and Materials Characterization and Engineering, 2009. **8**(9): p. 693-700.

[75] Sidhu, T., et al., *Oxidation and hot corrosion resistance of HVOF WC-NiCrFeSiB coating on Ni-and Fe-based superalloys at 800 C.* Journal of Thermal Spray Technology, 2007. **16**(5-

[76] Moody, N., et al., *Temperature effects on hydrogen-induced crack growth susceptibility of iron-based superalloys.* Engineering Fracture Mechanics, 2001.

[77] Chung, D.-W., et al., *Effects of Cr on the properties of multicomponent cobalt-based superalloys with ultra high γ'volume fraction.* Journal of Alloys and

Compounds, 2020: p. 154790.

2020: p. 1-7.

**3**(4): p. 936-941.

**53**(17): p. 4623-4633.

p. 21-29.

6): p. 844-849.

**68**(6): p. 731-750.

*DOI: http://dx.doi.org/10.5772/intechopen.96661*

*94Ti2Al4 high-entropy alloy.* Acta Materialia, 2018. **147**: p. 184-194.

[61] Xiao, K., et al., *A scanning AC calorimetry technique for the analysis of nano-scale quantities of materials.* Review of Scientific Instruments, 2012. **83**(11):

[62] Wang, Z., et al., *Effect of coherent L12 nanoprecipitates on the tensile behavior of a fcc-based high-entropy alloy.* Materials Science and Engineering: A,

[63] Gwalani, B., et al., *Stability of ordered L12 and B2 precipitates in face centered cubic based high entropy alloys-Al0. 3CoFeCrNi and Al0. 3CuFeCrNi2.* Scripta Materialia, 2016.

[64] Borkar, T., et al., *A combinatorial assessment of AlxCrCuFeNi2 (0< x< 1.5) complex concentrated alloys: Microstructure, microhardness, and magnetic properties.* Acta Materialia,

[65] Li, Y., et al., *Microstructure and elevated-temperature mechanical properties of refractory AlMo0. 5NbTa0. 5TiZr High Entropy Alloy fabricated by powder metallurgy.* arXiv preprint

[66] Kai, W., et al., *The oxidation behavior of a Ni2FeCoCrAl0. 5 highentropy superalloy in O2-containing environments.* Corrosion Science, 2019.

[67] Shafiee, A., et al., *Development and microstructural characterization of a new wrought high entropy superalloy.* Metals and Materials International, 2019:

[68] Saito, T., et al., *Effect of Heat Treatments on the Microstructural Evolution of a Single Crystal High-*

*Entropy Superalloy.* Metals, 2020. **10**(12):

p. 114901.

2017. **696**: p. 503-510.

**123**: p. 130-134.

2016. **116**: p. 63-76.

arXiv:1801.00263, 2017.

**158**: p. 108093.

*Recent Advances of High Entropy Alloys: High Entropy Superalloys DOI: http://dx.doi.org/10.5772/intechopen.96661*

*94Ti2Al4 high-entropy alloy.* Acta Materialia, 2018. **147**: p. 184-194.

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

*applications.* Science, 2014. **345**(6201):

[51] Mishra, R.K. and R.R. Shahi, *Magnetic characteristics of high entropy alloys.* Magnetism and magnetic materials. IntechOpen, Rijeka, 2018:

[52] Senkov, O., et al., *Refractory highentropy alloys.* Intermetallics, 2010.

[53] Hemphill, e.a., *Fatigue behavior of Al0. 5CoCrCuFeNi high entropy alloys.* Acta Materialia, 2012. **60**(16): p.

[54] Tsao, T.-K., et al., *The high temperature tensile and creep behaviors of high entropy superalloy.* Scientific

[55] Senkov, O.N., et al., *Development of a refractory high entropy superalloy.*

[56] Chen, Y.-T., et al., *Hierarchical microstructure strengthening in a single crystal high entropy superalloy.* Scientific

[57] Tsao, T.-K., A.-C. Yeh, and H. Murakami, *The microstructure stability of precipitation strengthened medium to high entropy superalloys.* Metallurgical and Materials Transactions A, 2017. **48**(5): p.

[58] Daoud, H., et al., *Microstructure and tensile behavior of Al 8 Co 17 Cr 17 Cu 8 Fe 17 Ni 33 (at.%) high-entropy alloy.* Jom, 2013. **65**(12): p. 1805-1814.

[59] He, J., et al., *A precipitation-hardened high-entropy alloy with outstanding tensile properties.* Acta Materialia, 2016. **102**: p.

[60] Zhao, Y., et al., *Thermal stability and coarsening of coherent particles in a precipitation-hardened (NiCoFeCr)* 

reports, 2017. **7**(1): p. 1-9.

Entropy, 2016. **18**(3): p. 102.

reports, 2020. **10**(1): p. 1-11.

p. 1153-1158.

p. 67-80.

5723-5734.

2435-2442.

187-196.

**18**(9): p. 1758-1765.

[40] Gao, M.C. and D.E. Alman, *Searching for next single-phase highentropy alloy compositions.* Entropy, 2013.

[41] Chatterjee, P., V.M. Athawale, and S. Chakraborty, *Selection of materials using compromise ranking and outranking methods.* Materials & Design, 2009.

*fundamental of high entropy alloys with promising high–temperature properties.* Journal of Alloys and Compounds, 2018.

**15**(10): p. 4504-4519.

**30**(10): p. 4043-4053.

**760**: p. 15-30.

Business Media.

213-218.

[42] Chen, J., et al., *A review on* 

[43] Hummel, R.E., *Understanding materials science: history, properties, applications*. 2004: Springer Science &

[44] Cantor, B., et al., *Microstructural* 

[45] Ye, Y., et al., *High-entropy alloy: challenges and prospects.* Materials Today,

[47] Joseph, J., *Study of direct laser fabricated high entropy alloys*. 2016,

[48] Kuehl, R.O. and R. Kuehl, *Design of experiments: statistical principles of research design and analysis.* 2000.

[49] Wang, Y.P., B.S. Li, and H.Z. Fu, *Solid solution or intermetallics in a highentropy alloy.* Advanced engineering materials, 2009. **11**(8): p. 641-644.

[50] Gludovatz, B., et al., *A fractureresistant high-entropy alloy for cryogenic* 

*multicomponent alloys.* Materials Science and Engineering: A, 2004. **375**: p.

[46] Yeh, J.W., et al., *Nanostructured highentropy alloys with multiple principal elements: novel alloy design concepts and outcomes.* Advanced Engineering Materials, 2004. **6**(5): p. 299-303.

*development in equiatomic* 

2016. **19**(6): p. 349-362.

Deakin University.

**108**

[61] Xiao, K., et al., *A scanning AC calorimetry technique for the analysis of nano-scale quantities of materials.* Review of Scientific Instruments, 2012. **83**(11): p. 114901.

[62] Wang, Z., et al., *Effect of coherent L12 nanoprecipitates on the tensile behavior of a fcc-based high-entropy alloy.* Materials Science and Engineering: A, 2017. **696**: p. 503-510.

[63] Gwalani, B., et al., *Stability of ordered L12 and B2 precipitates in face centered cubic based high entropy alloys-Al0. 3CoFeCrNi and Al0. 3CuFeCrNi2.* Scripta Materialia, 2016. **123**: p. 130-134.

[64] Borkar, T., et al., *A combinatorial assessment of AlxCrCuFeNi2 (0< x< 1.5) complex concentrated alloys: Microstructure, microhardness, and magnetic properties.* Acta Materialia, 2016. **116**: p. 63-76.

[65] Li, Y., et al., *Microstructure and elevated-temperature mechanical properties of refractory AlMo0. 5NbTa0. 5TiZr High Entropy Alloy fabricated by powder metallurgy.* arXiv preprint arXiv:1801.00263, 2017.

[66] Kai, W., et al., *The oxidation behavior of a Ni2FeCoCrAl0. 5 highentropy superalloy in O2-containing environments.* Corrosion Science, 2019. **158**: p. 108093.

[67] Shafiee, A., et al., *Development and microstructural characterization of a new wrought high entropy superalloy.* Metals and Materials International, 2019: p. 1-12.

[68] Saito, T., et al., *Effect of Heat Treatments on the Microstructural Evolution of a Single Crystal High-Entropy Superalloy.* Metals, 2020. **10**(12): p. 1600.

[69] Zhang, L., et al., *Microstructure and mechanical properties of precipitationhardened cast high-entropy superalloys.* Materials Science and Technology, 2020: p. 1-7.

[70] Belan, J., *GCP and TCP phases presented in nickel-base superalloys.* Materials Today: Proceedings, 2016. **3**(4): p. 936-941.

[71] Yeh, A.-c. and T.-K. Tsao, *Highentropy superalloy*. 2019, Google Patents.

[72] Zhang, J., et al., *The effect of lattice misfit on the dislocation motion in superalloys during high-temperature low-stress creep.* Acta Materialia, 2005. **53**(17): p. 4623-4633.

[73] Mughrabi, H., *The importance of sign and magnitude of γ/γ' lattice misfit in superalloys—with special reference to the new γ'-hardened cobalt-base superalloys.* Acta materialia, 2014. **81**: p. 21-29.

[74] Chawla, V., et al., *Corrosion Behavior of Nanostructured TiAlN and AlCrN Hard Coatings on Superfer 800H Superalloy in Simulated Marine Environment.* Journal of Minerals and Materials Characterization and Engineering, 2009. **8**(9): p. 693-700.

[75] Sidhu, T., et al., *Oxidation and hot corrosion resistance of HVOF WC-NiCrFeSiB coating on Ni-and Fe-based superalloys at 800 C.* Journal of Thermal Spray Technology, 2007. **16**(5- 6): p. 844-849.

[76] Moody, N., et al., *Temperature effects on hydrogen-induced crack growth susceptibility of iron-based superalloys.* Engineering Fracture Mechanics, 2001. **68**(6): p. 731-750.

[77] Chung, D.-W., et al., *Effects of Cr on the properties of multicomponent cobalt-based superalloys with ultra high γ'volume fraction.* Journal of Alloys and Compounds, 2020: p. 154790.

[78] Makineni, S., B. Nithin, and K. Chattopadhyay, *Synthesis of a new tungsten-free γ–γ' cobalt-based superalloy by tuning alloying additions.* Acta Materialia, 2015. **85**: p. 85-94.

[79] Peters, M., et al., *Titanium alloys for aerospace applications.* Advanced engineering materials, 2003. **5**(6): p. 419-427.

[80] Clemens, H. and W. Smarsly. *Lightweight intermetallic titanium aluminides– status of research and development*. in *Advanced materials research*. 2011. Trans Tech Publ.

[81] Wee, S., et al., *Review on Mechanical Thermal Properties of Superalloys and Thermal Barrier Coating Used in Gas Turbines.* Applied Sciences, 2020. **10**(16): p. 5476.

[82] Goward, G.W., *Current research on the surface protection of superalloys for gas turbine engines.* JOM, 1970. **22**(10): p. 31-39.

[83] Couturier, R. and C. Escaravage, *High temperature alloys for the HTGR Gas Turbine: Required properties and development needs*. 2001.

[84] Sims, C.T., N.S. Stoloff, and W.C. Hagel, *superalloys II*. 1987: Wiley New York.

[85] Zhu, S. and F. Wang, *Nanocrystalline, Enamel and Composite Coatings for Superalloys*, in *Production, Properties, and Applications of High Temperature Coatings*. 2018, IGI Global. p. 160-186.

[86] Levi, C.G., *Emerging materials and processes for thermal barrier systems.* Current Opinion in Solid State and Materials Science, 2004. **8**(1): p. 77-91.

[87] Goward, G., *Progress in coatings for gas turbine airfoils.* Surface and coatings technology, 1998. **108**: p. 73-79.

[88] Galetz, M.C., *Coatings for superalloys*, in *Superalloys*. 2015, InTech. p. 277-298.

[89] Chatterji, D., R. DeVries, and G. Romeo, *Protection of superalloys for turbine application*, in *Advances in corrosion science and technology*. 1976, Springer. p. 1-87.

[90] Wu, C., et al., *Phase evolution and cavitation erosion-corrosion behavior of FeCoCrAlNiTix high entropy alloy coatings on 304 stainless steel by laser surface alloying.* Journal of Alloys and Compounds, 2017. **698**: p. 761-770.

[91] Zhang, S., et al., *Synthesis and characterization of FeCoCrAlCu high-entropy alloy coating by laser surface alloying.* Surface and Coatings Technology, 2015. **262**: p. 64-69.

[92] Huang, C., et al., *Dry sliding wear behavior of laser clad TiVCrAlSi high entropy alloy coatings on Ti–6Al–4V substrate.* Materials & Design, 2012. **41**: p. 338-343.

[93] Nahmany, M., et al., *Al x CrFeCoNi High-Entropy Alloys: Surface Modification by Electron Beam Bead-on-Plate Melting.* Metallography, Microstructure, and Analysis, 2016. **5**(3): p. 229-240.

[94] Tian, Y., et al., *Research progress on laser surface modification of titanium alloys.* Applied Surface Science, 2005. **242**(1-2): p. 177-184.

[95] Herd, R.M., J.S. Dover, and K.A. Arndt, *Basic laser principles.* Dermatologic clinics, 1997. **15**(3): p. 355-372.

[96] Natto, Z.S., et al., *Comparison of the efficacy of different types of lasers for the treatment of peri-implantitis: a systematic review.* International Journal of Oral & Maxillofacial Implants, 2015. **30**(2).

[97] Zhang, H., et al. *Synthesis and characterization of NiCoFeCrAl3 high* 

**111**

*Recent Advances of High Entropy Alloys: High Entropy Superalloys*

*alloy.* Journal of Laser Applications,

[107] Rekha, M., N. Mallik, and C. Srivastava, *First report on high entropy alloy nanoparticle decorated graphene.* Scientific reports, 2018. **8**(1): p. 1-10.

[108] Xu, Y., et al., *Microstructure Evolution and Properties of Laser Cladding CoCrFeNiTiAlx High-Entropy* 

*Alloy Coatings.* Coatings, 2020.

[109] Chen, C., et al., *Influences of laser surface melting on microstructure, mechanical properties and corrosion resistance of dual-phase Cr–Fe–Co–Ni–Al high entropy alloys.* Journal of Alloys and Compounds, 2020. **826**: p. 154100.

[110] Cai, Z., et al., *Microstructure and wear resistance of laser cladded Ni-Cr-Co-Ti-V high-entropy alloy coating after laser remelting processing.* Optics & Laser

Technology, 2018. **99**: p. 276-281.

[111] Ion, J., *Laser transformation hardening.* Surface engineering, 2002.

Technology, 2017. **4**(03): p. 4.

[113] Zhang, S., et al., *Laser surface alloying of FeCoCrAlNi high-entropy alloy on 304 stainless steel to enhance corrosion and cavitation erosion resistance.* Optics & Laser Technology, 2016. **84**:

[114] Jiang, P., et al., *Microstructure and Properties of CeO 2-Modified FeCoCrAlNiTi High-Entropy Alloy Coatings by Laser Surface Alloying.* Journal of Materials Engineering and

[115] Pawlowski, L., *Thick laser coatings: A review.* Journal of thermal spray technology, 1999. **8**(2): p. 279-295.

Performance, 2020: p. 1-10.

[112] Manilal, K.M., et al., *A Review on Laser Surface Alloying.* International Research Journal of Engineering and

**18**(1): p. 14-31.

p. 23-31.

**10**(4): p. 373.

2020. **32**(2): p. 022005.

*DOI: http://dx.doi.org/10.5772/intechopen.96661*

*entropy alloy coating by laser cladding*. in *Advanced Materials Research*. 2010.

[98] Zheng, B., Q.B. Liu, and L.Y. Zhang. *Microstructure and properties of MoFeCrTiW high-entropy alloy coating prepared by laser cladding*. in *Advanced Materials Research*. 2013. Trans

[99] Ye, X., et al., *The property research on high-entropy alloy AlxFeCoNiCuCr coating by laser cladding.* Physics Procedia, 2011. **12**: p. 303-312.

[100] Guo, X., et al., *Corrosion behavior of aluminum in fluoride-containing discharge condition for excimer laser structure application.* Materials Research

Express, 2019. **6**(10): p. 106519.

[101] Sharma, Y., et al., *Magnetic anisotropy in single-crystal high-entropy perovskite oxide La (C r 0.2 M n 0.2 F e 0.2 C o 0.2 N i 0.2) O 3 films.* Physical Review Materials, 2020. **4**(1):

[102] Dobbelstein, H., et al., *Direct metal deposition of refractory high entropy alloy MoNbTaW.* Physics Procedia, 2016. **83**:

[103] Nam, H., et al., *Effect of post weld heat treatment on weldability of high entropy alloy welds.* Science and Technology of Welding and Joining,

[104] Rafique, M.M.A., *Additive Manufacturing of Bulk Metallic Glasses and their composites–Recent trends and* 

[105] Ocelík, V., et al., *Additive manufacturing of high-entropy alloys by laser processing.* Jom, 2016. **68**(7): p.

[106] Fan, Y., et al., *Effect of fiber laser welding on solute segregation and proprieties of CoCrCuFeNi high entropy* 

2018. **23**(5): p. 420-427.

Trans Tech Publ.

Tech Publ.

p. 014404.

p. 624-633.

*approaches.*

1810-1818.

*Recent Advances of High Entropy Alloys: High Entropy Superalloys DOI: http://dx.doi.org/10.5772/intechopen.96661*

*entropy alloy coating by laser cladding*. in *Advanced Materials Research*. 2010. Trans Tech Publ.

*Advances in High-Entropy Alloys - Materials Research, Exotic Properties and Applications*

[88] Galetz, M.C., *Coatings for* 

p. 277-298.

Springer. p. 1-87.

p. 338-343.

*superalloys*, in *Superalloys*. 2015, InTech.

[89] Chatterji, D., R. DeVries, and G. Romeo, *Protection of superalloys for turbine application*, in *Advances in corrosion science and technology*. 1976,

[90] Wu, C., et al., *Phase evolution and cavitation erosion-corrosion behavior of FeCoCrAlNiTix high entropy alloy coatings on 304 stainless steel by laser surface alloying.* Journal of Alloys and Compounds, 2017. **698**: p. 761-770.

[91] Zhang, S., et al., *Synthesis and characterization of FeCoCrAlCu high-entropy alloy coating by laser surface alloying.* Surface and Coatings Technology, 2015. **262**: p. 64-69.

[92] Huang, C., et al., *Dry sliding wear behavior of laser clad TiVCrAlSi high entropy alloy coatings on Ti–6Al–4V substrate.* Materials & Design, 2012. **41**:

[93] Nahmany, M., et al., *Al x CrFeCoNi High-Entropy Alloys: Surface Modification by Electron Beam Bead-on-Plate Melting.* Metallography, Microstructure, and Analysis, 2016. **5**(3): p. 229-240.

[94] Tian, Y., et al., *Research progress on laser surface modification of titanium alloys.* Applied Surface Science, 2005.

[96] Natto, Z.S., et al., *Comparison of the efficacy of different types of lasers for the treatment of peri-implantitis: a systematic review.* International Journal of Oral & Maxillofacial Implants, 2015. **30**(2).

[97] Zhang, H., et al. *Synthesis and characterization of NiCoFeCrAl3 high* 

[95] Herd, R.M., J.S. Dover, and K.A. Arndt, *Basic laser principles.* Dermatologic clinics, 1997. **15**(3): p.

**242**(1-2): p. 177-184.

355-372.

[78] Makineni, S., B. Nithin, and K. Chattopadhyay, *Synthesis of a new tungsten-free γ–γ' cobalt-based superalloy by tuning alloying additions.* Acta Materialia, 2015. **85**: p. 85-94.

[79] Peters, M., et al., *Titanium alloys for aerospace applications.* Advanced engineering materials, 2003. **5**(6): p.

[80] Clemens, H. and W. Smarsly. *Lightweight intermetallic titanium aluminides– status of research and development*. in *Advanced materials research*. 2011. Trans

[81] Wee, S., et al., *Review on Mechanical Thermal Properties of Superalloys and Thermal Barrier Coating Used in Gas Turbines.* Applied Sciences, 2020.

[82] Goward, G.W., *Current research on the surface protection of superalloys for gas turbine engines.* JOM, 1970. **22**(10):

[83] Couturier, R. and C. Escaravage, *High temperature alloys for the HTGR Gas Turbine: Required properties and* 

[84] Sims, C.T., N.S. Stoloff, and W.C. Hagel, *superalloys II*. 1987: Wiley

*Nanocrystalline, Enamel and Composite Coatings for Superalloys*, in *Production, Properties, and Applications of High Temperature Coatings*. 2018, IGI Global.

[86] Levi, C.G., *Emerging materials and processes for thermal barrier systems.* Current Opinion in Solid State and Materials Science, 2004. **8**(1):

[87] Goward, G., *Progress in coatings for gas turbine airfoils.* Surface and coatings

technology, 1998. **108**: p. 73-79.

*development needs*. 2001.

[85] Zhu, S. and F. Wang,

419-427.

Tech Publ.

**10**(16): p. 5476.

p. 31-39.

New York.

p. 160-186.

p. 77-91.

**110**

[98] Zheng, B., Q.B. Liu, and L.Y. Zhang. *Microstructure and properties of MoFeCrTiW high-entropy alloy coating prepared by laser cladding*. in *Advanced Materials Research*. 2013. Trans Tech Publ.

[99] Ye, X., et al., *The property research on high-entropy alloy AlxFeCoNiCuCr coating by laser cladding.* Physics Procedia, 2011. **12**: p. 303-312.

[100] Guo, X., et al., *Corrosion behavior of aluminum in fluoride-containing discharge condition for excimer laser structure application.* Materials Research Express, 2019. **6**(10): p. 106519.

[101] Sharma, Y., et al., *Magnetic anisotropy in single-crystal high-entropy perovskite oxide La (C r 0.2 M n 0.2 F e 0.2 C o 0.2 N i 0.2) O 3 films.* Physical Review Materials, 2020. **4**(1): p. 014404.

[102] Dobbelstein, H., et al., *Direct metal deposition of refractory high entropy alloy MoNbTaW.* Physics Procedia, 2016. **83**: p. 624-633.

[103] Nam, H., et al., *Effect of post weld heat treatment on weldability of high entropy alloy welds.* Science and Technology of Welding and Joining, 2018. **23**(5): p. 420-427.

[104] Rafique, M.M.A., *Additive Manufacturing of Bulk Metallic Glasses and their composites–Recent trends and approaches.*

[105] Ocelík, V., et al., *Additive manufacturing of high-entropy alloys by laser processing.* Jom, 2016. **68**(7): p. 1810-1818.

[106] Fan, Y., et al., *Effect of fiber laser welding on solute segregation and proprieties of CoCrCuFeNi high entropy*  *alloy.* Journal of Laser Applications, 2020. **32**(2): p. 022005.

[107] Rekha, M., N. Mallik, and C. Srivastava, *First report on high entropy alloy nanoparticle decorated graphene.* Scientific reports, 2018. **8**(1): p. 1-10.

[108] Xu, Y., et al., *Microstructure Evolution and Properties of Laser Cladding CoCrFeNiTiAlx High-Entropy Alloy Coatings.* Coatings, 2020. **10**(4): p. 373.

[109] Chen, C., et al., *Influences of laser surface melting on microstructure, mechanical properties and corrosion resistance of dual-phase Cr–Fe–Co–Ni–Al high entropy alloys.* Journal of Alloys and Compounds, 2020. **826**: p. 154100.

[110] Cai, Z., et al., *Microstructure and wear resistance of laser cladded Ni-Cr-Co-Ti-V high-entropy alloy coating after laser remelting processing.* Optics & Laser Technology, 2018. **99**: p. 276-281.

[111] Ion, J., *Laser transformation hardening.* Surface engineering, 2002. **18**(1): p. 14-31.

[112] Manilal, K.M., et al., *A Review on Laser Surface Alloying.* International Research Journal of Engineering and Technology, 2017. **4**(03): p. 4.

[113] Zhang, S., et al., *Laser surface alloying of FeCoCrAlNi high-entropy alloy on 304 stainless steel to enhance corrosion and cavitation erosion resistance.* Optics & Laser Technology, 2016. **84**: p. 23-31.

[114] Jiang, P., et al., *Microstructure and Properties of CeO 2-Modified FeCoCrAlNiTi High-Entropy Alloy Coatings by Laser Surface Alloying.* Journal of Materials Engineering and Performance, 2020: p. 1-10.

[115] Pawlowski, L., *Thick laser coatings: A review.* Journal of thermal spray technology, 1999. **8**(2): p. 279-295.

Section 3

Design of High-Entropy

Alloys

**113**

## Section 3
