Metallic Amorphous Alloy Reinforcements in Light Metal Matrices

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S P R I N G E R B R I E F S I N M AT E R I A L S

S. JayalakshmiM. Gupta

Metallic Amorphous Alloy Reinforcements in Light Metal Matrices

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S. Jayalakshmi • M. Gupta

Metallic Amorphous Alloy Reinforcements in Light Metal Matrices

ISSN 2192-1091 ISSN 2192-1105 (electronic) SpringerBriefs in Materials ISBN 978-3-319-15015-4 ISBN 978-3-319-15016-1 (eBook) DOI 10.1007/978-3-319-15016-1

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Springer Cham Heidelberg New York Dordrecht London © S. Jayalakshmi and M. Gupta 2015 This work is subject to copyright. All rights are reserved by the Publisher, whether the whole or part of the material is concerned, specifi cally the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfi lms or in any other physical way, and transmission or information storage and retrieval, electronic adaptation, computer software, or by similar or dissimilar methodology now known or hereafter developed. The use of general descriptive names, registered names, trademarks, service marks, etc. in this publication does not imply, even in the absence of a specifi c statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. The publisher, the authors and the editors are safe to assume that the advice and information in this book are believed to be true and accurate at the date of publication. Neither the publisher nor the authors or the editors give a warranty, express or implied, with respect to the material contained herein or for any errors or omissions that may have been made.

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S. Jayalakshmi Department of Mechanical EngineeringBannari Amman Institute of Technology Sathyamangalam, Tamil Nadu, India

M. Gupta Department of Mechanical EngineeringNational University of Singapore Singapore, Singapore

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Contents

1 Introduction ................................................................................................. 1 1.1 Background .......................................................................................... 1 1.2 Metal Matrix Composites ..................................................................... 2

1.2.1 Processing of MMCs/LMMCs ................................................. 2 1.2.2 Properties, Applications, and Limitations

of MMCs/LMMCs ................................................................... 3 1.3 Metallic Amorphous Alloys ................................................................. 4 References ..................................................................................................... 5

2 Light Metal Matrix Composites ................................................................ 7 2.1 Background .......................................................................................... 7 2.2 Characteristics of MMCs ..................................................................... 8

2.2.1 Importance of Interfacial Bonding ........................................... 11 2.2.2 Role of the Metallic Matrix ...................................................... 12

2.3 Processing of LMMCs ......................................................................... 13 2.3.1 Liquid-State Processes ............................................................. 13 2.3.2 Solid State Processes ................................................................ 21 2.3.3 Semisolid State Processes ........................................................ 25 2.3.4 Other Processes ........................................................................ 28

2.4 Strengthening Mechanisms in LMMCs ............................................... 31 2.4.1 Strength Prediction ................................................................... 32 2.4.2 Fracture Mechanisms ............................................................... 33

2.5 Microstructural and Mechanical Properties of LMMCs ...................... 34 2.5.1 Al-Composites ......................................................................... 34 2.5.2 Mg Composites ........................................................................ 40

2.6 Limitations ........................................................................................... 53 References ..................................................................................................... 53

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3 Amorphous Alloys/Bulk Metallic Glasses (BMG) .................................. 59 3.1 Background ........................................................................................ 59

3.1.1 Formation and Characteristics of Amorphous Alloys/ BMG ................................................. 60

3.2 Preparation Methods .......................................................................... 67 3.2.1 Rapid Solidifi cation Processes ............................................... 67 3.2.2 Methods for Producing BMG ................................................ 69 3.2.3 Mechanical Alloying .............................................................. 70

3.3 Structural, Thermal, and Mechanical Properties ................................ 72 3.4 Limitations ......................................................................................... 80 References ................................................................................................... 82

4 Light Metal Matrix Composites with Amorphous Alloys/Bulk Metallic Glass Reinforcements (BMG) .............................. 85 4.1 Introduction ........................................................................................ 85 4.2 Synthesis, Matrix Reinforcement Selection, and Properties .............. 86

4.2.1 Liquid State Processing: Infi ltration Method ......................... 86 4.2.2 Solid State Processing: Powder

Metallurgy-Based Methods .................................................... 89 4.2.3 Other Methods/Systems ......................................................... 104

4.3 Conclusions ........................................................................................ 105 References ................................................................................................... 105

5 Future Work .............................................................................................. 107 5.1 Future Research and Applicative Prospects ....................................... 107

Index ................................................................................................................. 109

Contents

1© S. Jayalakshmi and M. Gupta 2015 S. Jayalakshmi, M. Gupta, Metallic Amorphous Alloy Reinforcements in Light Metal Matrices, SpringerBriefs in Materials, DOI 10.1007/978-3-319-15016-1_1

Chapter 1 Introduction

Abstract Light metal matrix composites are briefl y introduced. The various conventional ceramic reinforcements (micro-/nano-size) that are used to make the composites and the MMC processing techniques are mentioned. The properties, applications, and limitations of metal matrix composites are summarized. In view of the limitations of ceramic reinforcements, metallic amorphous alloys are proposed as alternate reinforcement materials. Their inherent superior properties that can contribute to the enhanced properties of the composites are highlighted.

Keywords Light metals • Metal matrix composites • Ceramic micro-/nano-scale reinforcements • Processing and properties • Metallic amorphous alloys reinforcements

1.1 Background

Global concern over the energy crisis that is being faced worldwide has seriously pushed research to identify effi cient and robust solutions to meet the need. Rapid depletion of oil reserves, increasing demand for fuel effi ciency, and regulations on emission reduction has turned the attention towards lightweight materials. Research on these materials is largely focused to achieve multiple-performance reliability, along with easier material processing, machinability/formability, and high load- bearing capacity/structural strength. Energy effi ciency, recyclability, and sustai-nability are also in the focus. Given this context, R&D of Al and Mg is of great interest, especially for weight-critical applications such as in automotive, aviation, sports, electronics, and communication sectors (Rohatgi 1996 ; Miracle 2005 ; Kainer 2006 ; Sharon and Gupta 2011 ). In comparison with the density of steel (8.1 g/cc), Al and Mg have densities of 2.74 and 1.74 g/cc, respectively, and are therefore the lightest among the structural metals. To note, they offer high specifi c strength properties, provide energy/fuel effi ciency, and are recyclable. Al-alloys exhibit excellent ductility, superior resistance to corrosion, and have good thermal and electrical conductivities. Mg-alloys possess excellent castability, machinability, damping capacity, impact, and dent resistance. Both Al- and Mg-alloys can be

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precipitation hardened which gives them fl exibility in attaining properties suitable for a wide variety of commercial/industrial applications (Miller et al. 2000 ; Mordike and Ebert 2001 ).

Applications for high-strength light metal components made of Al and Mg are constantly growing. For instance, using these metals in automobile engine and chassis components reduces weight and lowers fuel consumption/emissions, thereby signifi cantly contributing towards fuel economy. As an example, replacing the existing cast iron engine block (~84.6 kg) in a V6 3.0 L six-cylinder car by Mg (~30 kg) would result in an overall weight reduction of ~65 %. The increasing use of Al- and Mg-based materials in the automotive industry is an excellent example of materials selection, wherein factors such as material availability, processability, cost, properties, environmental issues, recyclability, and fuel effi ciency are all taken into account, together.

1.2 Metal Matrix Composites

Despite the fact that Al- and Mg-alloys have several attractive properties, they do not completely satisfy the overall requirement in applications, where the components are required to withstand high mechanical/thermal stresses and also under tribological conditions. Hence, there arises a need to improve their properties so to realize their full potential in commercial applications. A reliable route to achieve this objective is to make composites from the light metals/alloys (Chawla and Chawla 2006 ).

Composites are combinations of dissimilar materials produced to achieve desired properties. Owing to the increasing demand for lightweight materials for structural applications in automobile and aerospace industries, metal matrix composites (MMCs) have become popular. MMCs are based on the principle of incorporating a high performance second phase (oxides, carbide ceramics such as SiC, Al 2 O 3 , B 4 C, and TiN of known volume fraction, V f ) in a conventional engineering material (here, light metal matrix such as Al, Ti, and Mg). The reinforcement phase in the MMC is the secondary phase. The reinforcements are classifi ed into two major categories—continuous reinforcements and discontinuous reinforcements. Due to its high stiff-ness and high strength, the reinforcement is the main load-bearing member in the composite. Reinforcements in MMCs are usually ceramics in the form of fi bers, whiskers, or particles. In the composite, the morphology and nature of the reinforce-ments are very important in controlling the fi nal properties, as their interaction with the matrix would alter the MMC’s microstructure, properties, and performance.

1.2.1 Processing of MMCs/LMMCs

Conventional composites are produced by liquid, solid, and semi-solid state processes. The choice of the processing route depends on several factors such as the reinforcement type, its distribution, matrix-particle bonding, control of matrix micro-structure, process simplicity, and cost-effectiveness (Chawla and Chawla 2006 ).

1 Introduction

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Liquid-state processes . Liquid-state processing routes are attractive as they are relatively simple, cost-effective, and are potentially scalable to industrial level. Liquid-state processes involve the incorporation of reinforcement into the molten metal followed by casting. Direct addition of reinforcements into the melt or incor-poration in the form of fi ber/particle preforms is usually employed. Some of the liquid-state processing routes include stir casting, ultrasonic-assisted casting, infi l-tration techniques, and disintegrated melt deposition method.

Solid-state processes . Solid-state processes are based on powder metallurgy (PM) methods, capable of producing near-net shape components. PM processes usually involve: blending of matrix alloy and reinforcing phase powders; compacting the blend, usually by cold pressing. The green compacts are then consolidated by different routes such as direct sintering, hot isostatic pressing (HIP), vacuum HIP, hot extrusion or cold sintering, microwave-assisted sintering, and spark plasma sintering.

Semi-solid state processes . In semi-solid state processes, incorporation of the rein-forcements in matrix metal is conducted when a matrix is in semi-solid state (slurry), as it facilitates: (1) uniform dispersion of reinforcements, (2) eliminates settling of denser reinforcement due to slurry-state, and (3) eliminates porosity (that usually occurs due to voids generated at the molten metal/particle interface).

1.2.2 Properties, Applications, and Limitations of MMCs/LMMCs

Conventional light metal matrix composites (LMMCs) are incorporated with various micron-sized stronger/stiffer nonmetallic/ceramic reinforcements (e.g., Al 2 O 3 , SiC, C) (Polmear 1995 ; Brook 1998 ; Avedesian and Baker 1999 ) which provide:

• Increased mechanical strength • Higher wear resistance • Improved thermal expansion • Enhanced thermal stability

Due in part to these reasons, light metal-based MMCs have been used in some products related to automotive, electronic packaging, industrial products, and recre-ational goods, such as pickup truck drive shafts, brake rotors/pads and drums, diesel engine pistons, aeronautic engine fan, exit guide vanes, aircraft ventral fi ns, fuel access covers, bicycle components, golf clubs, and electronic packaging applications. Examples of tribological applications include cylinder sleeves in engines, piston-recess walls, and brake discs/pads (Suresh et al. 1993 ; Clyne and Withers 1995 ; Deuis et al. 1997 ; Surappa 2003 ).

Although such MMCs have been proposed/used for several applications, it should be noted that their low ductility along with their poor machinability and weldability has restricted its complete use. The low ductility is caused by the poor interfacial characteristics between the reinforcement and matrix. Due to this drawback, they exhibit low fracture toughness. Further, under tribological conditions, the brittleness and hardness of these materials can cause damage to the counterfaces.

1.2 Metal Matrix Composites

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One practical solution to overcome the drawback of low ductility of light metal- based MMCs is to incorporate nano-sized reinforcements, i.e., to make nanocomposites (LMMNCs) (Choi and Awaji 2005 ; Dieringa 2010 ). Research investigations on both Al- and Mg-nano-composites have shown that the nanocomposites show: (1) improvement in yield strength and (2) enhancement/retention in ductility, thereby improving the fracture toughness. The superior properties of LMMNCs when compared to MMCs are mainly because of the “nanometer size” of the reinforcement which gives rise to dispersion strengthening-like effect (strength-ening arising due to uniform distribution of fi ne/hard particles). The production routes of conventional LMMCs are also suitable for nano-composites production.

However, not all processing methods can be directly employed to produce nanocomposites, as several critical factors such as particle agglomeration and dis-tribution govern the fi nal properties. Hence, other alternative processes to incorpo-rate nano reinforcement particles into metal matrices are being developed. Friction stir process (FSP) is one such method, which is based on friction stir welding. Other processes are accumulative roll bonding (ARB) and equal channel angular pressing methods. These are severe plastic deformation process (SPD) used to produce fi ne- grained materials. Research on full utilization of these methods to produce nanocomposites is still in the developmental stage.

1.3 Metallic Amorphous Alloys

Considering the limitations of ceramic reinforcements, viz., extreme brittleness and poor interfacial characteristics in the composite (formation of undesired intermetal-lic phases, structural, mechanical, thermal, and chemical incompatibility between matrix-reinforcement-interfacial products), new materials that can retain the posi-tive characteristics of ceramic reinforcements and at the same time overcome their limitations are sought as effective reinforcements.

In this view point, metallic amorphous alloys/bulk metallic glasses (BMG) are new class of metallic materials that are distinctly different from conventional met-als/alloys in terms of their structure and thermal behavior, and exhibit extremely high strength (~1–2 GPa) and large elastic strain limit (~1–2 %) (Miller and Liaw 2008 ). Given these unique properties, upon their incorporation into light metal matrices (Al/Mg-matrices), they will provide superior interfacial properties, i.e., high degree of compatibility with the matrix due to their metallic nature, when com-pared to conventional ceramic reinforcements. Further, the closer thermal coeffi -cient of expansion values between the metal matrix and the amorphous reinforcement will result in better interfacial stability, unlike that observed in conventional ceramic reinforcements. These together can signifi cantly enhance the mechanical perfor-mance of LMMCs.

Amorphous/BMG-reinforced composite is an emerging research fi eld and the existing literature is meager. The various processing methods that would be suit-able for producing amorphous alloy-reinforced MMCs should be identifi ed. Upon its successful synthesis, a comparison of mechanical properties and strengthening

1 Introduction

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mechanisms of amorphous/BMG-reinforced composites with those of the conventional ceramic composites is essential so as to elucidate the advantages of the amorphous alloy/BMG reinforcement. A discussion on future research direc-tions and wider research potential of the novel materials are required to identify the prospective applications of these materials in automobiles, aerospace, sports, and consumer electronics industries.

References

Avedesian MM, Baker H (1999) Magnesium and magnesium alloys. ASM International, Materials Park

Brook G (1998) Smithells light metals handbook. Butterworth-Heinemann, Oxford Chawla N, Chawla K (2006) Metal matrix composites. Springer, New York Choi S-M, Awaji H (2005) Nanocomposites—a new material design concept. Sci Technol Adv

Mater 6:2–10 Clyne T, Withers P (1995) An introduction to metal matrix composites. Cambridge University

Press, Cambridge Deuis R, Subramanian C, Yellup J (1997) Dry sliding wear of aluminium composites—a review.

Compos Sci Technol 57:415–435 Dieringa H (2010) Properties of magnesium alloys reinforced with nanoparticles and carbon nano-

tubes: a review. J Mater Sci 46:289–306. doi: 10.1007/s10853-010-5010-6 Kainer K (2006) Metal matrix composites: custom-made materials for automotive and aerospace

engineering. Wiley, Weinheim Miller M, Liaw P (2008) Bulk metallic glasses. Springer, New York Miller W, Zhuang L, Bottema J et al (2000) Recent development in aluminium alloys for the auto-

motive industry. Mater Sci Eng A 280:37–49 Miracle D (2005) Metal matrix composites—from science to technological signifi cance. Compos

Sci Technol 65:2526–2540. doi: 10.1016/j.compscitech.2005.05.027 Mordike B, Ebert T (2001) Magnesium properties—applications—potential. Mater Sci Eng A

302:37–45 Polmear I (1995) Light alloys: metallurgy of the light metals. Wiley, New York Rohatgi P (1996) Processing, properties, and applications of cast metal matrix composites. The

Minerals, Metals & Materials Society, Warrendale Sharon NM, Gupta M (2011) Magnesium, magnesium alloys, and magnesium composites. Wiley,

Hoboken Surappa M (2003) Aluminium matrix composites: challenges and opportunities. Sadhana 28:319 Suresh S, Mortensen A, Needleman A (1993) Fundamentals of metal-matrix composites.

Butterworth-Heinemann, Boston

References

http://dx.doi.org/10.1007/s10853-010-5010-6

http://dx.doi.org/10.1016/j.compscitech.2005.05.027

7© S. Jayalakshmi and M. Gupta 2015S. Jayalakshmi, M. Gupta, Metallic Amorphous Alloy Reinforcements in Light Metal Matrices, SpringerBriefs in Materials, DOI 10.1007/978-3-319-15016-1_2

Chapter 2Light Metal Matrix Composites

Abstract Fundamentals of metal matrix composites are overviewed. The various light metal matrix systems (particularly Al and Mg) and the different types of rein-forcements used are mentioned. The various composite production methods are described. Strengthening mechanisms that define the enhancement in properties of composites are discussed. The microstructural and mechanical properties of the composites are summarized. In addition, the several disadvantages encountered in MMCs due to ceramic reinforcement addition are understood from interfacial characteristic/properties.

Keywords Light metal matrix composites • Aluminum and magnesium matrices •Ceramic reinforcements • Liquid-state processing methods • Solid state and semi-solid state processes • Strengthening mechanisms • Microstructural and mechanicalproperties • Limitations

2.1 Background

For the past few decades, there has been a sharp demand for light weight structural materials, especially in the rapidly growing automotive and aerospace industries. In recent times, alloys of Al, Ti, and Mg have gained importance in such applica-tions because of their unique properties and comparative advantages such as light-weight, good machinability, dimensional stability, and energy efficiency/low power consumption. Despite the fact that Al- and Mg-alloys have several attractive proper-ties, they do not completely satisfy the overall requirements in applications wherethe components are required to withstand high mechanical/thermal stresses and alsounder tribological conditions. Their alloys present a noticeable decrease of mechan-ical properties at relatively low temperatures, less than about 200 °C (Surappa 2003; Friedrich and Mordike 2006) which strongly limits their application for critical components in the automotive and aerospace sectors. Hence, there arises a need to improve their properties so to realize their full potential in commercial applications. A reliable route to realize this objective is to make composites from the light metals/alloys. These are called light metal matrix composites (LMMCs) wherein the light

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metal matrix is incorporated with a hard reinforcing phase (typically ceramic or carbon based). In LMMCs, the synergistic effect of ductility and toughness of the metal matrix and the high strength and stiffness of the reinforcement are utilized. The introduction of such reinforcements into the light metal matrices significantly improve the mechanical properties such as hardness, tensile strength, elastic modu-lus, and yield strength, and also give rise to excellent wear resistance, Further improvement in properties can be achieved by proper heat/thermal treatments. Given these advantages, LMMCs are seen to have enormous potential in diverse industrial/commercial sectors such as automotive, aviation, biomedical, sporting equipments, consumer electronics, etc.

2.2 Characteristics of MMCs

In a metal matrix composite, three important features determine its characteristics: viz., the matrix, the reinforcement, and the matrix/reinforcement interface.

Matrix. Matrix is the continuous phase and its properties are improvised by convert-ing it into a composite with the introduction of an appropriate reinforcement. Selection of the matrix and the reinforcement largely depends upon the end use and the amenability for production. For a long time it was assumed that the only func-tion of the matrix was to hold the reinforcement in position. Over the years, the importance of the function of matrix and its influence on the properties has been well established. The metallic matrix being highly structure-sensitive, as any change in its microstructure (by the incorporation of reinforcement) would alter the overall performance of the composite.

Reinforcement. The reinforcement phase in the metal matrix composite is the sec-ondary phase. Due to its high stiffness and high strength, it is the main load bearing member in the composite. Reinforcements are usually ceramics in the form of fibers/whiskers/particles. The morphology and nature of the reinforcements are very important in controlling the final properties, as their interaction with the matrix would alter the composite microstructure, properties, and performance. Also, given the heat-treatable nature of both Al and Mg matrices, it is quite important to alsoconsider the effect of reinforcement on the precipitation mechanisms. Reinforcements are classified into two major categories—continuous reinforcements and discon-tinuous reinforcements.

1. Continuous reinforcements: Continuous fibers are those filamentary materials whose lengths are greater than 100 μm. They can be either amorphous, single crystalline, or polycrystalline. The properties of various continuous fibers are given in Table 2.1 (Harris 1988). Due to their unidirectional nature, the proper-ties of continuous fiber reinforced composites are anisotropic.

2. Discontinuous reinforcements: Short fibers, particles, and whiskers are classi-fied as discontinuous reinforcements. Tables 2.2, 2.3, and 2.4 list the properties

2 Light Metal Matrix Composites

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Table 2.1 Properties of various continuous fiber reinforcements (Harris 1988)

Fiber typeDiameter (μm)

Density (g/cm3)

Modulus (GPa)

UTS (MPa)

Breaking strain (%)

Boron, single fiber, SiC coated

100–150 2.70 400 3,100 0.77

α—Alumina, tow 20 3.95 379 1,380 0.36γ—Alumina, tow 17 3.25 210 1,800 0.85Carbon, high modulus, tow

8 1.85 400 2,300 0.58

Carbon, high tension, tow

8 1.75 23 2,800 1.10

SiC, tow 13 2.55 196 2,550 1.00SiC, single fiber 100–140 3.5 400 2,700 0.68

Table 2.2 Properties of particulate reinforcing materials used in MMCs (Girot et al. 1987)

Material Normal size used (μm) Density (g/cm3)

SiC 15–340 3.2SiO2 40–60 2.3MgO 40–60 2.7–3.6Si3N4 40–60 3.2TiC 40–50 2.25Al2O3 40–340 4.0B4C 40–300 2.5ZrO2 75–180 5.65–6.15BN2 40–50 2.25Graphite 40–250 1.6–2.2

Table 2.3 Properties of whiskers used in MMCs (Stacey 1988)

WhiskersLength, l (μm)

Diameter, d (μm)

Density (g/cm3)

Ultimate tensile strength, UTS (GPa)

Modulus of elasticity (GPa)

SiC (Tokai) 50–200 0.1–1 3.2 3–14 400–700SiC (Arco) 50 0.2–1 3.2 13 700Al2O3 100 2 3.97 14 2,275Si3N4 5–200 0.1–1.6 3.18 13.8 379Al2O3·B2O4 10–30 0.5–1 2.93 8 400

of the commonly used discontinuous reinforcements (Girot et al. 1987; Stacey 1988). The reinforcement of particles have gained major interest as they enable obtaining a strong enhancement of mechanical properties while maintaining an isotropic behavior, with relatively simple production routes (Maruyama 1998; Miracle 2005) and a possibility to use secondary processes (Ellis 1996; Manna and Bhattacharayya 2003). Whiskers are elongated single crystals that have a high degree of structural and chemical perfection and provide high strength and


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modulus of elasticity. Short/discontinuous fibers have much lower l/d ratio than continuous fibers. Both discontinuous fibers/particles have the following uniquefeatures when used as reinforcements.

• Isotropic properties• Easy fabrication by conventional methods• Sustainability of high operating temperatures• Increased modulus of elasticity and strength• Increased thermal stability• Increased wear resistance• Less expensive (cost-effective)

Recently nano-sized reinforcements (SiC, Al2O3, B4C, etc., of size usually <50 nm) have been used. Carbon nanotube (CNT) is also a promising reinforce-ment. The major advantage of using nano-reinforcements is that the enhancement in properties (including ductility) can be attained at lower volume fractions (<2 %), whereas for micron-scale particle reinforced MMCs higher volume fractions (≫10 %) are required. However, in order to produce sound nanocomposites withenhanced mechanical properties, good dispersion of the nano-reinforcement phase within the matrix is necessary, which is in turn strongly governed by the selection of a suitable production process.

Interface. Metcalfe (1974) defined the interface in a composite as “a surface formed by the common boundary of reinforcement and matrix in contact, which constitutes the bond in between for the transfer of loads.” The interface has physical and mechanical properties which are unique and not that of either the fiber/particle orthe matrix. The interface occupies a large area in the composite. Being in between the matrix and the reinforcement, the interface is thermodynamically unstable, and exerts an influence on the overall performance of the composite. It is to be further noted that the interface causes a large discontinuity in material parameters such as the modulus of elasticity (E), coefficient of thermal expansion (CTE-α), and the chemical potential of a composite. Such a difference, e.g., in the CTE between the reinforcement and the matrix, would cause a residual stress field and would affect

Table 2.4 Properties of short discontinuous fibers used in MMCs (Stacey 1988)

FibersLength, l (μm)

Diameter, d (μm)

Density (g/cm3) UTS (GPa)

Modulus of elasticity (GPa)

Carbon T300 2.5 7.8 1.75 3.45 230SiC (Nicalon) 1–6 10–16 2.55 3 195Al2O3 FP 3–6 15–25 3.96 1.7 380Al2O3 Saffil −3 1–5 3.30 2 300Saffil HA 0.1–3 1–5 3.40 1.5 300Fiberfrax 1–3 1–7 2.73 1.5 105Alumino silicate 2–5 1–7 3.00 0.8 150


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the properties in many complex ways. Similarly, a chemical potential discontinuity would lead to interdiffusion zone. This would provide preferential sites for any segregation and precipitation.

2.2.1 Importance of Interfacial Bonding

The basic requirement of a composite is the intimate contact and efficient bondingbetween the matrix phase and the reinforcement phase. Such a bond would pro-vide required efficiency in load transfer, usually from the matrix to the reinforce-ment. The factors influencing the qualities of an interfacial bond are brieflydiscussed below.

Wettability. Wettability is the ability of a liquid to spread on a solid surface. It isdependent on the contact angle attained, which describes the extent of intimate con-tact between a liquid and solid. In composites, the wettability is measured by theamount of work required by the molten metal to engulf the solid reinforcement.A contact angle between the liquid (molten metal) and the solid surface (reinforce-ment) close to 0° indicates perfect wetting, while the angle 180° indicates nonwetting (Fig. 2.1). The wettability, therefore, indicates the possibility of realizing an inti-mate contact between the liquid and the solid constituents in the composite(Mortenson et al. 1988). As most of the metal/reinforcement systems in MMCs show poor wettability, various measures have been attempted to improve wettabil-ity. Some such measures are pretreatment of reinforcements, alloying modifica-tions, and coating of reinforcements (Krishnan et al. 1981; Kimura et al. 1984; Rohatgi et al. 1986; Delanney et al. 1987).

Mechanical bonding. Mechanical bonding occurs due to the mechanical-keying effect between the surfaces. Increasing the surface roughness of the reinforcement usually produces such mechanical-keying effect. Rough surfaces could be achieved by the process of etching. One such example was observed in W/Al composite (Metcalfe 1974) where etching of tungsten wires was carried out to produce a rough interface in the composite. This caused the desired mechanical bonding and increased the stress levels. Artificially produced interface roughness by etch pitting or by microhardness indentations are some of the other methods employed to improve mechanical bonding (Chawla and Metzger 1978).

Fig. 2.1 Schematic diagram showing the contact angle between a liquid and solidsurface. When (a) θ = 0°—perfect wettability, (b) θ = 180°—no wetting, and (c) 0° < θ < 180°—partial wetting


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Chemical reactions. The chemical gradient across the reinforcement-matrix inter-face (produced under favorable conditions) would cause diffusion to occur or may cause chemical reactions between the components. The interface layer thus formed would then have characteristics that are much different from the individual constitu-tive components (Metcalfe 1974). Occasionally, these would prove beneficial and facilitate an improvement in the nature of the bond between the reinforcement and the matrix. This however depends on the critical thickness of the interface layer above which the effect is adverse. In most cases, the interfacial chemical reactions result in brittle oxides/carbides/intermetallics which act as local stress concentrators resulting in the brittle behavior of the composite.

Thermal stresses. Apart from wettability, mechanical bonding, and chemical effects, thermal stress is yet another important factor that affects the interfacial bonding. The linear coefficient of thermal expansion (CTE) of the metal matrix is usually higher than that of the ceramic reinforcement. Such thermal incompatibility (mis-match) between the matrix and the reinforcement leads to the formation of a large thermal residual stress field along the interface, which affects the characteristics of the composite in a complex manner. If the stresses so produced were of the order of the yield stress of the matrix metal, then it would cause local plastic deformation. This deformation would then introduce dislocations and vacancies at the interface; subsequently responsible for activating chemical reactions at those sites. The ther-mal mismatch is also believed to accelerate the precipitation kinetics of the matrix. Evidences of the existence of a hard zone of high dislocation density around the reinforcement and a soft zone of low dislocation density away from the reinforce-ment have been observed due to the thermal mismatch across the interface. This influences the property of the matrix, leading to a strengthening effect in nonprec-ipitation hardenable alloys and accelerating the aging kinetics in precipitation hard-enable alloy systems (Metcalfe 1974). Often such thermal mismatch would result in localized plastic deformation (inhomogeneous deformation in the composite).

2.2.2 Role of the Metallic Matrix

As mentioned earlier, the matrix plays an important role in the behavior of the com-posite. The matrix not only holds the reinforcements (fibers/particles) and aids in the transfer of load to the reinforcements, but also affects the overall performance of a composite. The chemical and thermal reactions of the matrix with the reinforcement result in a change in the matrix microstructure that would then alter the properties of the composite. The two factors that influence the matrix behavior are the following:

1. The difference in the CTE at the reinforcement/matrix interface (due to the ther-mal mismatch between them) leads to the formation of residual stress zones.

2. The chemical incompatibility at the reinforcement/matrix interface forms a new chemical zone. This new zone will have a CTE which is different from that of the matrix and the reinforcement making the overall behavior and response of the


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system more complicated. Possible events that are predictable from such reactions that are induced into the matrix are (Metcalfe 1974):

(a) Plastic deformation mechanisms such as slip, twinning, grain boundary slid-ing, and grain boundary migration, in the case of ductile matrix.

(b) Cracking and failure in the case of brittle matrix. (c) Failure at the reinforcement/matrix interface in both kinds of matrices.

2.3 Processing of LMMCs

LMMCs manufacturing processes are classified into ex situ routes, when the rein-forcing phase is produced at an earlier stage and then added to the matrix, or in situ routes, when the reinforcement is generated during the composite production, typi-cally through controlled reactions (Fridlyander 1994; Ye and Liu 2004; Tjong 2007). Ex situ processing techniques are commonly employed to make LMMCs andcan be further classified into liquid, solid, and semisolid state processes. Some ofthese production routes are also suitable for nanocomposites production. The choice of the processing route depends on several factors such as the reinforcement type, its distribution, matrix-particle bonding, control of matrix microstructure, process simplicity, and cost-effectiveness. The ex situ processing routes are described below.

2.3.1 Liquid-State Processes

Liquid-state processing routes are attractive as they are relatively simple, cost-effective, and are potentially scalable to industrial level. These routes include stir casting, ultrasonic-assisted casting, centrifugal casting, infiltration techniques, anddisintegrated melt deposition (DMD) method.

2.3.1.1 Stir Casting

Stir casting is also known as the vortex method and is widely used to produce LMMCs. It involves incorporation of reinforcements (mainly in the form of parti-cles) into the molten metal, followed by casting. Homogenous distribution of rein-forcement is achieved by:

1. A rotor rotating in the liquid metal that creates a vortex or2. By injection of a gas carrying the reinforcement into liquid metal (Fig. 2.2)

(Ezatpour et al. 2014)

The finely distributed slurry so produced is shaped by conventional casting tech-niques, viz. sand casting, permanent mould casting, pressure die casting or squeezecasting (Evans et al. 2003; Noguchi et al. 2008).


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Issues associated with stir casting process are: (1) gas entrapment and (2) slag in the melt. These lead to high porosity and microdefects. (3) Undesired chemical reactions at the matrix/reinforcement interface and (4) low wettability of reinforce-ments with the molten matrix that increases the tendency of the particles to agglom-erate (especially in nano-reinforcements that could result in the formation of nanoparticle clusters and nonuniform distribution of reinforcement). These issues if unchecked would cause severe deterioration of material properties. To successfully implement this process at both laboratory and industrial scales, careful standardiza-tion of process parameters such as temperature of molten metal, melt stirring time, stirring velocity, melt holding temperature, choice of matrix-reinforcements, etc., should be rightly selected (Koli et al. 2013).

2.3.1.2 Centrifugal Casting

Centrifugal casting is a relatively inexpensive process in which optimal reinforce-ment placements are achieved by inducing a centrifugal force immediately during casting (Fig. 2.3, http://www.adityainc.com/casting/centrifugal-casting.html 2014).

Fig. 2.2 Stir casting (vortex) method (adapted from Ezatpour et al. 2014) (© 2014, Elsevier. Used with permission). This method is cost-effective, simple, and effective in producing nanocomposites


http://www.adityainc.com/casting/centrifugal-casting.html

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This ensures an intentional variation in volume fraction (functionally gradient) within the matrix material. A typical example for use of centrifugal casting is in brake rotors, wherein the rotor face is expected to be of high wear resistance when compared to the hub. When processed by regular casting methods, difficulties in machining are encountered due to the high hardness of the composite; these are eliminated by centrifugal casting process (Chawla and Chawla 2006).

2.3.1.3 Squeeze Casting/Infiltration Process

This process involves the infiltration of a molten alloy into a ceramic fiber/particle preform followed by solidification (Ghomashchi and Vikhrov 2000). The introduc-tion of molten metal into a preform could be achieved either through pressureless infiltration or by infiltration under pressure. In pressureless infiltration, ceramic fiber bundles are first placed in the die. The molten metal is then poured on to it and allowed to solidify. The solidified composites are then hot pressed to achieve 100 % density. The initial infiltration occurs without any application of external pressure and the wettability of fibers ensures efficient infiltration. The pressure infiltration process can be employed in two different ways, namely via gas infiltration and pres-sure infiltration. In gas infiltration, vacuum or inert gas atmosphere is utilized to bring forth infiltration. Advantages of this method include increase in the wettability due to the increased surface activity of reinforcement in vacuum environment,

Fig. 2.3 Schematic of centrifugal casting process (http://www.adityainc.com/casting/centrifugal- casting.html 2014). Last accessed 23 Oct 2014




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elimination of gas entrapment or porosity and achieving near-net shaped compo-nents. Its main disadvantages are segregation of phases and reaction between matrix/fiber due to the slow nature of the process.

The squeeze infiltration process involves the infiltration of molten metal into aceramic preform using hydraulic pressure (Fig. 2.4).By this method, the drawbacks

Fig. 2.4 Schematic of squeeze casting/infiltration process


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of phase segregation and matrix/interface reaction encountered in gas infiltration can be eliminated due to the application of hydraulic pressure (as it increases the solidi-fication rate). To note, preforms are usually prepared using ceramic reinforcements having aspect ratios (e.g., alumina short fibers). For nanoscale reinforcements, CNTs are used as preforms for they have dimensional anisotropy (i.e., aspect ratio) (Uozumi et al. 2008). However, preparation of CNT preforms is a major challenge. In both cases, an improperly made preform can cause local inhomogeneous distribution of reinforcement fibers causing large variation in the volume fraction within a solidified composite. Also, in case that the preform is not well prepared (such as insufficient binder), it has a tendency to break during the application of squeeze pressure.

Squeeze Casting/Infiltration Process Parameters

The quality of composite castings produced in this process largely depends upon thecontrol of the processing parameters discussed below.

1. Metal casting temperature: The temperature at which the metal is poured into the die cavity plays an important role on the casting quality and die life. The castingtemperature depends on the liquidus temperature, the freezing range of the alloy,and the die configuration. Low casting temperatures result in inadequate fluidlife and incomplete die fill. On the other hand, high casting temperatures would force the metal to penetrate between the die and punch leading to metal flash or jamming of the tooling.

2. Tooling temperature: High tooling temperatures would cause surface defects in the castings. In some cases, welding may occur between the casting and the die wall. Low temperature leads to premature solidification, thermal fatigue, and cold laps. Temperatures of ~300 ° C or less are maintained for nonferrous alloys.

3. Melt quality and quantity: In this process, the metal is directly poured into the cavity that has no gating or feeding system. Hence adequate precautions shouldbe taken to ensure that the material is free from any dross or suspended impuri-ties. Due to the absence of a gating system, the process requires precise quantityof metal to be poured into the die cavity, which will otherwise change the casting dimensions.

4. Die coating/lubricant: Die coating serves as a releasing agent and is selected based upon the die material and the alloy composition. A commonly used die coat is water-based colloidal graphite that is sprayed onto the die surface and the punch. At high squeeze pressures, the coating may get stripped from the diesurface causing surface contamination. Hence, precaution should be taken in applying the coating to the right thickness to ensure its desired performance.

5. Temperature for pressure application: The squeeze casting/infiltration process is fullyeffective when the metal is completely in a liquid state (Das and Chatterjee 1981). Hence, the control of time delay between pouring and squeezing is very importantand should be minimized. Else, it may lead to reduction in the melt temperature lead-ing to premature solidification or incomplete infiltration of the preform.


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6. Duration of pressure and pressure level: The applied pressure level depends on the fluidity and the freezing range of the alloy (Das and Chatterjee 1981). It also depends on the component geometry and the alloy characteristics. Pressure in the range of about 30–110 MPa is usually required to eliminate gas orshrinkage porosity. The duration of pressure application depends on the alloy type and the heat transfer conditions. Prolonged holding under pressure would cause die wall cracking and difficulty in the withdrawal of the punch. A high applied pressure may also result in breakage of the preform. Usually, a maxi-mum pressure holding time of approximately 1 s/mm section thickness of cast-ing is employed.

7. Press speed: At high press speeds, the following may occur: metal flash at joints, dilation of die parts and premature solidification at the impact zones due to instantaneous peak pressures. Normally a two-speed action of rapid approach to the metal surface followed by a slower impact speed is adopted.

8. Preform properties: Major factors to be considered are the full infiltration of the preform (without damage or distortion) and the freezing of metal in the preform with good interfacial contact between the matrix and the fiber. Insufficient infil-tration produces shrinkage pores that deteriorate the properties of the composite. Further, the degradation of fibers by the melt depends on the exposure time, squeeze infiltration temperature, wettability, and the bonding between the fiber/matrix at the interface. In addition, nonuniform distribution of fibers (clustering) and large increase in applied pressure result in fiber degradation.

2.3.1.4 DMD Technique

The DMD technique is a liquid-state processing method, which has the combinedadvantages of both gravity die casting and spray forming processes (Gupta and Sharon 2011). Unlike in the spray process, DMD employs higher superheat tem-peratures and lower impinging gas jet velocity. It is usually employed for particle reinforcements, and is quite successful for reinforcing nanoparticles. The processinvolves stirring of reinforcement particles with a predetermined stirring velocity and time using an impeller when the metal/alloy is in the molten state. The resulting composite slurry is then made to exit from the bottom of a crucible, followed by disintegration of the melt by jets of inert gas at a superheat temperature of 750–850 °C (depends on the alloy) and is finally deposited onto a metallic substrate (Fig. 2.5). The disintegration process of the composite melt ensures higher solidifi-cation rate and fine-grained structure.

Generally during Mg materials production via conventional methods such as gravity die casting, critical issues encountered are:

1. Presence of oxides in the final product (molten Mg is highly oxidizable in nature) 2. Retention of reinforcement particles in crucibles (i.e., most of the reinforcement

particles are denser than Mg, and hence tend to settle at the bottom of crucibles)


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These together give rise to impurities, insufficient reinforcement volume frac-tion, and nonuniform reinforcement dispersion. Given that DMD is a bottom- pouring technique, it ensures:

1. Effective elimination of oxide entry into deposited products 2. Complete utilization of the reinforcement 3. Higher solidification rates due to disintegration of molten metal by inert gas

Salient features of the process are:

• Combined advantages of casting and spray forming processes.• Eliminates the requirement for separate melting and pouring units.• Removes oxides and slag/dross and least metal wastage.• Flexibility of addition/incorporation of alloying/reinforcing particle elements.• Eliminates the retention/settling of the reinforcements in the crucible.• High process yield and gives rise to fine-grained materials with minimal porosity.

Thermocouple

Crucible Lid

Stirrer

Resistance Furnace

Graphite Crucible

Pouring Nozzle

MoltenSlurry

Argon-filledChamber

Mold

Substrate

Ar Ar

ArgonGasTank

Motor

750 °C

Fig. 2.5 Disintegrated melt deposition technique. This method has the advantage of both conven-tional die casting and spray forming processes. It is very suitable to produce Al- and Mg-based nanocomposites


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DMD is a primary process, after which a secondary process such as extrusion is usually employed. Following this process, light metal composites/nanocomposites have been successfully produced at laboratory scale. To make components/products directly using DMD, the disintegrating gas velocity and the distance between the melt exit stream and a substrate need to be standardized according to the requireddimensions of the final product/component. This being arduous, the process has remained suitable to produce composite ingots, which can then be used as precur-sors for making wrought products.

2.3.1.5 Ultrasonic-Assisted Casting

The ultrasonic-assisted casting method is effective in mitigating particle cluster for-mation in composites that occurs (nanocomposites in particular) due to the low wettability and high tendency of agglomeration of nanoparticles (Donthamsetty et al. 2009; Mula et al. 2009). Agglomeration is usually encountered in conventional stirring methods such as mechanical stirring/vortex methods. In contrast, the ultrasonic- assisted method employs subjecting the melts with ultrasonic waves (fre-quency range: 18–20 kHz) during or after adding a reinforcing phase. This is fol-lowed by casting. A schematic of the setup is shown in Fig. 2.6 (Yang et al. 2004).

High-intensity ultrasonic waves can generate transient cavitation and acoustic streaming in liquids (Abramov 1994). Acoustic streaming causes pressure gradient within the bulk of molten metal that produces stirring effect. In cavitation, cyclic high-intensity ultrasonic waves induce the formation, growth (during the negative pressure cycle), pulsating, and collapsing (during the positive pressure cycle) of tiny bubbles in the liquid phase. At every cavitation cycle, bubbles implosively collapse

Fig. 2.6 Schematic of Ultrasonic-assisted casting (adapted from Yang et al. 2004) (© 2004, Elsevier. Used with permission). The method is very effective to homogenously disperse nano- sized reinforcements in light metal matrices


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in less than 10−6 s, producing micro “hot spots” that can reach temperatures of ~5,000 °C, pressures of ~1,000 atm, and heating/cooling rates >1,010 K/S during microseconds transient (Suslick et al. 1999). During cavitation of the composite melt, the air entrapped in voids of particle clusters serves as nuclei for cavitation, which is strong enough to break the clusters thereby providing uniform dispersion. The high pressure and temperature developed also removes gases/impurities and enhances the wettability of the reinforcement with matrix. This method is extremely successful in producing composites with uniform dispersion of nano- reinforcements, at laboratory scale. For any large-scale production, it requires to up-scale the probesize with higher source power to ensure its effect over large volume of melts.

2.3.1.6 High Pressure Die Casting

Among the traditional liquid processes, aside from stir casting technique, high pres-sure die casting (HPDC) has also been used for LMMC production. In general, the process enables obtaining more precise components as compared to gravity and low pressure die casting methods. The molten metal is forced into the die cavity under pressure, and both filling speeds and solidification rates are particularly high. For this reason, HPDC method is characterized by fast cycle times, which may range from seconds to several minutes, depending on the size and wall thickness of the casting. On the other hand, the process inevitably induces gas entrapment due to the highly turbulent flow of metal in the cavity (Wang et al. 2011; Long et al. 2012). Very few works have been reported on the application of HPDC for the production of MMCs, especially in the case of nanocomposites. HPDC was applied by Li et al. (2010) to manufacture CNT-reinforced Al-based composites. The composites exhibited an increase of both tensile stress and elongation to failure when compared to the unreinforced alloy.

2.3.2 Solid State Processes

Solid state processes involve production of materials in solid state form (such as powders in powder metallurgy (PM)) (Suryanarayana and Al-Aqeeli 2013). The first step in powder metallurgy method is the blending/mixing of metal matrix and reinforcement powder, which is conducted using a ball-milling machine (without milling media, e.g., steel balls). In some cases, reactive mixing is employed, in which reinforcement particles and/or the alloying/catalyst element are milled together (with a milling media at selected rpm and time duration). This step is undertaken so as to improve the wettability. The powders so blended are then com-pacted usually by cold pressing (called “green compacts”), and in some cases by using hot/vacuum hot pressing. The green compacts are then further sintered using either of the methods of direct sintering using resistance furnace, microwave- assisted sintering, spark plasma sintering, hot extrusion, severe plastic deformation, etc.


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PM methods can be used for producing near-net shaped composite components and large-batches of small components. As the various PM processes are exhaustive, only a brief mention on the process, microstructure, and mechanical properties are presented in this chapter. For more detailed information, the readers can refer to Suryanarayana and Al-Aqeeli (2013), which deals exclusively on solid state pro-cesses and properties of composites. Major advantages of PM techniques are:

1. Minimal wettability-related problems 2. Incorporation of higher volume fraction of reinforcements3. Possibility of using reinforcement-metal combinations not viable by liquid-state

routes

Some of the drawbacks include cost-ineffectiveness, oxidization of powders which demands inert/protective gas conditions, and high porosity content (inevita-ble in powder metallurgy techniques) that leads to degradation in properties, espe-cially the drastic reduction in ductility. As a part of PM process, sintering of green compacts is usually performed using conventional resistance furnaces. Other tech-niques such as microwave-assisted rapid sintering and spark plasma sintering (SPS)have also been recently developed and are being used to produce composites and have been briefly mentioned.

2.3.2.1 Mechanical Alloying

In the traditional PM process, the aim of blending is simply to mix the powders without inducing material transfer between the mixed components. It is possible, however, to perform a high energy mixing through milling media, so as to eliminate the voids between the matrix and the reinforcement powders through a solid state bonding (Ye et al. 2005). For example, in mechanical alloying (MA), matrix and reinforcement are alloyed together by inducing cold welding, fracturing, and re- welding of the powder particles (Suryanarayana 2001, 2011; Suryanarayana and Al-Aqeeli 2013). The strengthening of metallic alloys is achieved through grain size refinement and dispersion of particles. During the process, a small quantity of thebase powder is loaded into a sealed container, together with the grinding media, then blended through agitation at high speed for a predetermined amount of time (Fig. 2.7a). As the kinetic energy of the grinding balls depends on their mass and velocity, dense materials such as stainless steel or tungsten carbide are preferred to ceramic. Main process parameters, influencing the quality of the composite, com-prise ball-to-powder ratio (BPR), time and rotational speed of milling. After being milled, powders are compacted, degassed, and consolidated.

A process control agent (PCA, usually referred to as lubricant or surfactant) is usually added while blending the powders, aimed to minimize the effect of cold welding and consequent formation of large powder clusters. Methanol, stearic acid,and paraffin compounds are used for this purpose (Witkin and Lavernia 2006). During continuous severe plastic deformation, refinement of the internal structure of the powders to the fine scales may occur, resulting in the production of


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nanostructured powders (Fecht 1995; Suryanarayana 2011). In the entire process, contamination of the powders must be carefully controlled. Possible sources of con-tamination are the milling tools, milling atmosphere, as well as the PCA. During consolidation, impurities may influence microstructural evolution and grain growth, leading to a possible decrease of mechanical properties of the resulting composite. In some cases, the milled powders obtained from mechanical alloying are also employed as reinforcing particles for casting processes.

2.3.2.2 Reaction Milling

Mechanical alloying can also be accompanied by a solid state reaction, aimed to produce fine dispersion of oxides, nitrides, and carbides in the light alloy matrix. In this case, the process is usually defined as reaction milling (RM). In order to allow the reaction to occur, the PCA can be absent or a suitable milling atmosphere can be used to introduce reagents, i.e., oxygen, argon, nitrogen, or simply air (Suryanarayana 2001). Sometimes the PCA could be itself part of the reaction process.

2.3.2.3 Cryomilling

In cryomilling, the milling phase is carried out at cryogenic temperatures (Fig. 2.7b) or, in some cases, within a cryogenic medium, as liquid nitrogen (Witkin andLavernia 2006). A PCA (e.g., stearic acid) can be used to avoid severe sticking. During traditional milling process, the temperature increases due to the attrition. As a result, severe recovery and recrystallization of fine microstructures occur.

Fig. 2.7 (a) Mechanical alloying processing technique: milling action on the powders (adapted from Suryanarayana 2001) (© 2001, Elsevier. Used with permission). (b) Schematic of cryomilling process (adapted from Witkin and Lavernia 2006) (© 2006, Elsevier. Used with permission)


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On the contrary, when cryomilling is applied, recovery and recrystallization are suppressed by the extremely low milling temperature, enhancing the beneficial effects of mechanical milling and leading to finer grain structures and more rapid grain refinement. As a result, grain structures of the order of nanocrystallites may be obtained. Moreover, detrimental chemical reactions between matrix and reinforce-ment are also suppressed at such low temperatures.

2.3.2.4 Microwave Sintering

Microwave heating is a volumetric heating process that involves the conversion of electromagnetic energy into thermal energy (Gupta and Eugene 2007). In conven-tional sintering processes, during sintering, the thermal energy is transferred from the outer surface of materials to their inner surface. In contrast, in microwave sinter-ing, heat is generated from within the materials. It is then radiated outwards due to the penetrative power of microwaves. Due to this phenomenon, during microwave sintering, higher temperatures exist at the core of materials whereas their surfaces experience lower temperatures (thermal gradient), which results in variation of microstructure and hence the properties. To avoid such an occurrence hybrid micro-wave heating referred to as the “bi-directional hybrid microwave-assisted rapid sin-tering” has been developed (Fig. 2.8) (Tun and Gupta 2009).

In this process, microwave susceptors such as SiC particles/rods are used to assist in the reduction of thermal gradient during sintering. The compacted metal/composite powder billets are placed in the inner crucible and SiC powder is placed in between the inner and outer crucibles. As SiC powder absorbs microwave readily, it heats up quickly. This provides radiant heat that can in turn externally heats thecompacted billets. In addition, the compacted billets themselves absorb microwave and get heated from within/internally. Thereby, uniform heat is experienced along

Controlled atmosphere chamber

InsulationMicrowave oven

Lid

Outer ceramic crucible

Inner ceramic crucible

SiC powder susceptor

Compacted sample

Dummy block

Turntable

Fig. 2.8 Bi-directional hybrid microwave-assisted rapid sintering (adapted from Tun and Gupta 2009) (© 2009, Elsevier. Used with permission). This technique is advanced than conventionalsintering in terms of distribution of heat during sintering and energy efficiency


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the entire section of a specimen, thus reducing any core-to-periphery thermal varia-tion (Gupta and Eugene 2007). Due to this reason, high sintering temperatures (~620–650 °C) can be generated within a short period of time (12–14 min), that are almost close to the melting points of Al and Mg, by the virtue of which enhanced wettability and reduced porosity can be achieved. The bi-directional hybrid microwave- assisted rapid sintering has been successfully used for nanocomposites production, especially for Mg materials. Advantages of this process include:

1. Rapid heating rates 2. Low sintering time due to which in most cases (even for Mg) the process does

not require inert atmosphere 3. Lesser porosity when compared to conventional sintering 4. Fine microstructures and improved mechanical properties (Gupta and Eugene 2007)

2.3.2.5 Spark Plasma Sintering

As mentioned before, the main drawback of conventional sintering is the occurrence of high porosity in a product; likewise, when the green compacts are hot pressed, hot extruded, or hot isostatically pressed, it often results in matrix grain growth that weakens mechanical properties. In this regard, SPS, also known as field assisted sintering (FAST), is an effective nonconventional sintering method for obtaining fully dense materials with refined grain size (Saheb et al. 2012). In SPS, the densi-fication is facilitated by the use of a current. A pulsed DC current is directly passed through a graphite die and composite powder compact. Joule’s heating effect plays the role in densifying powder compacts achieving near theoretical density. In SPS, the heat generation is internal in contrast to the conventional hot pressing (where heat is provided by external heating elements). This facilitates high heating rates (up to ~1,000 K/min) and hence the sintering process is very fast (within a few min-utes). The speed of the process ensures densification of powders while avoiding coarsening that occurs in standard densification routes (Saheb et al. 2012; Sairam et al. 2013). Figure 2.9 shows a SPS process setup (Saheb et al. 2012). A detailed review on SPS process can be found in ref. Saheb et al. (2012).

2.3.3 Semisolid State Processes

In semisolid processes, incorporation of particles in matrix metal is conducted when a matrix is in semisolid state (slurry), as it:

1. Facilitates uniform dispersion of reinforcements 2. Eliminates settling of denser reinforcement due to slurry state 3. Eliminates porosity (that usually occurs due to voids generated at the molten

metal/particle interface)


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Semisolid processes are of two types, namely thixo-process/casting and rheo- processes (compocasting). In thixo-casting process, addition of particles is car-ried out in matrix molten state followed by agitation/vortex stirring of particles in slurry state, and then casting. In rheocasting, incorporation of nanoparticles is done in the semisolid state of a matrix. A combination of techniques such as semisolid pro-cessing (stirring) assisted by ultrasonic vibration has been utilized to achieve effective dispersion when using nano-reinforcements (Nie et al. 2011; Kandemir et al. 2012).

2.3.3.1 Thixo-processing

In thixo-processes a proper solid feedstock is reheated and partially melted. The base material is generated by allowing a liquid melt to partially solidify under con-trolled conditions (low superheat and rapid cooling, usually combined with signifi-cant convection in the liquid), so as to induce the formation of crystals in the slurry.The feedstock may be produced in a variety of ways such as with mechanical stir-ring during solidification as in rheocasting, continuous casting combined with magneto- hydrodynamic stirring for grain refining and ultrasonic treatment for grain refinement. Other methods to prepare fine-grained nondendritic material are by spray casting and low-superheat casting processes. Many of the processes employ intense chemical inoculation to maximize the efficiency of the above mentioned processes, particularly magneto-hydrodynamic stirring and low-superheat casting. Usually, the semisolid material is then injected into hardened steels, dies as final stage process. The advantages of semisolid processing include low shrinkage and porosity, nonturbulent filling and lower processing temperature (Abbasipour et al. 2010; Heinrich and Gonasagren 2012).

Upper electrode

Lower electrode

Thermocouple

Vacuumchamber

Powder

Lowerpunch

Upperpunch

Die

On-

off D

C p

ulse

gene

rato

r

Oil pressure,pneumatic system

Controller/computing

TemperaturePressure

Current-voltageVacuum

Longitudinal displacement

P

P

Fig. 2.9 Spark plasma sintering (SPS). This process is effective for producing fully dense light metal nanocomposites (adapted from Saheb et al. 2012) (©2012, Hindawi. Open access)


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2.3.3.2 Rheo-processing

Unlike the thixo-process, in the rheo-processes a special feedstock is not required,and the semisolid slurry is generated starting from the liquid state by cooling themolten metal during the casting process itself. In common for all rheo-processes is that they are easier to implement in a foundry as they involve standard equip-ment for melting, transport, treatment, degassing, and handling. The key differ-ence between the various approaches is in the slurry-making process, where great efforts are being made to create a robust on-demand slurry-making capability. Among rheo-processes techniques, it is worthwhile mentioning the New RheoCasting process (NRC), which relies on a cooling slope to generate the initial slurry (Yasunori et al. 1996). In this process, the molten metal is poured at low superheat (about 10 K) onto the side of a holing cup and a large amount of very small crystals are formed. The slurry is then held for a pre-set time in the cup, allowing the crystals to grow and spherodize without additional shearing or stir-ring. Just before pouring the temperature of the slurry is homogenized (Kaufmann and Uggowitzer 2001). As a variation to the precedent, the Hong-Nano Casting method (H-NCM) (Hong and Kim 2006), uses an electromagnetic field in the pouring and cooling stages. This modification helps in homogenizing the tem-perature and increases the overall heat transfer, resulting in fast cooling and copi-ous nucleation—approximately 1,000 times higher than in the NRC process. Further, the Rheo Die Casting process (RDC), also known as Twin Screw Rheo Moulding (TSRM), involves the use of twin screws for mixing, providing a high amount of shearing. The molten metal is cooled at a controlled rate. The high level of shear is thought to break oxides into small, round particles which are well dis-persed in the entire cast component. The slurry may be generated by letting the melt passing through a conversion reactor (a cooled copper or iron block with a twisting channel inside, causing the melt to cool and partially solidify under shear) in the so called Continuous Rheo-conversion Process (CRP). Other pro-cesses developed so far for semisolid metal processing are the Sub-Liquidus (SLC)and Semi Solid Rheo (SSR) casting processes, GISS process, Rapid Slurry Forming (RSF), Semisolid Metal (SSM), and ATS processes.

2.3.3.3 Compocasting

Compocasting route, is a rheocasting process that involves the injection of rein-forcement particles into semisolid state alloys (Fig. 2.10). Compocasting is gener-ally thought to be a processing route allowing to obtain quite uniform distribution ofreinforcing particles, as well as to enhance particle wettability (Kamali Ardakani et al. 2014). It has been reported that the primary solid particles which are formed in the semisolid slurry are able to mechanically entrap the reinforcing phase and to prevent their gravity segregation, as well as to reduce their agglomeration tendency (Naher et al. 2005). Moreover, the lower porosity which is usually observed in experimental studies is attributed to the better wettability between the matrix and the reinforcement particles as well as the lower volume shrinkage of the matrix


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alloy. Despite these advantages, some agglomerates, inevitably induced by the high surface-to-volume ratio, as well as to Van der Waals interactions are still reported (Sajjadi et al. 2012).

2.3.4 Other Processes

Alternative processes to incorporate reinforcement particles into metal matrices to form bulk composites are also being developed. Friction stir process (FSP) is one such method, which is based on friction stir welding. Other processes are accumulative roll bonding (ARB) and equal channel angular pressing methods. These are severe plastic

Resistancefurnace

Stirrer

Stopper

Baffle

Coatedinjector

10 cm

Mold

11 c

m

Valve

The

rmoc

oupl

e

Ar

Powder

Fig. 2.10 Compocasting experimental setup (adapted from Abbasipour et al. 2010) (© 2010, Elsevier. Used with permission)


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deformation (SPD) process used to produce fine-grained materials. Research on full utilization of these methods to produce LMMCs is in the developmental stage.

2.3.4.1 Friction Stir Processing

FSP is based on friction stir welding. It was initially used to produce surface- reinforced composites (Sun et al. 2012). During FSP, a rotating tool with a shoulder and a pin is plunged into the surface of a work piece (desired base matrix) with grooves filled with particles of required volume fraction (Fig. 2.11) (Lee et al. 2006). When the tool rotates, it covers the region of interest. In recent years, efforts are being made to use this process as an alternative route to incorporate nanoparti-cles into metal matrices to form bulk nanocomposites (Lee et al. 2006; Sun et al. 2012). However, obtaining uniform dispersion of nano-sized reinforcements still remains a challenge.

2.3.4.2 Equi-Channel Angular Pressing

Equi-Channel Angular Pressing (ECAP) is a SPD which involves processing/extru-sion of a metal billet through an angled (typically 90°) channel (Segal 1999; Balog et al. 2013; Sklenicka et al. 2013) as shown in Fig. 2.12. To achieve optimal results,

Fig. 2.11 Friction stir process (FSP): (a) cutting groove and inserting nanoparticles, (b) using a flat tool for surface repair, (c) applying a tool with a fixed pin that causes FS, and (d) conducting multiple FSP passes (adapted from Lee et al. 2006) (© 2006, Elsevier. Used with permission)


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this process is repeated several times, changing the orientation of a billet with each pass. This produces uniform shear throughout the bulk of material. The process is widely used for fine-grained/ultrafine-grained materials, and has also been adopted to produce nanocomposites as it can provide fine grains, uniform distribution of reinforcements, and elimination of porosity (Segal 1999).

2.3.4.3 Accumulative Roll Bonding

ARB is a solid state method in which composites can be produced in the form of sheets. This process is also an SPD process. In this process two sheets of the same material are stacked, heated (to below their crystallization temperature), and rolled to bind the sheets together (with up to 50 % thickness reduction) (Fig. 2.13) (Saito et al. 1999). The bonded sheet is then cut into half, the two halves are stacked together, and the entire process is repeated all over again several times. Compared to other SPD processes, ARB does not require specialized equipment or tooling, butonly a conventional rolling mill. During the ARB process care should be taken to

Sample

Plunger

Die

Pressed sample

P

Φ

Ψ

Fig. 2.12 Equi-channel angular pressing (ECAP). The die hasan internal angle of 90° and an outerarc of curvature of ~20° (adapted from Sklenicka et al. 2013) (©2013, InTech. Open access)


31

thoroughly clean the surfaces before rolling them together so as to achieve good bonding (Saito et al. 1999).

2.4 Strengthening Mechanisms in LMMCs

For engineering materials design and selection, yield strength and ductility of a material are the most important parameters. These properties in LMMCs are enhanced by the strengthening mechanisms (Arsenault 1983; Nardone and Prewo 1986; Dieter 1988; Miller and Humphreys 1991; Brown and Stobbs 2006; Zhang and Chen 2006) mentioned below:

1. Hall–Petch effect (grain refinement): Reinforcement (fibers/particles) when introduced in a molten matrix can act as heterogeneous nucleation sites during solidification, thus giving rise to refined and equiaxed grains. When secondaryprocesses such as hot extrusion are undertaken, grain growth during recrystalliza-tion is hindered due to the grain boundary pinning effect by the reinforcements.

Surface treatment

Degreasing,Wire brushing

Cutting

Roll bonding

Heating

Stacking

+

+

Fig. 2.13 Accumulative roll bonding (ARB) process, based on intense plastic straining (adapted from Saito et al. 1999) (© 1999, Elsevier. Used with permission). This process can produce nano-composite in the form of sheets


32

2. Orowan strengthening: Strengthening arising due to obstacle posed by closely- spaced hard particles to the dislocation motion. Highly dispersed reinforcements significantly increase the yield strength.

3. Enhanced dislocation density: Usually, matrix and reinforcement materials have different CTE, due to which thermal stresses are generated during processing (e.g., during cooling in casting process). Such stress levels are large enough to increase dislocation density, especially at the matrix/reinforcement interface resulting in higher yield strengths.

4. Load bearing effect: Under external loads, matrix transfers the load to the rein-forcement. Reinforcements have strong interfacial bonding with matrix and so good load bearing capacity is exhibited.

2.4.1 Strength Prediction

Theoretical yield strength of metal matrix composites can be estimated using the relation formulated by Nardone and Prewo (Nardone and Prewo 1986), which is based on a modified shear-lag model extended to the case of particulate composites, and is given by (2.2) as follows:

s syc ym p p= +( )éë ùû + éë ùû{ }V S V1 2 1/ -

(2.1)

This equation takes into account the load transfer on particle from matrix, where σym is the yield stress of unreinforced matrix, Vp is the particle volume fraction, S is the aspect ratio of particle. Equation (2.1) does not take into account, the parameters involved in strengthening of composites. Strengthening of reinforcement particles on the yield strength (σym) of metallic matrix arises due to: (1) Orowan strengthen-ing (ΔσOrowan), the stress increase needed to move a dislocation through an array of impeding particles, (2) stress contribution due to statistically stored dislocations introduced by the thermal expansion mismatch between the matrix and reinforce-ment (ΔσCTE), and (3) generation of geometrically necessary dislocations (GND) to accommodate the plastic deformation mismatch between the matrix and particles (ΔσGND). When several strengthening effects are combined, the rule of addition and the root of the sum of the squares may be used to predict the yield strength, as men-tioned below in (2.2). The strength of a reinforced matrix is given by,

s s Ds

my my1 = +

(2.2)

where smy1

and σmy are the yield strength of reinforced and unreinforced matrices respectively. Δσ represents the total increment in yield stress of reinforced matrix and is estimated as follows,

Ds Ds Ds Ds= ( ) + ( ) + ( )Orowan CTE GND

2 2 2

(2.3)


33

1. ΔσOrowan, can be estimated using (2.4a) and (2.4b),

Ds f mOrowan m m= × b L/ (2.4a)

where ϕ is a constant of order 2, μm and bm are the shear modulus of metal matrix and its Burgers vector, respectively (for Al, μm: 26 GPa, bm: 0.286 nm). L is the interparticle spacing of the second phase particles, which is given by

L DV

=æ

èçç

ö

ø÷÷

p6

1 3

p

/

(2.4b)

where Vp and D are the volume fraction and diameter of particles, respectively. 2. ΔσCTE can be estimated using (2.5a) as given below:

Ds h mCTE m m= × b r (2.5a)

where,

rDaD

=-( )

12

1

TV

b D V

p

mp

(2.5b)

Here, η is a constant of order 1, and ρ is the dislocation density. Δα is the differ-ence in thermal expansion coefficients between the matrix. ΔT is the temperature change from the processing temperature to room temperature.

3. ΔσGND is calculated using (2.6),

Ds b m eGND m p

m m= × ( )V b D/

(2.6)

β is a geometric factor with a numerical value ~0.4, and εm is the plastic strain of the matrix.

2.4.2 Fracture Mechanisms

The fracture mechanisms of MMCs (such as in particulate composites) is dependent on interfacial strength and particle strength. If the interface strength is stronger than the particle strength (as in precipitation hardened composite), then the particle frac-ture occurs before the interface. The failure of the composite hence occurs by shear localization of fractured particles, giving rise to brittle fracture features. In short fiber composite, this type of damage gives rise to fiber cracking and fiber breakage. On the other hand, when the particle strength is higher than the interfacial strength,


34

it results in decohesion of particles due to void nucleation and growth in the matrix. This would be evident by ductile shear failure in the fracture surface of the compos-ite. In short fiber composites, such a phenomenon leads to fiber pull-out in fibers that are oriented along the loading direction. Fracture also occurs by particle crack-ing and void formation in the matrix within clusters of particles. Particle fracture is more prevalent in coarser particles due to the following reasons: (1) the larger the particle the more it will be loaded by conventional fiber loading and end loading mechanisms and (2) coarser ceramic particles will have a higher probability of con-taining fracture initiating defects. Further, particle clusters give rise to a triaxial state of stress causing severe reduction in ductility resulting in catastrophic failure of the composite components.

2.5 Microstructural and Mechanical Properties of LMMCs

The attributes of Al and Mg composites are a combination of high specific stiffness, excellent wear resistance, and the potential for relatively low-cost conventional pro-cessing. Al and Mg composites have been under development for many years during which time a vast number of different types of reinforcement have been attempted with varying degrees of success. These include continuous fibers, short fibers, whis-kers, and particles. Many different matrices have been tried over the years and these have a bearing on some of the properties that can be achieved in the composite. Tailoring the mechanical properties with respect to a specific application can also be achieved by varying the type and amount of reinforcement, the choice of matrix alloy, and the composite processing route. All these factors are interrelated and should not be considered in isolation when developing a new material. Among the various types of Al, Mg MMCs, the advantage of particulate reinforced composites are due to processing, wherein conventional metal manufacturing methods and machining techniques can be used (Lloyd 1994). Given that the literature available on Al and Mg composites/nanocomposites are exhaustive, only representative microstructures/mechanical data are presented here.

2.5.1 Al-Composites

Different processing methods have been used to produce various micro- and nano- reinforced Al-composites. As mentioned earlier, literature review indicates that Al2O3, SiC, B4C, BN, AlN, etc. in the form of particles/whiskers/fibers have been incorporated in various Al-matrices. Daoud and Reif (2002) studied the influence of Al2O3 particulates on the aging response of A356 alloy prepared using the vortex method. In these composites, the formation of MgAl2O spinel on the interface caused depletion of Mg in the matrix, which eventually resulted in less age- hardening behavior (Fig. 2.14). The role of grain refinement treatment and the


35

degree of grain refinement was studied in Al-7Si-10 % Vf SiC particulate reinforced composites prepared by vortex method (Garcia-Hinojosa et al. 2004). While com-plete dewetting of the SiC particles and rejection from the melt was observed when K2TiF6 salt was used as grain refiner, the wettability improved with Al-6Ti and Al-5Ti-1B master alloy grain refiners. Significant grain refinement and uniform dis-tribution of the reinforcement was achieved. The mechanical properties of Al–AlN composites prepared by squeeze infiltration method (Lii et al. 2002), showed that the hardness and the compressive strength increased, with a significant reduction ductility. It was identified that the crack growth path was prominent along the Al–Al, AlN–AlN, and Al–AlN interfaces. The tensile property data of Al–AlN com-posite showed increased yield and tensile strength, accompanied by a drastic reduc-tion in ductility both at room and high temperatures (Fig. 2.15) (Zhang et al. 2003). Vicens et al. (2002) studied the interfacial phenomena of 2024, 6060, 5754 Al-matrices with, 45 % Vf AlN, produced by squeeze infiltration. TEM investigation

Fig. 2.14 SEM micrographs showing the morphological changes in the surface of Al2O3 particles in A356 alloy during melt processing (a) before addition into the molten A356 alloy (b) after addi-tion with A356 alloy in the composite and (c) magnified view of the reaction products (adapted from Daoud & Reif 2002) (© 2002, Elsevier. Used with permission)

400

a b100

80

60

40

5

0400

300

300

200

200

Temperature,°C

Str

engt

h, M

Pa

100

1000

0 400300200

Temperature,°C

Elo

ngat

ion,

%

1000

0.2% proof stress, aluminum matrixAluminum matrix

0.2% proof stress, AINp/AI compositeAINp/AI composite

UTS, AINp/AI composite

UTS, aluminum matrix

Fig. 2.15 (a, b) Room and high temperature tensile properties of Al–AlN composites. Significant improvement in (a) strength properties is achieved when compared to the unreinforced matrix, however (b) with a drastic reduction in ductility (adapted from Zhang et al. 2003) (© 2003, Elsevier. Used with permission)


36

revealed MgAl2O4 spinel crystals in 6060 and 5754 composites. In these compos-ites, the magnesium element of the matrix reacts with a very thin alumina layer which is deposited on the AlN surfaces during the liquid infiltration step. The 5754composite exhibited a stronger reaction leading to the formation of MgO phases with the spinel that resulted in degradation of the mechanical properties. Also, microcracking occurred mainly at the interface between two AlN grains where MgO and MgAl2O4 precipitated (Fig. 2.16a, b). Interfacial studies on pure Al rein-forced with SiC, Al2O3, and B4C show that while an interfacial layer is formed in Al–SiC composites, none were observed in the other two composites (Shorowordi et al. 2003). However, in the Al–B4C composite, two secondary phases consisting of carbon, aluminum, and boron were observed which based on the fracture studies give rise to a strong interface due to effective interfacial bonding.

Investigation on squeeze cast Al–Mg–BN-reinforced composites showed the for-mation of interfacial AlB2 phase and in situ AlN phase which resulted in significant grain refinement (Lee et al. 2002). The interfacial reaction products of the Al/B4Cp composite fabricated by the pressureless infiltration method were analyzed by Lee et al. (2001) using XRD, SEM, EDS, AES, and TEM. From the SADP and CBED analysis, it was identified that AlB2, β-AlB12 (Al3B48C2), AlB10 (AlB24C4), and Al3BC were formed as interfacial reaction products (Fig. 2.17). Similarly, interfacial study on a SiCp/Al composite fabricated by pressureless infiltration of an Al–Mg–Si alloy into a preoxidized SiC preform confirmed the presence of MgAl2O4 interfacial reac-tion product (Fig. 2.18), which is expected to improve the wettability between the

Fig. 2.16 (a) TEM image of MgAl2O4 precipitates formed at the interface in Al–Mg–Si/AlN com-posite and (b) MgO and MgAl2O4 interface precipitates giving rise to crack formation between two AlN grains (after compression tests) (adapted from Vicens et al. 2002) (© 2002, Elsevier. Used with permission)


37

SiC particles and the matrix (Zhang et al. 2013). The microstructure of a SiC par-ticulate reinforced Al 2080 alloy was characterized by Evans and Boyd (2003) using focused ion beam microscopy and transmission electron microscopy. A 40 nm thick amorphous interface layer was seen to develop during processing of the compos-ite material by diffusion of Al and Mg into the preexisting SiO2 layer on the SiC particles. During tensile deformation, damage occurs in the near-interface zone of

Fig. 2.17 TEM images of interfacial products between Al matrix and B4C particles (adapted from Lee et al. 2001) (© 2001, Elsevier. Used with permission)

Fig. 2.18 High resolution TEM (HRTEM) image showing the formation of MgAl2O4 at the interface between Al-matrix and SiC particle. The interfacial product is expected to improve the wettability (adapted from Zhang et al. 2013) (© 2013, Elsevier. Used with permission)


38

the matrix, at the particle–matrix interface and within the SiC particles. TEM study of interface in Al–Si–Mg–Zn reinforced with SiC particulates prepared by pressure-less infiltration method (Luo et al. 2001) showed that the reinforcement particles were coated with reaction layers of MgAl2O4 spinel particles, in addition to the Mg2Si intermetallic phases found in the eutectic along the Al-grains. Tekmen et al. (2003) studied the effect of porosity on the mechanical and fracture behavior in Al–Si matrix alloy and composites reinforced with SiC particles of 10 and 20 % Vf in the as-cast state and after extrusion process. It was identified that the increase in reinforcement content increases the overall porosity content; however it reduced to low levels upon extrusion, but effective only for porosities lower than a certain size. While the increase in porosity content decreases both the yield and ultimate tensile strength value, fracture behavior was influenced by large-sized pores rather than the overall porosity content. Tee et al. (2000) investigated in situ TiB2 particulate rein-forced composites produced by the direct addition of Ti and B to molten Al, which resulted in a significant improvement in yield and tensile strengths. However, the formation of additional brittle phase (Al3Ti flakes) during the process caused a large reduction in ductility of the in situ composite. Interfacial reaction studies on Al–TiC composites by Kennedy et al. (2001) showed the formation of large Al3Ti precipi-tates in the bulk of the matrix and Al4C3 blocks at the particle–matrix interface.

For nanocomposites the most commonly used reinforcements are Al2O3, SiC, and CNTs, while other particles such as Si3N4, B4C have also been recently consid-ered. When nanoscale Al2O3 is used as the reinforcement, one major concern is its poor wettability with the Al matrix. As below 1,000 °C, the contact angle between aluminum and Al2O3 is greater than 90° (Schultz et al. 2011). Such poor wetting behavior results in clustering of particles, which eventually leads to deterioration in properties. In this work, Al2O3 nanoparticles were mixed with Mg particles so as to improve their wettability (addition of Mg increases the surface energy of nanopar-ticles). Hossein-Zadeh et al. (2014) studied the effect of heat treating Al2O3 nanopar-ticles prior to its addition into molten Al matrix on matrix/particle wettability. The heat treatment improved the wettability due to grooving at the grain boundary junctions, thereby improving the microstructure and mechanical properties when compared to the nanocomposite having particles without any heat treatment. Mula et al. (2009) prepared pure Al with 2 wt% nano-sized Al2O3 particles by ultrasonic cavitation and observed continuous nano-alumina dispersed zones near the grain boundaries in the nanocomposite. The addition of nanoparticles contributed to the significant increase in the hardness (~92 %), and tensile strength (~48 %). Su et al. (2012) reported improvements in grain refinement, ultimate tensile strength, and yield strength (37 % and 81 % respectively) but with reduction in ductility in Al-2024 nano-Al2O3 composite produced using casting with ultrasonic dispersion. Mazahery and Shabani (2012) prepared nano-SiC-reinforced A356 composite by compocasting and reported improvement in microstructural and mechanical proper-ties. Xiong et al. (2011) reinforced nano-SiC particles in Al-3.0 wt % Mg matrix by pressureless infiltration and investigated the effect of SiC nanoparticles on interfa-cial reactions. Results showed that the formation of MgO at the interface between SiC particulates and molten Al act as a barrier for the diffusion of Si, C, and Al (Fig. 2.19), thereby improving the interfacial properties. Bathula et al. (2012)


39

synthesized Al–SiC nanocomposite with Al-5083 alloy powder and SiC nano par-ticulates using powder metallurgy route followed by SPS. High resolution transmis-sion microscopy (HR-TEM) analysis revealed that the matrix grain size slightly coarsened upon sintering that reduced the ductility (Fig. 2.20). Ahmed et al. (2010) prepared nanocomposites with Al-7075 alloy matrix (Al Zn–Mg Cu alloy) and nano-SiC particles and observed drastic microstructural change (i.e., grain coarsen-ing) that caused deterioration in mechanical properties (~30 % and 40 % reduction in hardness and tensile strength, respectively). With regard to CNTs as reinforce-ments, the major issue that needs to be overcome is their agglomeration during processing. As an example, Esawi et al. (2010) used ball-milling to disperse up to 5 wt% CNT in Al matrix. They reported poor dispersion of CNTs (for >2 wt%) and formation of carbides in the 5 wt% sample. Liao et al. (2010) treated MWCNTs with sodium dodecyl sulfate (SDS) surfactant (as it decreases the van der Waals force between the CNTs) in order to facilitate their easy dispersion in Al-metal. They formed the Al-MWCNT composites via SPS, followed by hot extrusion. It was observed that the mechanical properties of the Al-0.5 wt% MWCNT composites

Fig. 2.19 TEM images of Al–SiC nanocomposite showing interfacial reaction products (a) TEM image of interface in composite with 10 % Vf nano SiC, (b) TEM image of interface in composite with 14% Vf nano SiC, (c) SAD pattern of Mg2Si, (d) SAD pattern of MgO, and (e) SAD pattern of Al4C3 (adapted from Xiong et al. 2011) (© 2011, Elsevier. Used with permission)


40

enhanced noticeably (in both tension and compression). Any increase in wt % of CNTs thereafter decreased the mechanical properties due to agglomeration and insufficient interfacial bonding. The mechanical properties of selected Al-micro/nanocomposites are listed in Table 2.5.

2.5.2 Mg Composites

Mg, unlike Al has h.c.p. structure and possesses limited number of slip systems. Its deformation behavior is governed by both slip and twinning (Dieter 1988). The strengthening mechanisms in Mg materials are strongly influenced by their texture, i.e., preferred crystallographic orientation. Texture evolution is influenced by the form in which the material exists (cast or wrought form), i.e., while casting can give rise to random crystallographic orientation and hence isotropic properties, wrought materials (strong orientation of crystallographic planes) give rise to anisotropic deformation behavior (Wang and Huang 2003). Further, slip/twinning deformation modes are determined by the direction of loading. For instance, in extruded Mg materials, the basal planes are aligned strongly in the extrusion direction, which are highly unfavorable for the basal slip to occur. Due to this reason, the tensile ductility of extruded Mg materials is limited/lower when the testing is carried out in the direction parallel to the extrusion direction (Reed-Hill 1973; Wang and Huang

800

600

400

200

00 2 4 6 8

Strain, %

Str

ess,

MP

a

10 12 14

AI 5083/10wt.% SiC nanocompositeAI 5083 alloy - As receivedAI 5083 alloy - 15 h of milling

Fig. 2.20 Engineering stress–strain curves of 5083 Al-alloy in the as-received, milled (nanostruc-tured), and nano-reinforced condition. Note the large improvement in strength due to nano- reinforcement addition, with little or no ductility (adapted from Bathula et al. 2012) (© 2012, Elsevier. Used with permission)


41

Table 2.5 Mechanical properties of several Al-metal matrix composites/nanocomposites

MaterialMicrohardness (Hv)

Yield strength (MPa)

Ultimate strength (MPa)

% Ductility

Elastic modulus (GPa)

Al–Si (Lloyd 1994) A356 (T6) – 205 280 6 76A356-10%SiCp (T61) – 287 308 0.5 82A356-15%SiCp (T61) – 329 336 0.2 91A356-20%SiCp (T61) – 336 357 0.4 98A380 (F) – 160 320 3.7 72A380-10%SiCp (F) – 245 332 1.0 9.5A380-20%SiCp (F) – 308 356 0.4 114

Al–Mg–Si (Lloyd 1994)6061 (T6) – 275 310 20 696061-10 %Al2O3p (T6) – 296 338 7.8 816061-15 %Al2O3p (T6) – 317 359 7.4 876061-20 %Al2O3p (T6) – 359 379 2.1 986061-15 %SiCp (T6) – 342 364 3.2 916061-15 % SiCp (T4) – 405 460 7.0 986061-20 % SiCp (T4) – 420 500 5.0 1.56061-25 % SiCp (T4) – 430 515 4.0 115

Al–Cu (Lloyd 1994)2014 (T6) – 476 524 13 732014-10 %Al2O3p (T6) – 483 517 3.3 842014-15 %Al2O3p (T6) – 476 503 3.2 922014-20 %Al2O3p (T6) – 483 503 1.0 1012014-15 %SiCp (T6) – 466 493 2.0 1002124 (T6) – 325 472 12 722124-17.8 %SiCp (T4) – 400 610 5–7 1002124-25 % SiCp (T4) – 490 630 2–4 1162124-20 % SiCp (T4) – 4.5 560 7 105

Al–Zn–Mg (Lloyd 1994)7075 (T6) – 505 570 10 727075-15 % SiCp (T651) – 556 601 3 957049-15 % SiCp (T6) – 598 643 2 907090-20 % SiCp (T6) – 665 735 2 105

Al–Li (Lloyd 1994)8090 (T6) – 415 485 7 908090-13 %SiCp (T4) – 454 520 5 1018090-13 % SiCp (T6) – 499 547 3 1018090-17 % SiCp (T4) – 310 460 4–7 1038090-17 % SiCp (T6) – 450 540 3–4 103

(continued)


42

Table 2.5 (continued)




% Ductility


Lii et al. (2002)Al-51 % AlN 92 – – – –Al-60 % AlN 120 – – – –Al-63 % AlN 130 – – – –Al-70 %AlN 150 – – – –

Nanocomposites

Hossein-Zadeh et al. (2014)A356 (C) 53.2 75 213 22 –A356 + 1.0 wt% nAl2O3

(non-HT) (C)63.4 125 265 42 –

A356 + 1.0 wt% nAl2O3 (HT) (C)

70.6 180 375 48 –

Mula et al. (2009)Al (conventional cast) (T) 36 30 62 47 –Al (USC) (T) 51 47 92 36 –Al + 2 % nAl2O3

(USC + CR) ratio 1.1118 – – – –

Al + 2 % nAl2O3 (USC + CR) ratio 2.0

139 – – – –

Su et al. (2012)AA2024 (USC) (T) – 83 154 1.4 –AA2024 + 0.5 wt%

nAl2O3(USC) (T)– 139 194 1.2 –

AA2024 + 1 wt% nAl2O3(USC) (T)

– 154 212 1 –

AA2024 + 1.5 wt% nAl2O3(USC) (T)

– 149 208 0.6 –

AA2024 + 2 wt% nAl2O3(USC) (T)

– 144 203 0.5 –

Mazahery and Shabani (2012)

HB

A356 (SC) (T) 52 122 147 6.0 –A356 (CC) (T) 55 130 157 8.0 –A356 + 0.5 % Vf nSiC

(SC) (T)64 127 227 3.8 –

A356 + 0.5 % Vf n.SiC (CC) (T)

65 138 243 5.6 –

A356 + 1.5 % Vf nSiC (SC) (T)

67 135 240 3.7 –

A356 + 1.5 % Vf nSiC (CC) (T)

70 143 253 5.0 –

(continued)


43





% Ductility


A356 + 2.5 % Vf nSiC (SC) (T)

73 143 260 3.7 –

A356 + 2.5 % Vf nSiC (CC) (T)

75 147 273 4.5 –

A356 + 3.5 % Vf nSiC (SC) (T)

76 146 290 3.6 –

A356 + 3.5 % Vf nSiC (CC) (T)

80 149 295 4.3 –

A356 + 4.5 % Vf nSiC (SC) (T)

75 138 247 3.3 –

A356 + 4.5 % Vf nSiC(CC) (T)

82 151 303 4.1 –

Bathula et al. (2012) Al5083 (unmilled) (C) 148 – 280 13.8 – Al5083 (milled) (C) 250 – 400 3.12 – Al5083 + 10 wt% nSiC

(milled + SPS) (C)– 824 2.5 –

Esawi et al. (2010) Al (T) – – 170 – – Al + 0.5 wt% CNT

(ball-mill + hot ext.) (T)– – 220 – –

Al + 1.0 wt% CNT (ball-mill + hot ext.) (T)

– – 240 – –


– – 255 – –


– – 250 – –

2003) giving rise to large difference between tensile and compression yield strengths (tension-compression yield asymmetry). Favorable orientations of crystallographic planes can give rise to better properties in terms of both strength and ductility, and can be obtained by varying the alloying additions, reinforcement incorporation, and heat treatments/precipitation.

Mg composites are usually reinforced with SiC and Alumina in the form of fibers/whiskers/particulates. Some studies also use carbon fibers; however carbon has poor wetting characteristics with Mg, and in the absence of surface modification of the fiber it results in interface debonding, and is ineffective as reinforcement. Hu et al. (2010) investigated the influence of interfacial properties on the mechanical and in situ fracture behavior of saffil alumina short-fiber-reinforced AE44 (Mg–4.0Al–4.1RE–0.3Mn) composite. Interfacial studies indicated that the SiO2 binder in the preform reacted with molten Mg during infiltration and formed MgO. Further, lamellar and particle intermetallic phases composed of Al–RE


44

(Al11RE3 and Al2RE) were formed on the fiber surface as well as in the matrix. While there was an increase in the elastic modulus, yield, and tensile strengths, a large reduction in ductility (~10 % in unreinforced alloy reduced to ~1 % in the composite) was reported. Fracture studies showed that microcracks initiated in the region of interface leading to interfacial debonding and/or fiber breakage (Fig. 2.21). Contreras (2004) prepared high volume fraction in situ TiC/AZ91D composites using infiltration technique. The room and high temperature tensile properties indi-cated improved strengthening due to reinforcement but with a loss in ductility (brit-tle characteristics). Hu and Wang (2000) prepared ZK51/SiCw reinforced composites by a two-step squeeze infiltration process. Their study revealed that the casting tem-perature determines the interfacial characteristics/reactions and hence the final properties. While a lower casting temperature caused interfacial debonding, a higher casting temperature resulted in melt oxidation and grain coarsening.

Microstructure and interface studies of Mg2B2O5 whisker reinforced AZ91D com-posite by Chen et al. (2010) showed that the Mg2B2O5 whiskers have a twinned struc-ture with the (2 0 2) as the twin plane and growth direction along [0 1 0] and that MgB4O7 particles and globular Mg2Si particles were observed within the Mg2B2O5whisker, and Mg2B2O5whiskers/Mg-matrix interface. Further, MgO and MgB2 phase formed at the matrix-whisker interface during vacuum–gas pressure infiltration process due to the interfacial reaction (Fig. 2.22). Further the rough and uneven interfacial layer would have negative effect on the interfacial bonding. The

Fig. 2.21 Formation of interfacial MgO layer in AE44/Saffil Al2O3 fiber-reinforced composites (adapted from Hu et al. 2010) (© 2010, Elsevier. Used with permission)


45

mechanical properties indicate a dramatic increase of modulus, yield, and tensile strengths in the composite when compared to that of the AZ91D cast alloy, however, at the expense of ductility (large reduction in % elongation). Similar increase in strength and drastic reduction in ductility was observed in hybrid Mg-composite with SiC particles and mullite fibers (Gu et al. 2004). Comparatively, in Mg-hydroxyapatite (HA) composite, not much improvement in strength was observed, but ductility reduced significantly (Gu et al. 2010). Wang et al. (2006) produced Mg–Li matrix composite with 5 wt% Al2Y particulates by stir casting technique and identifiedhomogenous distribution of Al2Y intermetallics particles in the matrix, along with a stable interface with no reaction products. Mechanical properties showed better hardness, modulus, and tensile strength, with ~42 % reduction in ductility. Similar behavior was observed in stir cast Mg–Zn–Ca/SiCp reinforced composites (Fig. 2.23) (Wang et al. 2012) and squeeze infiltrated ZK60A/(SiCw+B4Cp) reinforced hybrid composites (Zhang et al. 1997). The effect of high temperature exposure on the inter-face and tensile properties of AZ91/Al18B4O33w composite fabricated by squeezecasting was studied by Zheng et al. (2004). It was identified that the interfacial region was composed of MgO (10–20 nm thick), which prevented the formation of reaction products between the whisker and the matrix. Improvement in elastic modulus and strength properties was observed with an inevitable drastic reduction in ductility. The effect of acid aluminum phosphate binder was reported in AZ91/SiCw composites fabricated by squeeze casting process (Zheng et al. 2001). For the composites with no binders, the SiCw-AZ91 interfaces are very clean, whereas for the composites

Fig. 2.22 TEM image showing the existence of two interfacial layers in Mg2B2O5 whisker rein-forced Mg composite (a) TEM BF image of Mg2B2O5 whisker. Two dash lines delineate the trace of the twin planes; (b) SAED pattern of the Mg2B2O5 whisker along [010] Mg2B2O5, [1 1 0] MgO and [2-1-10] MgB2 zone axes detected at the connection area of position C (adapted from Chen et al. 2010) (© 2010, Elsevier. Used with permission)


46

with acid phosphate binders, fine, uniform, and discrete interfacial MgO reaction products at the interface exist due to reaction between the binder and molten Mg (Fig. 2.24). The definite crystallographic orientation relationship between MgO and SiCw lead to a compatible interface with good lattice matching, and increase whis-ker–matrix interfacial bonding strength. Due to efficient load transfer from the matrix to whisker, improved modulus and tensile strength of the composites are achieved, at the expense of ductility. Jayalakshmi et al. studied the mechanical behavior of Al2O3 short-fiber-reinforced composites with two different matrices, AM100 alloy and ZC63 alloy (Jayalakshmi et al. 2002, 2006). It was identified that the composites show no or little improvement in strength at room temperature along with a large reduction in ductility. Further, the properties at high temperatures showed that the tensile behavior was not dependent on temperature and volume fraction, and that they were influenced by the nature of the base matrix (i.e., brittle/ductile).

250

150

50

00 1 2 3 4

Strain (%)

Stre

ss (

MPa)

5

5%

10%

15%

6 7 8

100

200

alloy

300 70 8

6

4

2

0

60

50

400 5 10 15

250YS

Elastic ModulusElongation

Elo

ngat

ion

(%)UTS

200

150

500 5 10

Volume fraction of SiC particles (%) Volume fraction of SiC particles (%)

Ela

stic

Mod

ulus

(G

Pa)

Stre

ngth

(M

Pa)

15

100

a

b c

Fig. 2.23 (a) stress–strain curves (b) effect of volume fraction on strength and (c) effect of volume fraction on elongation in Mg–Zn–Ca/SiC particulate reinforced composites. Note that the variation in properties is influenced by the volume fraction, and that the ductility decreases drastically in the composites (adapted from Wang et al. 2012) (© 2012, Elsevier. Used with permission)


47

Considering nanoscale reinforcements, nano-Al2O3 has limited reactivity with molten Mg and therefore Mg–Al2O3 nanocomposites were the first of the Mg MMNCs to be prepared. Hassan and Gupta (2005) synthesized pure Mg-based nanocomposite with 1.1 Vf (%) of Al2O3 particulates (size 50 nm) using DMD tech-nique, followed by hot extrusion. The tensile properties showed an increase in yieldstrength by ~90 % and increase in ductility by ~45 %, when compared to the unre-inforced Mg. In AZ31B-1.5 % Vf Al2O3 nanocomposites produced by DMD, Nguyen and Gupta (2008) observed near-equiaxed grains with decreasing grain size, androundness of second phase with increasing nano-Al2O3 content. Tensile properties (assessed in terms of work of fracture) improved more than four times with enhanced ductility. CNTs have been used in Mg matrices to make composites. Goh et al. (2006) fabricated Mg nanocomposites containing 0.3, 1.3, 1.6, and 2 wt% CNTs using DMD process, that were hot extruded and characterized for their material behavior. TEM investigations (Goh et al. 2008) of dislocation structures in pure Mg and Mg 1.3CNT nanocomposites (Figs. 2.25 and 2.26) showed that only basal dis-locations were present in pure Mg (no prismatic or pyramidal slip planes), whereas in Mg 1.3 wt% CNT nanocomposite, basal and prismatic slip planes were observed. Additional slip systems observed in the CNT-reinforced nanocomposites showed the propensity of nano-reinforcement addition in activating the nonbasal slip systems, which strongly influence the mechanical behavior. Mg matrices have been incorpo-rated with SiC nanoparticles by various methods. Microstructural analyses of the as-cast specimens of pure Mg nanocomposites reinforced with 2 wt% nano- SiC

Fig. 2.24 Interfacial reaction resulting in the formation of MgO in AZ91/SiC-whisker reinforced composite with Al(PO3)3 binder (adapted from Zheng et al. 2001) (© 2001, Elsevier. Used with permission)


48

particles by ultrasonic cavitation-based dispersion method (Erman et al. 2012) showed about 60 % reduction in grain size of the composite. Mg-(2-4)Al-1SiC with 2 % nano-SiC by ultrasonic cavitation method (Cao et al. 2008) showed good dis-persion of nanoparticles (embedded inside the grains) with good interface bonding (no intermetallic phase formation). The yield strength of the Mg-2Al-1Si-2%SiC

Fig. 2.25 TEM investigation on pure Mg produced by the DMD method. The images show the dislocation structures on: (a) (2-1-10) foil plane and (b) (2-1-10) foil plane. Basal plane trace com-mon to (a) and (b) is indicated by white line (adapted from Goh et al. 2008) (© 2008, Elsevier. Used with permission)

Fig. 2.26 TEM images of dislocation structures in Mg-1.3 wt% CNT nanocomposite observed on: (a) (2-1-10) foil plane and (b) (2-1-10) foil plane. Basal plane trace common to (a) and (b) is indicated by black line. The composite was produced using the DMD method (adapted from Goh et al. 2008) (© 2008, Elsevier. Used with permission)


49

and Mg-4Al-1Si-2%SiC nanocomposites improved significantly, while retaining the ductility. AZ63 reinforced with nano-SiC particles (size 40 nm) using FSP (Sun et al. 2012) showed no interfacial reaction products at the matrix/particle interface. The ultimate tensile strength of the composite increased significantly (~312 MPa) when compared to ~160 MPa of the unreinforced alloy. Other representative studies that investigate nanocomposites properties are given in refs. Paramsothy et al. (2009, 2010), Jayaramanavar et al. (2009), Nguyen and Gupta (2010), Sankaranarayanan et al. (2011), Thakur et al. (2007a, b, c), Tun et al. (2012). The mechanical properties of selected Mg-micro/nanocomposites are listed in Table 2.6.

Table 2.6 Mechanical properties of several Mg-metal matrix composites/nanocomposites




% Ductility


Hu et al. (2010)AE44 (infiltration) – 98.05 197.69 10.26 43.6AE44-15 % Vf Al2O3

(infiltration)– 163.26 221.62 1.28 72.3

Lloyd (1994)AZ61 (cast) – 157 198 3 38AZ61-20 % Vf SiC

(cast)– 260 328 2.5 47.5

AZ91 (cast) – 168 311 21 49AZ91-9.4 % Vf SiC

(cast)– 191 276 2 47.5

AZ91-15.1 % Vf SiC (cast)

– 260 328 2.5 80

Contreras (2004)Mg-56 % Vf TiC

(pressureless infiltration)

195 – 200 – 130

Chen et al. (2010)AZ91D (vacuum–gas

pressure infiltration)

– 97 165 7.3 45

AZ91D-50 % Vf Mg2B2O5

– 262 265 0.98 54

Gu et al. (2004)Pure Mg (liquid pressure

infiltration)– – 98 16 –

Mg-8 % Vf (SiCp- Al2O3·SiO2-f)

– – 210 2.13 –

Mg-18 % Vf (SiCp- Al2O3·SiO2-f)

– – 267 0.64 –

(continued)


50





% Ductility


Gu et al. (2010)Pure Mg (PM + Ext) – 107 197 10.3 –Mg-10 wt% HA

(PM + Ext)– 117 171 6.7 –

Mg-20 wt% HA (PM + Ext)

– 105 146 4.3 –

Mg-30 wt% HA (PM + Ext)

– 71 92 2.6 –

Wang et al. (2006) HBMg–Li (As-cast) – 67.9 130 12 34.6Mg-5 wt%Al2Y

(As-cast)– 107.4 189 7 50.1

Zhang et al. (1997)ZK60A (liquid pressure

infiltration)– – 315 10.4 42.2

ZK60A-24 % Vf (SiCw-B4Cp)

– – 428 2.3 80.6

Zheng et al. (2004)AZ91 (As-cast) – 90 188 11.72 45AZ91-Al18B4O33w

(As-cast)– 266 352 0.97 67

AZ91-Al18B4O33w (250 °C/100 h)z

– 270 368 0.96 71

AZ91-Al18B4O33w (400 °C/50 h)

– 261 353 0.96 65

Zheng et al. 2001)AZ91 (squeeze cast) 54 87 189 8.14 45AZ91-20 % Vf SiCw

(Al(PO3)3binder)175 220 355 1.38 85

AZ91-20 % Vf SiCw (without binder)

174 202 314 1.29 77

Jayalakshmi et al. (2002) BHNAM100 (T6) squeeze

infiltration86 – 230 2 46

AM100-15%Vf Al2O3−SF (T6) (SI)

145 – 220 1.5 69


156 – 220 0.9 76


163 – 218 0.43 85

Jayalakshmi et al. (2006) BHNZC63 (T6) squeeze

infiltration (SI)60 195 270 8.2 47

(continued)


51





% Ductility


ZC63-15 %Vf Al2O3−SF (T6) (SI)

115 202 210 0.55 72

ZC63-20%Vf Al2O3−SF (T6) (SI)

138 202 212 0.73 80

ZC63-25%Vf Al2O3−SF (T6) (SI)

144 140 150 0.62 88

Nanocomposites

Hassan and Gupta (2005)Pure Mg (DMD + hot

ext.)40 97 173 7.4 –

Mg + 1.1 % Vf nAl2O3(DMD + hot ext.)

65 175 246 14 –

Nguyen and Gupta (2008)AZ31B (DMD + hot

ext.)63 201 270 5.6 –

AZ31B + 1.5 % Vf nAl2O3(DMD + hot ext.)

86 144 214 29.5 –

Jayaramanavar et al. (2009)ZK60A (DMD + hot

ext.) (T)97 139 ± 4 246 ± 4 20.2 ± 2 –

ZK60A + 1.5 Vf nAl2O3 (DMD + hot ext.) (T)

92 147 ± 8 252 ± 5 26.0 ± 1 –

Paramsothy et al. (2009)AZ31 (DMD + hot ext.)

(T)64 172 ± 15 263 ± 12 10.4 ± 3.9 –

AZ31 + 1.5 Vf nAl2O3 (DMD + hot ext.) (T)

83 204 ± 10 317 ± 5 22.2 ± 2.4 –

Sankaranarayanan et al. (2011)Mg (DMD + hot ext.) (T) 48 125 ± 9 169 ± 11 6.2 ± 0.7 –Mg 5.6 Ti-2.5 nAl2O3

(wt. %) (DMD) (T)74 175 ± 4 227 ± 10 3.3 ± 0.2 –

Mg 5.6 Ti-2.5 nAl2O3 (BM) (DMD) (T)

69 168 ± 8 214 ± 8 6.8 ± 0.8 –

Paramsothy et al. (2010)AZ31 (DMD + hot ext.) (T) 64 172 ± 15 263 ± 12 10.4 ± 3.9 –AZ31 + 1 Vf

CNT(DMD + hot ext.) (T)

95 190 ± 13 307 ± 10 18.0 ± 2.6 –

Nguyen and Gupta (2008)AZ31B (DMD + hot

ext.) (T)63 201 ± 7 270 ± 6 5.6 ± 1.4 –

(continued)


52





% Ductility


AZ31B + 1.5 Vf nAl2O3+ 3Ca (DMD) (T)

113 235 ± 7 285 ± 14 7.3 ± 0.2 –

Tun et al. (2012)Mg (microwave sint.)

(MWS + hot ext.) (T)

45 121 ± 5 179 ± 6 11 ± 1 –

Mg 1.0 Vf nZnO(MWS + hot ext.) (T)

62 125 ± 4 231 ± 13 17 ± 2 –

Thakur et al. (2007a)Mg (MWS + hot ext.)

(T)39 119 ± 7 170 ± 4 5.4 ± 1.5 –

Mg 1.0 nSiC(MWS + hot ext.) (T)

43 131 ± 12 182 ± 9 5 ± 0.4 –

Mg 0.5 nSiC + 0.5 nAl2O3 (MWS + hot ext.) (T)

46 155 ± 7 197 ± 1 4.6 ± 2 –

Thakur et al. (2007c)Mg 1.0 CNT(MWS hot

ext.) (T)43 113 ± 2.8 146 ± 6.5 1.9 ± 0.9 –

Mg + 0.3 CNT-0.7 nAl2O3(MWS) (T)

44 153 ± 2.1 196 ± 3.3 2.5 ± 0.8 –

Thakur et al. (2007b)Mg (MWS hot ext.) (T) 41 112 ± 7.7 155 ± 2.1 5.9 ± 1.2 –Mg 1.0 CNT (MWS hot

ext.) (T)43 117 ± 6.2 153 ± 2.8 1.5 ± 0.3 –

Mg 0.5 CNT + 0.5 nSiC(MWS) (T)

45 152 ± 1.2 188 ± 2.7 2.3 ± 0.6 –

Erman et al. (2012)Mg (USC) – 47 120 12.3 –Mg-1 wt% nSiC (USC) – 67 133 6.3 –

Cao et al. (2008)Mg (2,4)Al 1Si (USC) (T) – 62 145 8.8 –Mg (2,4)Al 1Si-2 nSiC

(USC) (T)– 82 178 9.5 –

Sun et al. (2012)AZ63 (FSP) (T) 80 – 150 6 –AZ63-nSiC (FSP) (T) 110 – 300 11 –


53

2.6 Limitations

In spite of the various advantages that the conventional ceramic reinforced LMMCs can provide, there exists a critical limitation that has restricted their application potential—which is their low ductility. As seen from Tables 2.5 and 2.6, the tensile elongation decreases rapidly with the addition of reinforcing particles. The reduc-tion in ductility arises due to the inherent brittle nature of the reinforcement and the poor interfacial characteristics between the reinforcement/matrix. Poor interface occurs because of the low wettability between the matrix and the particles/binder or reactions between binder/particle and the matrix. Owing to this drawback, particle decohesion, particle agglomeration/clusters, and nonbeneficial interfacial reactions occur. As mentioned earlier in Sect. 2.4.2, fracture behavior of composite is associ-ated with the interfacial strength or lack thereof.

In this context, there lies a great demand to explore better alternative reinforce-ments. The most promising solution is to incorporate bulk metallic glasses (BMG) as replacements for conventional ceramic reinforcements. BMGs are new class of metallic materials that exhibit superior mechanical properties, due to their structural properties (absence of long-range atomic order) and thermal behavior (i.e., exis-tence of glass-transition, Tg, and crystallization temperature, Tx). With these uniqueproperties, amorphous alloys exhibit high strength (~1–2 GPa) and large elastic strain limit (~1–2 %). When these novel materials are used as reinforcements in LMMCs, they can provide excellent interfacial properties and impart superior strength due to them being metallic in nature (metallic bonding), when compared to conventional ceramic reinforcements (covalent/ionic bonds). This will invoke syn-ergistic effect of the inherent structural/mechanical/thermal properties of the matrix and reinforcements, in order to achieve superior performance of LMMCs.

References

Abbasipour B, Niroumand B, Monir Vaghefi SM (2010) Compocasting of A356-CNT composite. Trans Nonferrous Met Soc China 20:1561–1566. doi:10.1016/S1003-6326(09)60339-3

Abramov O (1994) Ultrasound in liquid and solid metals. CRC Press, Boca RatonAhmed A, Neely AJ, Shankar K et al (2010) Synthesis, tensile testing, and microstructural charac-

terization of nanometric SiC particulate-reinforced Al 7075 matrix composites. Metall Mater Trans A 41:1582–1591. doi:10.1007/s11661-010-0201-y

Arsenault R (1983) Particulate mcirostructure of fiber and SiC in 6061 Al composites. Scr Metall 17:67–71

Balog M, Yu P, Qian M et al (2013) Nanoscaled Al–AlN composites consolidated by equal channelangular pressing (ECAP) of partially in situ nitrided Al powder. Mater Sci Eng A 562:190–195

Bathula S, Anandani RC, Dhar A, Srivastava K (2012) Microstructural features and mechanical properties of Al 5083/SiCp metal matrix nanocomposites produced by high energy ball milling and spark plasma sintering. Mater Sci Eng A 545:97–102. doi:10.1016/j.msea.2012.02.095

Brown L, Stobbs W (2006) The work-hardening of copper-silica v. equilibrium plastic relaxationby secondary dislocations. Philos Mag 34:351–372

References

http://dx.doi.org/10.1016/S1003-6326(09)60339-3

http://dx.doi.org/10.1007/s11661-010-0201-y

http://dx.doi.org/10.1016/j.msea.2012.02.095

54

Cao G, Konishi H, Li X (2008) Mechanical properties and microstructure of SiC-reinforced Mg-(2,4)Al-1Si nanocomposites fabricated by ultrasonic cavitation based solidification pro-cessing. Mater Sci Eng A 486:357–362. doi:10.1016/j.msea.2007.09.054

Chawla N, Chawla K (2006) Metal matrix composites. Springer, New YorkChawla K, Metzger M (1978) Advances in research on strength and fracture of materials, vol 3.

Pergamon Press, New York, p 1039Chen SH, Jin PP, Schumacher G, Wanderka N (2010) Microstructure and interface characteriza-

tion of a cast Mg2B2O5 whisker reinforced AZ91D magnesium alloy composite. Compos Sci Technol 70:123–129. doi:10.1016/j.compscitech.2009.09.015

Contreras A (2004) Mg/TiC composites manufactured by pressureless melt infiltration. Scr Mater 51:249–253. doi:10.1016/j.scriptamat.2004.04.007

Daoud A, Reif W (2002) Influence of Al2O3 particulate on the aging response of A356 Al-based composites. J Mater Process Technol 123:313–318. doi:10.1016/S0924-0136(02)00103-6

Das A, Chatterjee S (1981) Squeeze casting of an aluminium alloy containing small amounts ofSiC whiskers. Metall Mater Technol 137

Delanney F, Frozen L, Peryttere A (1987) Review—the wetting of solids by molten metals and its relation to the preparation of metal-matrix composites. J Mater Sci 22:1

Dieter GE (1988) Mechanical metallurgy. McGraw-Hill Higher Education, LondonDonthamsetty S, Damera NR, Jain PK (2009) Ultrasonic cavitation assisted fabrication and char-

acterization of A356 metal matrix nanocomposite reinforced with Sic, B4C, CNTs. Asian Int J Sci Technol Prod Manuf Eng 2:27–34

Ellis M (1996) Joining of aluminium based metal matrix composites. Int Mater Rev 41:41–58Erman A, Groza J, Li X et al (2012) Nanoparticle effects in cast Mg-1 wt% SiC nano-composites.

Mater Sci Eng A 558:39–43. doi:10.1016/j.msea.2012.07.048Esawi MK, Morsi K, Sayed A et al (2010) Effect of carbon nanotube (CNT) content on the

mechanical properties of CNT-reinforced aluminium composites. Compos Sci Technol 70:2237–2241. doi:10.1016/j.compscitech.2010.05.004

Evans R, Boyd J (2003) Near-interface microstructure in a SiC/Al composite. Scr Mater 49:59–63. doi:10.1016/S1359-6462(03)00180-5

Evans A, Marchi CS, Mortensen A (2003) Metal matrix composites in industry: an introduction and a survey, vol 1. Springer, New York

Ezatpour HR, Sajjadi SA, Sabzevar MH et al (2014) Investigation of microstructure and mechani-cal properties of Al6061-nanocomposite fabricated by stir casting. Mater Des 55:921–928

Fecht H (1995) Nanostructure formation by mechanical attrition. Nanostruct Mater 6:33–42Fridlyander J (1994) Metal matrix composites. Springer, New YorkFriedrich HE, Mordike BL (2006) Magnesium technology: metallurgy, design data, automotive

applications. Springer, BerlinGarcia-Hinojosa JA, Gonzalez CR, Juarez JAI, Surrapa MK (2004) Effect of grain refinement

treatment on the microstructure of cast Al–7Si–SiCp composites. Mater Sci Eng A 386:54–60. doi:10.1016/j.msea.2004.07.020

Ghomashchi M, Vikhrov A (2000) Squeeze casting: an overview. J Mater Process Technol 101:1–9Girot FA, Quenisset JM, Naslain R (1987) Discontinuously reinforced Al metal matrix composites.

Compos Sci Technol 30:155–184Goh CS, Wei J, Lee LC, Gupta M (2006) Simultaneous enhancement in strength and ductility by

reinforcing magnesium with carbon nanotubes. Mater Sci Eng A 423:153–156. doi:10.1016/j.msea.2005.10.071

Goh CS, Wei J, Lee LC, Gupta M (2008) Ductility improvement and fatigue studies in Mg-CNT nanocomposites. Compos Sci Technol 68:1432–1439. doi:10.1016/j.compscitech.2007.10.057

Gu J, Zhang X, Gu M (2004) Mechanical properties and damping capacity of (SiCp+Al2O3·SiO2f)/Mg hybrid metal matrix composite. J Alloys Compd 385:104–108. doi:10.1016/j.jallcom.2004.04.106

Gu X, Zhou W, Zheng Y et al (2010) Microstructure, mechanical property, bio-corrosion and cyto-toxicity evaluations of Mg/HA composites. Mater Sci Eng C 30:827–832. doi:10.1016/j.msec.2010.03.016

Gupta M, Eugene WWL (2007) Microwaves and metals. Wiley, Hoboken




http://dx.doi.org/10.1016/j.scriptamat.2004.04.007

http://dx.doi.org/10.1016/S0924-0136(02)00103-6



http://dx.doi.org/10.1016/S1359-6462(03)00180-5





http://dx.doi.org/10.1016/j.jallcom.2004.04.106

http://dx.doi.org/10.1016/j.msec.2010.03.016

http://dx.doi.org/10.1016/j.msec.2010.03.016

55

Gupta M, Sharon NML (2011) Magnesium, magnesium alloys, and magnesium composites. Wiley, Hoboken

Harris SJ (1988) Cast metal matrix composites. Mater Sci Technol 4:231Hassan SF, Gupta M (2005) Enhancing physical and mechanical properties of Mg using nanosized

Al2O3 particulates as reinforcement. Metall Mater Trans A 36:2253–2258Heinrich M, Gonasagren G (2012) Semi-solid processing of alloys and composites XII. Trans Tech

Publication, Durnten-ZurichHong C, Kim J (2006) Development of an advanced rheocasting process and its application. In:

Proceeding of the ninth international conference on the processing of semi-solid alloys and composite, Busan, Korea, pp 44–53

Hossein-Zadeh M, Mirzaee O, Saidi P (2014) Structural and mechanical characterization of Al-based composite reinforced with heat treated Al2O3 particles. Mater Des 54:245–250. doi:10.1016/j.matdes.2013.08.036

Hu L, Wang E (2000) Fabrication and mechanical properties of SiCw/ZK51A magnesium matrix composite by two-step squeeze casting. Mater Sci Eng A 278:267–271. doi:10.1016/S0921-5093(99)00608-5

Hu B, Peng L, Powell BR et al (2010) Interfacial and fracture behavior of short-fibers reinforced AE44 based magnesium matrix composites. J Alloys Compd 504:527–534. doi:10.1016/j.jallcom.2010.05.155

Jayalakshmi S, Kailas SV, Seshan S (2002) Tensile behaviour of squeeze cast AM100 magnesiumalloy and its Al2O3 fibre reinforced composites. Compos Part A Appl Sci Manuf 33:1135–1140

Jayalakshmi S, Kailas SV, Seshan S, Fleury E (2006) Properties of squeeze cast Mg-6Zn-3Cu alloyand its saffil alumina short fibre reinforced composites. J Mater Sci 41:3743–3752. doi:10.1007/s10853-005-4484-0

Jayaramanavar P, Paramsothy M, Balaji A, Gupta M (2009) Tailoring the tensile/compressive response of magnesium alloy ZK60A using Al2O3 nanoparticles. J Mater Sci 45:1170–1178. doi:10.1007/s10853-009-4059-6

Kamali Ardakani MR, Khorsand S, Amirkhanlou S, Javad Nayyeri M (2014) Application of compo-casting and cross accumulative roll bonding processes for manufacturing high-strength, highly uniform and ultra-fine structured Al/SiCp nanocomposite. Mater Sci Eng A 592:121–127

Kandemir S, Yalamanchili A, Atkinson H (2012) Production of aluminium matrix nanocomposite feedstock for thixoforming by an ultrasonic method. Key Eng Mater 504–506:339–344

Kaufmann H, Uggowitzer P (2001) Fundamentals of the new rheocasting process for magnesium alloys. Adv Eng Mater 3:963

Kennedy AR, Weston DP, Jones MI (2001) Reaction in Al–TiC metal matrix composites. Mater Sci Eng A 316:32–38. doi:10.1016/S0921-5093(01)01228-X

Kimura Y, Mishima Y, Umekawa S, Suzuki T (1984) Compatibility between carbon fibre and binary aluminium alloys. J Mater Sci 19:3107

Koli DK, Agnihotri G, Purohit R (2013) Properties and characterization of Al-Al2O3 composites pro-cessed by casting and powder metallurgy routes (review). Int J Latest Trends Eng Technol 2:486–496

Krishnan B, Surappa M, Rohatgi P (1981) The UPAL process: a direct method of preparing cast aluminium alloy-graphite particle composites. J Mater Sci 16:1209

Lee KB, Sim HS, Cho SY, Kwon H (2001) Reaction products of Al–Mg/B4C composite fabricated by pressureless infiltration technique. Mater Sci Eng A 302:227–234. doi:10.1016/S0921-5093(00)01831-1

Lee KB, Sim HS, Heo SW et al (2002) Tensile properties and microstructures of Al composite reinforced with BN particles. Compos Part A Appl Sci Manuf 33:709–715. doi:10.1016/S1359-835X(02)00011-8

Lee C, Huang J, Hsieh P (2006) Mg based nano-composites fabricated by friction stir processing. Scr Mater 54:1415–1420. doi:10.1016/j.scriptamat.2005.11.056

Li Q, Rottmair C, Singer R (2010) CNT reinforced light metal composites produced by melt stir-ring and by high pressure die casting. Compos Sci Technol 70:2242–2247

References

http://dx.doi.org/10.1016/j.matdes.2013.08.036

http://dx.doi.org/10.1016/S0921-5093(99)00608-5

http://dx.doi.org/10.1016/S0921-5093(99)00608-5



http://dx.doi.org/10.1007/s10853-005-4484-0

http://dx.doi.org/10.1007/s10853-005-4484-0

http://dx.doi.org/10.1007/s10853-009-4059-6

http://dx.doi.org/10.1016/S0921-5093(01)01228-X

http://dx.doi.org/10.1016/S0921-5093(00)01831-1

http://dx.doi.org/10.1016/S0921-5093(00)01831-1

http://dx.doi.org/10.1016/S1359-835X(02)00011-8

http://dx.doi.org/10.1016/S1359-835X(02)00011-8


56

Liao J, Tan M-J, Sridhar I (2010) Spark plasma sintered multi-wall carbon nanotube reinforced aluminum matrix composites. Mater Des 31:S96–S100. doi:10.1016/j.matdes.2009.10.022

Lii D-F, Huang J-L, Chang S-T (2002) The mechanical properties of AlN/Al composites manufac-tured by squeeze casting. J Eur Ceram Soc 22:253–261. doi:10.1016/S0955-2219(01)00255-2

Lloyd DJ (1994) Particle reinforced aluminium and magnesium matrix composites. Int Mater Rev 39:1–23. doi:10.1179/095066094790150982

Long A, Thornhill D, Armstrong A, Watson D (2012) Predicting die life from die temperature for high pressure dies casting aluminium alloy. Appl Therm Eng 44:100–107

Luo ZP, Song YG, Zhang SQ (2001) A TEM study of the microstructure of SiCp/Al composite prepared by pressureless infiltration method. Scr Mater 45:1183–1189

Manna A, Bhattacharayya B (2003) A study on machinability of Al/SiC-MMC. J Mater Process Technol 140:711–716

Maruyama B (1998) Progress and promise in aluminium metal matrix composites. AMPTIAC NewsLett 2(3)

Mazahery A, Shabani MO (2012) Mechanical properties of A356 matrix composites reinforced with nano-SiC particles. Strength Mater 44:686–692

Metcalfe AG (1974) Interfaces in metal matrix composites. In: Metcalfe AG (ed) Composite mate-rials. Academic, New York

Miller W, Humphreys FJ (1991) Strengthening mechanisms in particulate metal matrix compos-ites. Scr Metall Mater 25:33–38

Miracle D (2005) Metal matrix composites—from science to technological significance. Compos Sci Technol 65:2526–2540

Mortenson A, Flemmings JA, Cornie MC (1988) Solidification processing of metal matrix com-posites. J Appl Meteorol 40:12–19

Mula S, Padhi P, Panigrahi SC et al (2009) On structure and mechanical properties of ultrasoni-cally cast Al–2% Al2O3 nanocomposite. Mater Res Bull 44:1154–1160. doi:10.1016/j.materresbull.2008.09.040

Naher S, Brabazon D, Looney L (2005) Development and assessment of a new quick quench stircaster design for the production of metal matrix composites. J Mater Process Technol 166:430–439

Nardone V, Prewo K (1986) On the strength of discontinuous silicon carbide reinforced aluminium composites. Scr Metall 20:43–48

Nguyen QB, Gupta M (2008) Increasing significantly the failure strain and work of fracture of solidification processed AZ31B using nano-Al2O3 particulates. J Alloys Compd 459:244–250. doi:10.1016/j.jallcom.2007.05.038

Nguyen QB, Gupta M (2010) Enhancing mechanical response of AZ31B using Cu+nano-Al2O3 addition. Mater Sci Eng A 527:1411–1416. doi:10.1016/j.msea.2009.11.002

Nie KB, Wang XJ, Wu K et al (2011) Processing, microstructure and mechanical properties of magnesium matrix nanocomposites fabricated by semisolid stirring assisted ultrasonic vibra-tion. J Alloys Compd 509:8664–8669. doi:10.1016/j.jallcom.2011.06.091

Noguchi T, Asano K, Hiratsuka S, Miyahara H (2008) Trends of composite casting technology and joining technology for castings in Japan. Int J Cast Met Res 21:219–225

Paramsothy M, Hassan SF, Srikanth N, Gupta M (2009) Enhancing tensile/compressive response of magnesium alloy AZ31 by integrating with Al2O3 nanoparticles. Mater Sci Eng A 527:162–168. doi:10.1016/j.msea.2009.07.054

Paramsothy M, Hassan SF, Srikanth N, Gupta M (2010) Simultaneous enhancement of tensile/compressive strength and ductility of magnesium alloy AZ31 using carbon nanotubes. J Nanosci Nanotechnol 10:956–964. doi:10.1166/jnn.2010.1809

Reed-Hill R (1973) Role of deformation twinning in determining the mechanical properties of metals: In: The inhomogeneity of plastic deformation. ASM International, Materials Park

Rohatgi PK, Asthana R, Das S (1986) Solidification, structures and properties of cast metal- ceramic particle composites. Int Met Rev 31:115

Saheb N, Iqbal Z, Khalil A et al (2012) Spark plasma sintering of metals and metal matrix nano-composites: a review. J Nanomater 2012:1–13. doi:10.1155/2012/983470



http://dx.doi.org/10.1016/S0955-2219(01)00255-2

http://dx.doi.org/10.1179/095066094790150982

http://dx.doi.org/10.1016/j.materresbull.2008.09.040

http://dx.doi.org/10.1016/j.materresbull.2008.09.040





http://dx.doi.org/10.1166/jnn.2010.1809

http://dx.doi.org/10.1155/2012/983470

57

Sairam K, Sonber J, Murthy T et al (2013) Influence of spark plasma sintering parameters on den-sification and mechanical properties of boron carbide. Int J Refract Met Hard Mater 42:185–192

Saito Y, Utsunomiya H, Tsuji N, Sakai T (1999) Novel ultra-high straining process for bulk materi-als—development of the accumulative roll-bonding (ARB) process. Acta Mater 47:579–583

Sajjadi S, Torabi Parizi M, Ezatpour H, Sedghi A (2012) Fabrication of A356 composite reinforced with micro and nano Al2O3 particles by a developed compocasting method and study of its properties. J Alloys Compd 511:226–231

Sankaranarayanan S, Jayalakshmi S, Gupta M (2011) Effect of ball milling the hybrid reinforce-ments on the microstructure and mechanical properties of Mg–(Ti+n-Al2O3) composites. J Alloys Compd 509:7229–7237. doi:10.1016/j.jallcom.2011.04.083

Schultz BF, Ferguson JB, Rohatgi PK (2011) Microstructure and hardness of Al2O3 nanoparticle reinforced Al–Mg composites fabricated by reactive wetting and stir mixing. Mater Sci Eng A 530:87–97. doi:10.1016/j.msea.2011.09.042

Segal V (1999) Equal channel angular extrusion: from macromechanics to structure formation.Mater Sci Eng A 271:322–333

Shorowordi KM, Laoui T, Haseeb ASMA et al (2003) Microstructure and interface characteristics of B4C, SiC and Al2O3 reinforced Al matrix composites: a comparative study. J Mater Process Technol 142:738–743. doi:10.1016/S0924-0136(03)00815-X

Sklenicka V, Dvorak J, Svoboda M et al (2013) Equal-channel angular pressing and creep in ultra-fine-grained aluminium and its alloys. In: Ahmad Z (ed) Aluminium alloys—new trends in fabrication and applications. InTech, Rijeka

Stacey MH (1988) Production and characterization of fibres for MMCs. Mater Sci Technol 4:227–230

Su H, Gao W, Feng Z, Lu Z (2012) Processing, microstructure and tensile properties of nano-sized Al2O3 particle reinforced aluminum matrix composites. Mater Des 36:590–596. doi:10.1016/j.matdes.2011.11.064

Sun K, Shi QY, Sun YJ, Chen GQ (2012) Microstructure and mechanical property of nano-SiCp reinforced high strength Mg bulk composites produced by friction stir processing. Mater Sci Eng A 547:32–37. doi:10.1016/j.msea.2012.03.071

Surappa M (2003) Aluminium matrix composites: challenges and opportunities. Sadhana 28:319Suryanarayana C (2001) Mechanical alloying and milling. Prog Mater Sci 46:1–184Suryanarayana C (2011) Synthesis of nanocomposites by mechanical alloying. J Alloys Compd

509:S229–S234. doi:10.1016/j.jallcom.2010.09.063Suryanarayana C, Al-Aqeeli N (2013) Mechanically alloyed nanocomposites. Prog Mater Sci

58:383–502. doi:10.1016/j.pmatsci.2012.10.001Suslick KS, Didenko Y, Fang MM, Hyeon T, Kolbeck KJ, McNamara WB et al (1999) Acoustic

cavitation and its chemical consequences. Philos Trans R Soc Lond A 357:335–353Tee KL, Lu L, Lai MO (2000) Synthesis of in situ Al±TiB2 composites using stir cast route.

Compos Struct 47:589–593Tekmen C, Ozdemir I, Cocen U, Onel K (2003) The mechanical response of Al–Si–Mg/SiCp

composite: influence of porosity. Mater Sci Eng A 360:365–371. doi:10.1016/S0921-5093(03)00461-1

Thakur SK, Balasubramanian K, Gupta M (2007a) Microwave synthesis and characterization of magnesium based composites containing nanosized SiC and hybrid (SiC+Al[sub 2]O[sub 3]) reinforcements. J Eng Mater Technol 129:194. doi:10.1115/1.2400279

Thakur SK, Kwee GT, Gupta M (2007b) Development and characterization of magnesium com-posites containing nano-sized silicon carbide and carbon nanotubes as hybrid reinforcements. J Mater Sci 42:10040–10046. doi:10.1007/s10853-007-2004-0

Thakur SK, Srivatsan TS, Gupta M (2007c) Synthesis and mechanical behavior of carbon nano-tube–magnesium composites hybridized with nanoparticles of alumina. Mater Sci Eng A 466:32–37. doi:10.1016/j.msea.2007.02.122

References



http://dx.doi.org/10.1016/S0924-0136(03)00815-X





http://dx.doi.org/10.1016/j.pmatsci.2012.10.001

http://dx.doi.org/10.1016/S0921-5093(03)00461-1

http://dx.doi.org/10.1016/S0921-5093(03)00461-1

http://dx.doi.org/10.1115/1.2400279

http://dx.doi.org/10.1007/s10853-007-2004-0


58

Tjong S (2007) Novel nanoparticle-reinforced metal matrix composites with enhanced mechanical properties. Adv Eng Mater 9:639–652

Tun KS, Gupta M (2009) Development of magnesium/(yttria+nickel) hybrid nanocomposites using hybrid microwave sintering: microstructure and tensile properties. J Alloys Compd 487:76–82. doi:10.1016/j.jallcom.2009.07.117

Tun KS, Jayaramanavar P, Nguyen QB et al (2012) Investigation into tensile and compressive responses of Mg-ZnO composites. Mater Sci Technol 28:582–588. doi:10.1179/1743284711Y.0000000108

Uozumi H, Kobayashi K, Nakanishi K et al (2008) Fabrication process of carbon nanotube/light metal matrix composites by squeeze casting. Mater Sci Eng A 495:282–287

Vicens J, Chedru M, Chermant JL (2002) New Al–AlN composites fabricated by squeeze casting:interfacial phenomena. Compos Part A Appl Sci Manuf 33:1421–1423

Wang Y, Huang J (2003) Texture analysis in hexagonal materials. Mater Chem Phys 81:11–26. doi:10.1016/S0254-0584(03)00168-8

Wang SJ, Wu GQ, Li RH et al (2006) Microstructures and mechanical properties of 5 wt.% Al2Yp/Mg–Li composite. Mater Lett 60:1863–1865. doi:10.1016/j.matlet.2005.12.038

Wang L, Turnley P, Savage G (2011) Gas content in high pressure die castings. J Mater Process Technol 211:1510–1515

Wang XJ, Nie KB, Sa XJ et al (2012) Microstructure and mechanical properties of SiCp/MgZnCa composites fabricated by stir casting. Mater Sci Eng A 534:60–67. doi:10.1016/j.msea.2011.11.040

Witkin D, Lavernia E (2006) Synthesis and mechanical behavior of nanostructured materials via cryomilling. Prog Mater Sci 51:1–60

Xiong B, Xu Z, Yan Q et al (2011) Effects of SiC volume fraction and aluminum particulate size on interfacial reactions in SiC nanoparticulate reinforced aluminum matrix composites. J Alloys Compd 509:1187–1191. doi:10.1016/j.jallcom.2010.09.171

Yang Y, Lan J, Li X (2004) Study on bulk aluminum matrix nano-composite fabricated by ultra-sonic dispersion of nano-sized SiC particles in molten aluminum alloy. Mater Sci Eng A 380:378–383. doi:10.1016/j.msea.2004.03.073

Yasunori M, Hiroto T, Atsushi S (1996) Method and apparatus for shaping semisolid metals. EP0745694B1

Ye H, Liu X (2004) Review of recent studies in magnesium. J Mater Sci 9:6153–6171Ye J, He J, Schoenung J (2005) Cryomilling for the fabrication of a particulate B4C reinforced Al

nanocomposite: part I. Effects of process conditions on structure. Metall Mater Trans A 37(10):3099–3109

Zhang Z, Chen D (2006) Consideration of Orowan strengthening effect in particulate-reinforced metal matrix nanocomposites: a model for predicting their yield strength. Scr Mater 54:1321–1326. doi:10.1016/j.scriptamat.2005.12.017

Zhang X, Zhang D, We R et al (1997) (SiCw+B4Cp)/ZK60A Mg alloy matrix composite. Scr Mater 37:1631–1635

Zhang Q, Chen G, Wu G et al (2003) Property characteristics of a AlNp/Al composite fabricated by squeeze casting technology. Mater Lett 57:1453–1458. doi:10.1016/S0167-577X(02)01006-6

Zhang Q, Ma X, Wu G (2013) Interfacial microstructure of SiCp/Al composite produced by the pressureless infiltration technique. Ceram Int 39:4893–4897. doi:10.1016/j.ceramint.2012.11.082

Zheng M, Wu K, Yao C (2001) Characterization of interfacial reaction in squeeze cast SiCw/Mgcomposite. Mater Lett 47:118–124

Zheng MY, Wu K, Liang M et al (2004) The effect of thermal exposure on the interface and mechanical properties of Al18B4O33w/AZ91 magnesium matrix composite. Mater Sci Eng A 372:66–74. doi:10.1016/j.msea.2003.09.085



http://dx.doi.org/10.1179/1743284711Y.0000000108

http://dx.doi.org/10.1179/1743284711Y.0000000108

http://dx.doi.org/10.1016/S0254-0584(03)00168-8

http://dx.doi.org/10.1016/j.matlet.2005.12.038





http://dx.doi.org/10.1016/S0167-577X(02)01006-6

http://dx.doi.org/10.1016/j.ceramint.2012.11.082

http://dx.doi.org/10.1016/j.ceramint.2012.11.082



Chapter 3Amorphous Alloys/Bulk Metallic Glasses (BMG)

Abstract The background of metallic amorphous materials/bulk metallic glasses (BMGs) is briefly introduced. The various methods of processing that give rise to amorphous phase formation are explained. The properties of amorphous/BMG systems based on Zr, Ni, Ti, etc., are presented in terms of their structure, thermal, and mechanical properties. The major limitations that restrict the use of these novel materials are mentioned.

Keywords Amorphous materials • Bulk metallic glasses (BMGs) • Glass forming ability • Processing methods • Structural and thermal properties • Mechanical prop-erties • Limitations

3.1 Background

The advancement in technology has resulted in the research and development of advanced materials. For example, while the process of alloying has improved the properties of a pure metal, making metal matrix composites out of the metals/alloys has provided a synergistic effect of utilizing the advantages of both hard/stiff ceramic and relatively soft/ductile metal matrix, giving rise to improvements in properties such as stiffness, mechanical, structural, high temperature and wear behavior. Intermetallics and functionally graded materials are other notable exam-ples. The development of new materials has also given rise to novel processing/synthesizing methods. In this context, the development of metallic amorphous alloys/bulk metallic glasses (BMGs) is significant.

Metallic materials are crystalline in nature, i.e., metallic materials exhibit an ordered structure of atoms wherein there is a definite periodicity in their structure in three dimensions (presence of long-range atomic order). This ensures the possibility of prediction of the position of each atom in the crystal lattice, their nearest- neighbor distance, and the number of such nearest-neighbors. On the other hand, amorphous materials/BMGs are a relatively new class of metallic materials that show no atomic periodicity (noncrystalline solids). Similar to conventional oxide glasses, they are said to possess a random-network structure (absence of long-range atomic order),

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and hence different nearest-neighbors, resulting in a complex atomic structure (Fig. 3.1). Rapid solidification processes (RSP), high-energy ball-milling, severe plastic deformation, etc., are some of the processes that are used to produce amor-phous/BMG (mentioned in detail in the next section). Usually noncrystalline metal-lic materials prepared by continuous cooling from the liquid state are referred to as glass, whereas those prepared by other methods such as mechanical alloying are called amorphous alloys. Also, sometimes ribbons are referred to as amorphous, whereas those which can be prepared to at least ~1 mm thickness are called BMG (Suryanarayana and Inoue 2010).

3.1.1 Formation and Characteristics of Amorphous Alloys/ BMG

Figure 3.2 shows the formation of amorphous/glassy material from the liquid state. Crystalline metals follow a well-defined path that is thermodynamically stable, have low energy and at equilibrium conditions. Noncrystalline solids such as metallic glass follow nonequilibrium path favored by fast cooling. Apart from not having crystalline structure, amorphous alloys/BMG are distinctly different from conventional metals in various aspects. Some are mentioned here. (1) A typical X-ray diffractogram (XRD) of a conventional crystalline metal and that of an amorphous/BMG material is given in Fig. 3.3. While crystalline peaks are observed in the metal mixture, the amorphous alloy/BMG exhibit a diffused halo pattern. (2) Unlike crystalline metals, amorphous alloys/metallic glasses are metastable in nature. This is due to the nonequilibrium processes involved (such as RSP/high-energy ball-milling/mechanical alloying), wherein the initial crystalline material in equilibrium is transformed to an energized state (Fig. 3.4a), and given sufficient driving force (time/temperature, etc.,) it will

Fig. 3.1 Schematic showing the atomic structure in crystalline and amorphous alloys

3 Amorphous Alloys/Bulk Metallic Glasses (BMG)

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Fig. 3.2 Schematic showing the formation of amorphous alloys/metallic glass

Fig. 3.3 X-ray diffractogram showing the crystalline peaks of Ni, Nb in the as-mixed Ni60Nb40 (at.%) crystalline powder and its gradual evolution into amorphous state as a function of bal- milling time. Ni60Nb40 amorphous alloy powder is obtained after 87 h of high-energy ball-milling. Note the presence of prominent crystalline peaks in the as-mixed state and the presence of diffused halo in the amorphous state. The diffused halo is typical of an amorphous structure

3.1 Background

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transform back to its stable equilibrium crystalline phase. Figure 3.4b shows the schematic of free-energy relationship in crystal-to- amorphous/glass transformation (Suryanarayana 2001). (3) Unlike conventional crystalline metals, the thermal char-acteristic of amorphous/glassy metals (Fig. 3.5a, as obtained from differential scan-ning calorimeter, DSC) consists of glass transition temperature, Tg, below which it behaves like glass (Johnson 1999). At crystallization temperatures (Tx-onset and Tx-peak), amorphous/glassy state-to-nanocrystalline/crystalline state transformation occurs, and at Tm melting occurs. In the region of Tx − Tg (ΔTx—supercooled liquid region), most metallic glasses exhibit Newtonian viscous flow behavior/superplasticity (Kato et al. 1998). (4) The volume of most of the materials decreases with decreasing tem-perature from the liquid state (Fig. 3.5b). In contrast with that of a normal metal wherein the volume decreases suddenly at the freezing temperature, in a metallic glass (i.e., glass forming liquid compositions), there is a gradual decrease in volume even when the liquid is undercooled significantly (Suryanarayana and Inoue 2010). (5) As seen from Fig. 3.5c, for a conventional metal when cooled from the molten state, the viscosity increases slowly with decrease in temperature, and increases sud-denly when it reaches the freezing point (i.e., beginning of crystallization). However, for a metallic glass, the gradual increase in viscosity with decrease in temperature occurs even below the freezing point (Spaepen 2006). While the rate of increase is rapid with further decrease in temperature (supercooled liquid), at Tg, the viscosity if very high. Such high viscosity (slow atomic motion) prevents the formation of long-range atomic order, and the material behaves no more as a liquid and is completely a solid (Samwer et al. 1994). (6) Conventional metals exhibit defects (vacancies/ dislocations grain/grain boundaries, etc.) which govern their deformation processes and hence their strengthening mechanisms. On the contrary, grains/grain boundaries, etc. are absent in metallic glasses. Hence the deformation at low temperature/high stress is usually inhomogeneous (failure by shear band formation) and at high tempera-ture/low stress, the deformation is homogenous (Fig. 3.6) (Spaepen 2006). (7) Unlike conventional metals, amorphous alloys/BMGs consists of free volume (space between atoms formed due to the frozen-in state) (as shown in the schematic in Fig. 3.7),

Ene

rgy

Stablestate

Amorphous PhaseFormation

Stable

Unstable

Metastable

EnergyGain

EnergyBarrier

a b

Meta-stablestate

Solid, Liquid, or Vapor

Rapid SolidificationMechanical AlloyingVapor deposition

Metastable Phase

Equilibrium Crystalline State

T, P, E G1

G2

G

G0

Fig. 3.4 (a) Schematic showing energy states in the processing amorphous alloys/metallic glass and (b) shows the schematic of free-energy relationship in crystal-to-amorphous/glass transforma-tion (adapted from Suryanarayana 2001) (© 2001, Elsevier. Used with permission)


Fig. 3.5 (a) The thermal characteristic of amorphous/glassy metals as obtained from differential scanning calorimeter (DSC) consisting of glass transition temperature, Tg and crystallization tem-perature Tx. The DSC pattern shown here is that of Vitreloy 1 (b) The variation of volume with temperature for crystalline and amorphous materials. (c) Variation of viscosity with temperature between crystalline and amorphous materials (adapted from Spaepen 2006) (© 2006, Elsevier. Used with permission)

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Temperature

log

T/µ

Tg Tx Tm

“Viscous”

“Solid”

Homogeneous flow

Inhomogeneous flow

Unaccessible

“Fluid”

−20

−15

−10

0

−5

=10−1 s−1γ

=10−4 s−1γ

=10−8 s−1γ

Fig. 3.6 Deformation behavior of amorphous alloys/metallic glass showing inhomogeneous and homogenous characteristics. These are strongly dependent on temperatures and strain rates (adapted from Spaepen 2006) (© 2006, Elsevier. Used with permission)

Fig. 3.7 Schematic showing the presence of free volume in the amorphous structure


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which according to one of the theories of metallic glass, is believed to aid in the deformation process, especially at high temperatures (De Hey et al. 1997; Spaepen 2006; Bletry et al. 2006).

As mentioned earlier, for the formation of glass, the melt needs to be under-cooled below Tg. However, it should be noted that all alloy compositions cannot form BMG. Over years, several important aspects necessary for glass formation (glass forming ability, GFA) was put forward to establish glass forming criteria. It is known that, amorphous/glassy state can be achieved if the formation/nucleation of crystals is suppressed (by undercooling). Such large undercooling can be obtained by rapid solidification at faster rate, typically above a critical cooling rate (Rc) and depends on the alloy system and composition. The dependence of critical cooling rate on the alloy system/composition (Suryanarayana and Inoue 2010) can be explained by the TTT diagram (time–temperature–transformation) shown in Fig. 3.8 (Basu and Ranganathan 2003). In the figure, if the addition of alloying elements is more, the C-curve is shifted to the right and that the transformation can be slowed/delayed/longer times (i.e., lower critical cooling rate). On the other hand, if the number of alloying elements is less, the curve is shifted to the left (shorter times) (Suryanarayana and Inoue 2010). Hence in binary/ternary systems using processes like melt spinning (lower thickness—usually in the form of ribbons), the amorphous state can be obtained for Rc ~ 105–106 K s−1 whereas for alloys with multiple con-stituents, glass formation can be achieved at critical cooling rate as low as 10−3 K s−1 (e.g., suction casting). Figure 3.9 shows the variation of critical cooling rate/dimen-sion/thickness of amorphous/BMG with reduced glass transition temperature (Basu and Ranganathan 2003). Similarly, the ratio of the glass transition temperature to

Liquid

Time (s)

Tem

pera

ture

(K

)

010–2 102 10410–410–6

Slowcooling

RSP

Glass

Crystallization

Tg

Tl

Fig. 3.8 Time–temperature–transformation (TTT) diagram showing the dependence of critical cooling rate on the amorphous alloy system/composition (adapted from Basu and Ranganathan 2003) (© 2003, Springer. Used with permission)

3.1 Background

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the liquidus temperature (called the reduce glass transition temperature) should be high as it would increase the viscosity of the alloys, thereby promoting glass forma-tion at slower cooling rates.

A multitude of compositions (binary, ternary, quaternary, and multicomponent systems) have been formed, the details of which are available in literature. According to Inoue (1998), the probability of glass formation (GFA) increases if the following criteria are met: (1) the composition is composed of binary or higher-order alloys (multicomponent system), (2) the composition has a deep eutectic, i.e., the eutectic temperature of the alloy (binary and ternary phase diagram of the constituents in case of a multicomponent system) is significantly lower than the melting points of each component in the system, (3) there is large difference in the radii of the indi-vidual constituents (i.e., radius ratio, Δr/r > 0.12), and (4) the components have negative heat of mixing. These factors together (“confusion principle”) results in dense atomic packing, high viscosity, difficulty for atomic rearrangement thereby retarding/suppressing the nucleation and growth of the crystalline phase. However, while these conditions have been proposed for amorphous phase/glass formation, there are also some exceptions wherein glass formation occurred without satisfying these conditions.

Fig. 3.9 Variation of critical cooling rate and dimension/thickness of amorphous alloy/BMG as a function of temperature (adapted from Basu and Ranganathan 2003) (© 2003, Springer. Used with permission)


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3.2 Preparation Methods

As seen in Fig. 3.9, metallic amorphous materials in the form of ribbons, wires, powders, etc., have been very successfully produced by rapid solidification methods at a cooling rate of 106 K s−1. These were in tens of microns in thickness, and for glassy materials to be produced with at least >1 mm, other processes are required to be employed and at a lower cooling rate. In this section, the important processing methods used to produce amorphous alloys/metallic glasses are presented.

3.2.1 Rapid Solidification Processes

Pol Duwez and colleagues first produced rapid solidification of Au-25%Si metallic melts at the California Institute of Technology that led to the discovery of metallic glass (Duwez 1967). The process involves the rapid solidification of a molten metal/alloy at cooling rates as high as 106 K s−1 (rapid heat extraction rate). Such high cooling rates can be achieved by using: (1) droplet method, wherein atomized mol-ten metal droplet is splat cooled/air quenched/inert gas quenched on thermally con-ducting substrate (such as copper), so as to achieve high heat extraction rates (2) jet method, in which a continuous molten metal stream is allowed to solidify on a mov-ing chill surface (copper wheel), so as to produce rapidly solidified ribbons/ filaments/sheets and (3) surface melting, so as to rapidly melt the surface (e.g., laser melting) when compared to the bulk of the metal, which acts as a rapid heat sink. To achieve high solidification rates, formation of a thin layer of the molten metal, inti-mate thermal contact with the substrate, and rapid heat extraction by the substrate from the molten liquid metals are required. For rapid solidification to occur, the rate R, is inversely proportional to the square of the thickness of the solidified molten metal layer. The solidification rate for a molten metal layer with thickness x and heat transfer coefficient of ∞ is given by,

R

A

x=

2

(3.1)

where x is the distance from the exit/orifice of molten metal/substrate interface and A is a constant and is a function of materials properties and temperature. Under ideal cooling conditions wherein the heat transfer coefficient is ∞, the value of A is 8.1 × 10−3 m2 K s−1. Under nonideal conditions, the value is less than the above- mentioned value.

One of the most common and widely used methods for producing amorphous/glassy materials by rapid solidification is the melt spinning technique. The schematic of the method is shown in Fig. 3.10 (Sowjanya and Kishen Kumar Reddy 2014).


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In the melt spinning process, small quantity of the alloy composition is melted in a crucible (usually quartz tube). The molten metal is ejected under pressurization through an orifice and the exiting molten metal is allowed to solidify on a chilled substrate (such as a fast rotating copper wheel). Upon impingement/in contact with the moving substrate, rapid heat extraction occurs and solidification of the metal occurs. Various parameters such as the alloy composition, melt temperature, exit pressure, nozzle dimension, vacuum conditions, rpm of the rotating wheel influence the size, shape, thickness, and quality of the ribbons. Usually, the outer surface of the wheel is polished to remove surface roughness as the wheel side of the ribbon imi-tates that of the wheel. Also, the faster the rpm of the wheel, the thinner is the ribbon. The presence of vacuum/inert conditions during melt spinning is also very essential, as ribbons when produced in air atmosphere result in porosity (air side) thus resulting in ribbons of low quality. In most cases, the difference in heat extraction between the wheel side and air side result in inhomogeneous surface characteristics/composition, affecting their surface/mechanical properties. It should be noted that although various composition of melt spun amorphous alloy ribbons/metallic glasses have been pro-duced successfully, the high solidification rates involved (105–106 K s−1) pose a constraint in the dimension (thickness) of the specimens (10–60 μm in the form of ribbons,

Molten alloy

Induction coil

Crucible

Slit nozzle

Amorphousribbon

Polished Cuwheel

Fig. 3.10 Schematic of the melt spinning process (adapted from Sowjanya and Kishen Kumar Reddy 2014) (© 2014, Elsevier. Used with permission)


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wires, pulverized and consolidated powders), and hence on the applicative prospects of these materials. Some of the few but important applications of amorphous/glassy ribbons are in transformer cores (soft and hard magnetic properties) (Chrobak et al. 2006; Kane et al. 2012), hydrogen membrane separators, e.g., during coal gasifica-tion (Hara et al. 2002; Yamaura et al. 2005), bipolar plates, and catalytic membranes in fuel cells, etc. (Fleury et al. 2006; Jayalakshmi et al. 2012). In order to overcome this constraint, methods to increase the thickness/dimension of the materials with slower solidification rates are developed, wherein the solidification rates involved are 102–103 K s−1 (see Fig. 3.9) and the materials are referred to as BMGs.

3.2.2 Methods for Producing BMG

Given that rapid solidification techniques cannot be utilized to produce metallic glasses of relatively larger dimensions (solidification rate is inversely proportional to thickness), alternate methods need to be adopted to produce BMGs. Some of them are briefed here.

3.2.2.1 Melting and Casting Methods

Arc Melting/Induction Melting

Arc melting is usually used in alloy compositions that require low critical cooling rate to form glassy structure (Suryanarayana and Inoue 2010). Melting is carried out using a DC power to strike an arc between the electrode and the copper mold/hearth (water-cooled), wherein the alloy to be melted is placed. The intense heat generated by the arc causes melting while the copper mold/hearth acts as a heat sink thus extracting the heat from the melt resulting in solidification of the alloy. It should be noted that in this process, the metal at the bottom is always in contact with the cop-per mold, causing incomplete melting. This gives rise to heterogeneous nucleation and the resulting arc-melted alloy is not fully glassy (low volume fraction crystal-line phases). Due to this reason, the alloys produced by the arc-melting process are re-melted several times and used as master alloys. Similarly, induction melting is also widely used to produce BMG master alloy ingots. These are usually followed by processes such as those discussed below.

High Pressure Die-Casting

The molten alloy is injected using a plunger into a mold under high pressure (hydraulic pressure) to ensure immediate solidification of the alloy as soon it comes into contact with the mold, resulting in high solidification rates and short process time (in milliseconds).


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Squeeze Casting

Squeeze casting is another method to produce metallic glasses. It involves solidifi-cation of molten metal under pressure, so that it ensures higher solidification rates, efficient heat transfer between the mold and the metal, and elimination of porosities due to pressure application. More details regarding the squeeze casting process have been explained earlier in Chap. 2.

Water Quenching

This method involved melting of the glassy composition (using arc melting/induc-tion melting) in a quartz tube, and directly quenching the quartz tube in water, resulting in immediate solidification of the metal inside. Turnbull and co-workers prepared Pd40Ni40P20 of 5–10 mm size (Drehman et al. 1982) and Inoue et al. (1989) produced La55Al25Ni20 metallic glass of 1.2 mm diameter using this method. As the cooling rates attained can be ~102 K s−1, it is very suitable for alloy com-positions which have low critical cooling rate (i.e., high GFA). It should be noted that the wall thickness of the tube is usually ~1 mm and a length of few centimeters.

Suction Casting

In this process, the molten metal is sucked into a mold/die cavity (usually copper) using a pressure differential between the melting and casting chambers (Fig. 3.11). The process is usually combination of arc melting (melting chamber) followed by suction casting (casting chamber) (Fig. 3.11). The copper mold is connected to a vacuum source that creates a difference in pressure between the chambers, and enabling the molten metal to be sucked into the copper die/mold when the piston between the two chambers is removed. The process is suitable for both high and low GFA compositions as the suction process can force molten metal to be cast into mold cavities with even small dimensions.

3.2.3 Mechanical Alloying

Apart from the solidification of melt from the liquid state, amorphous metals/metal-lic glasses are widely processed using powder metallurgy techniques, viz., mechani-cal alloying. Mechanical alloying is a solid state processing technique; where in metal powders of required composition are initially blended and mixed in a steel vial in a ball-mill. The mixed elemental powder particles are then subjected to extensive grinding in the vial using a grinding medium (usually stainless steel/tung-sten carbide/alumina balls) and are agitated at high speeds for a desired period of


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time (high-energy ball-milling), until the required composition/structure/reaction is established. The process involves repeated cold welding, flattening, fracturing, and re-welding of powder particles. Figures 3.12 and 3.13 show a schematic of the pro-cess for crystalline-to-amorphous transformation and the mechanism of amorphiza-tion (Suryanarayana 2001). Figure 3.14 shows a representative picture of the vial, balls, and the ball-milling machine.

Mechanical alloying is a complex process and involves optimization of a number of variables (Suryanarayana 2001). Parameters such as the type of mill (attrition/planetary/shaker), milling container/vial, milling speed, milling time, type, size, and distribution of the grinding medium, ball-to-powder weight ratio, extent of fill-ing the container, milling atmosphere, process control agent and temperature of milling, etc. should be carefully selected in order to achieve the desired phase/microstructure.

Other liquid/vapor state methods include cap-cast technique, centrifugal casting, electromagnetic vibration, vapor-phase deposition, etc. Similarly, in solid state techniques, hydrogen-induced amorphization, pressure-induced amorphization,

Vacuum GaugeGas inlet

Vacuum chamber

Vacuumchamber

Rotary Pump

MeltingChamber

CastingChamber

Arc Melted and Suction CastBMG Ingots

Using Cu-molds

Fig. 3.11 Schematic of the arc-melting and suction-casting processes. The prepared BMG ingots are also shown in the figure

Fig. 3.12 Shows the schematic of the process from crystalline-to-amorphous transformation by high-energy ball-milling process (adapted from Suryanarayana 2001) (© 2001, Elsevier. Used with permission)


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irradiation processes, severe plastic deformation processes, accumulative roll bond-ing, etc. (some of these processes are mentioned with illustrations in Chap. 2) can also be used to produce amorphous materials.

3.3 Structural, Thermal, and Mechanical Properties

As mentioned in the previous sections, in the amorphous/BMG materials, the absence of long-range atomic order (structural properties) and the metastable state (thermal properties) gives rise to unique mechanical/surface and magnetic proper-ties that are unattainable in conventional crystalline materials.

With regard to the amorphous structure, a visual inspection of cast ingots of BMGs can quickly reveal whether the composition is amorphous or not. A high reflective surface generally indicates amorphous nature, whereas a dull gray surface generally indicates that the ingot has devitrified or crystallized. A variety of modern and advanced techniques such as XRD, TEM/HRTEM, X-ray absorption fine

Fig. 3.14 Shows representative pictures of the vial, balls, and the ball-milling machine

Fig. 3.13 Shows the mechanism of amorphization (adapted from Suryanarayana 2001) (© 2001, Elsevier. Used with permission)


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spectroscopy (XAFS), etc. are used to study the atomic structure, phase transitions, and shear band formation/deformation mechanisms. Methods such as positron anni-hilation spectroscopy is used to determine free volume content in the glass, while advanced methods like 3D atom probe can indicate the position of atoms and the nearest-neighbors, thus providing crucial insights on atomic positions, GFA, phase formation, and deformation mechanisms (Miller and Liaw 2008). Further, the TTT diagrams (can be obtained by simple isothermal measurements in a DSC) indicate the formation and stability of the amorphous phase, the driving force required, and the amorphous-to-crystal transformation. Usually, studies pertaining to the structure (atomic structure and stability) are conducted to investigate phase formations during processing/post-processing (quenching, annealing, devitrification, etc.). Through these studies, the composition, size/morphology, structure, and volume fraction of phases (if any) are identified, which eventually help in alloy design. It should be noted that in some cases, the atomic order seems to be longer than the short-range order. These are medium-range order (MRO) structures and are not representatives of nanoclusters. However, under sufficient driving force (annealing, etc.), these may act as nucleation sites to transform to crystalline/nanocrystalline structure (Sheng et al. 2006). Figures 3.15, 3.16, 3.17, 3.18, and 3.19 show representative XRD, DSC, and TEM/HRTEM images of Zr-, Ni-, Cu-, Ti-based amorphous materials/metallic glasses (Xu et al. 2004; Hofmann et al. 2006; Duan et al. 2008; Nagase et al. 2010; Zhu et al. 2012). Table 3.1 list the thermal properties of several amor-phous materials/metallic glasses available in literature (Xu et al. 2004; Duan et al. 2008; Huang et al. 2008; Zhu et al. 2012).

The unique advantage of amorphous materials/BMGs is their mechanical prop-erties and their deformation behavior with higher temperature. As mentioned ear-lier, amorphous alloys do not contain microstructural features that are observed in

Fig. 3.15 DSC curves of amorphous (a) Ti45Zr20Be35 (b) Ti45Zr20Be20Cr5 and (c) Ti40Zr25Be30Cr5 metallic glasses at constant heating rate of 0.33 K−1 (adapted from Duan et al. 2008) (© 2008, Elsevier. Used with permission)


74

crystalline materials, such as grain boundaries, dislocations, stacking faults, etc. It is important to briefly state the deformation theories proposed for metallic glasses: the free volume model proposed by Turnbull (Drehman et al. 1982), later developed by Speapen (2006), and the shear transformation zone (STZ) theory initially pro-posed by Argon and Kuo (1979), and further developed by Falk and Langer (Falk and Langer 1998; Langer 2006), are widely accepted. In the latter, the flow in amor-phous metals is accommodated by cooperative shearing of atomic clusters, referred

Fig. 3.16 HRTEM micrograph and diffraction pattern (inset) showing totally amorphous microstructure in Cu46Zr54 glass (adapted from Hofmann et al. 2006) (© 2006, Elsevier. Used with permission)

0.8a b0.5mmNi45

Ti20

Zr23

AI12 Ni

45Cu

6Ti

10.6Zr

23.5AI

10,

Ni40

Cu5Ti

10.5Zr

25.5AI

10,

Ni40

Cu5Ti

17Zr

28AI

10,

Ni40

Cu6Ti

16Zr

28AI

10,

Ni45

Ti20

Zr25

AI10

,

Ni45

Ti20

Zr25

AI10

Ni45

Ti20

Zr27

AI8

Ni45

Ti20

Zr35

2mm

2mm

0.5mm

3mm

3mm

4mm

5mm

splat quenched0.3

-0.2

-0.7

-1.2

-1.7

-2.2

-2.7

-3.2

700 750 800 850

Temperature (K) Two Theta (degree)

Nor

mal

ized

Hea

t Flo

w (

W/g

)E

xoth

erm

ic

Rel

ativ

e In

tens

ity (

a.u.

)

900

Heating Rate = 0.33 K/s

Tg

Tx1

950 30 40 50 60 70 80 90 100

Fig. 3.17 Thermal and structural properties of Ni-base metallic glass systems: (a) DSC pattern of ternary and quaternary Ni–Ti–Zr and Ni–Ti–Zr–Al alloys at a heating rate of 0.33 K−1 and (b) XRD pattern of NixCua−xTiyZrb−yAl10 (a ~ b ~ 45 at.%) quaternary and quinary alloys (adapted from Xu et al. 2004) (© 2004, Elsevier. Used with permission)


75

20

1100 1150 1200 1250Temperature, T/K Temperature, T /K

Tg

Tx

TmTl

1300 1350 650 700 750 800 850 900

30

β−Nb

Inte

nsity

/ a.

u.

Exo

ther

mic

/ a

.u.

Exo

ther

mic

/ a

.u.

2θ / °40 50

10mm

10mm

15mm

18mm

20mm

18mm

12mm

60

Zr65

Cu20

AI7Pd

5Nb

3

Zr62

Cu23

AI7Pd

5Nb

3

Zr60

Cu25

AI7Pd

5Nb

3

Zr55

Cu3AI

7Pd

5Nb

3

Zr50

Cu35

AI7Pd

5Nb

3

Zr50

Cu35

AI7Pd

5Nb

3

Zr45

Cu40

AI7Pd

5Nb

3

Zr45+xCu40-xAI7Pd5Nb3 Zr45+xCu40-xAI7Pd5Nb3

70

x = 22x = 22x = 20

x = 20x = 17

x = 17x = 15

x = 15

x = 10 x = 10

x = 5 x = 5

x = 0x = 0

0.33 K/s 0.67 K/s

80

a b

c d

Fig. 3.18 Structural and thermal properties of Ni- and Be-free Zr-based bulk metallic glasses with high glass forming ability and unusual plasticity (a) XRD patterns of as-cast ZrCuAlPdNb alloys (b) HRTEM image of Zr55Cu30Al7Pd5Nb3 alloy (c) and (d) DTA and DSC curves of ZrCuAlPdNb alloys (adapted from Zhu et al. 2012) (© 2012, Elsevier. Used with permission)

Fig. 3.19 In situ TEM studies on glass-to-liquid transition in Fe–Zr–B BMG containing dispersed crystalline Cu globules, at various temperatures of (a) 858 K, (b) 883 K and (c) 1,023 K. Aggregated crystalline Cu globules are visible at 883 K as shown by arrows (adapted from Nagase et al. 2010) (© 2010, Elsevier. Used with permission)


76

to as STZs. However, the free volume plays an important role in that model too; STZ are likely to initiate in less packed clusters and the excess free volume tends to lower the deformation resistance of the sheared regions (Argon 1979). These theo-ries are more suitable for BMGs (i.e., glasses with larger dimensions).

Table 3.1 Thermal properties of several selected amorphous/glassy materials (Duan et al. 2008; Huang et al. 2008; Xu et al. 2004; Zhu et al. 2012)

Compositions

Critical casting thickness, d, mm Tg, K Tx, K Tm, K Tl, K ΔT, K

Ni40Ti20Zr35 0.5 725 752 – – 27Ni45Ti20Zr27Al8 <0.5 761 802 – – 41Ni45Ti20Zr25Al10 2 773 818 – – 45Ni45Ti20Zr23Al12 <0.5 783 832 – – 49Ni40Cu6Ti16Zr28Al10 3 765 807 – – 42Ni40Cu5Ti17Zr28Al10 4 762 808 – – 46Ni40Cu5Ti16.5Zr28.5Al10 5 763 809 – – 46Ni39.8Cu5.97Ti15.92Zr27.86Al9.950Si0.5 5 768 815 – – 47Ti45Zr20Be35 6 597 654 – 1,123 57Ti40Zr25Be35 6 598 675 – 1,125 76Ti45Zr20Be30Cr5 7 602 678 – 1,135 77Ti40Zr25Be30Cr5 8 599 692 – 1,101 93Ti30Zr35Be35 6 595 713 – 1,201 118Zr465Cu12.5Be22.5 4 583 684 – 1,098 99Zr41.2Ti13.8Ni10Cu12.5Be22.5 >20 623 712 – 993 89Zr46.75Ti8.25Ni10Cu7.5Be27.5 >20 625 738 – 1,185 113Zr45Cu40Al7Pd5Nb3 – 724 774 1,141 1,190 50Zr50Cu35Al7Pd5Nb3 – 710 770 1,142 1,184 60Zr55Cu30Al7Pd5Nb3 – 693 763 1,141 1,205 70Zr60Cu25Al7Pd5Nb3 – 674 737 1,138 1,234 63Zr62Cu23Al7Pd5Nb3 – 670 724 1,184 1,231 54Zr65Cu20Al7Pd5Nb3 – 655 708 1,186 1,239 43Zr67Cu18Al7Pd5Nb3 – 653 690 1,187 1,235 37Fe72Y6B22 2 898 944 1,391 1,419 46Fe71Ni1Y6B22 2 893 926 1,382 1,507 43Fe70Ni2Y6B22 2 880 925 1,377 1,509 45Fe69Ni3Y6B22 1.5 874 910 1,370 1,503 36Fe68Ni4Y6B22 1.5 872 907 135 1,470 35Fe67Ni5Y6B22 1 866 891 1,345 1,469 25Fe70Co2Y6B22 2 898 944 1,390 1,420 46Fe68Co4Y6B22 2 896 941 1,385 1,414 45Fe66Co6Y6B22 2 887 925 1,379 1,509 38Fe64Co8Y6B22 2.5 884 927 1,376 1,505 43Fe62Co10Y6B22 2.5 885 932 1,375 1,503 47

Tg glass transition temperature, Tx crystallization temperature, Tm melting temperature, Tl liquidus temperature, ΔT supercooled liquid region


77

In amorphous ribbons, the content of free volume is somewhat high due to its rapidly quenched structure. The viscoplastic deformation is dependent on the atomic mobility involving the rearrangement of a small number of atoms and their surrounding free volume (Turnbull and Cohen 1961). The larger the free volume, the easier is the atomic mobility that would favor a homogeneous flow (Argon and Kuo 1979). Also at high temperatures, short-range diffusion of atoms would cause structural relaxation and therefore a simultaneous reduction in free volume could not be neglected. Patterson and Jones (Patterson and Jones 1979) reported a reduction in activation energy for crystallization upon application of a tensile stress.

The mechanical properties of amorphous alloys/metallic glasses are strongly dependent on temperature and strain rate, and follow inhomogeneous and homog-enous deformation behaviors (Fig. 3.6) (Spaepen 2006). At room/low tempera-tures (i.e., in the glassy state, below Tg) they exhibit very high strength (>1 GPa) and an elastic strain of 2 %, as opposed to 0.2 % in crystalline metals. However, the lack of crystal defects, grains/grain boundaries cause lack of plastic deforma-tion modes resulting in poor ductility. To elaborate much further, unlike conven-tional crystalline alloys where plastic deformation is explained by the theory of dislocation, plastic deformation in amorphous alloys occurs by shear localization at low temperature (Spaepen 2006). Under mechanical loading conditions and at temperatures well below the glass transition temperature, the stress is insensitive to strain rate and the flow is characterized by a plastic and inhomogeneous behav-ior (Wang et al. 2005a). The nonelastic deformation is confined to narrow bands called “shear bands” (inhomogeneous flow), and fracture eventually occurs due to excessive localized deformation of such dominant shear bands (Zhang et al. 2003). Typical tensile and compressive stress–strain curves at room (low) tem-perature and fracture surface morphology are shown in Figs. 3.20, 3.21, 3.22, and 3.23, respectively.

In contrast, at higher temperatures, as for example observed in the supercooled liquid region or at temperatures >0.70Tg, and at low stresses, the specimens deform uniformly (homogenous flow) and extensive plastic deformation occurs (Wang et al. 2005b). However, the transition temperature from inhomogeneous to homogeneous deformation (or brittle-to-ductile transition) is strongly dependent upon strain rate indicating that homogeneous deformation is associated with the rate/diffusion- dependent process. When the variation of stress is directly propor-tional to the variation of strain rate, the viscous flow is described as Newtonian, whereas for stress varying nonlinearly with strain rate, the stress can be correlated to the strain rate by the equation,

s = k p

m�e

(3.2)

with m < 1, which is known as the non-Newtonian flow (Spaepen 1976, 2006). Several studies have reported Newtonian flow for metallic glasses (superplastic behavior) and few examples are in ref. Lu et al. (2003), Chu et al. (2003), Laws et al. (2008). Figure 3.24a–d shows the representative superplastic flow behavior in Zr-based BMG in the supercooled liquid region (Wang et al. 2005b).


78

2.0a

b

1.5

1.0

0.5

0

2.0

1.5

1.0

0.5

0

Com

pres

sive

str

ess

(GPa)

Ten

sile

Str

ess

(GPa)

Tensile strain

Compressive strain

B

A: ε=4.5 x10-5s-1

A: ε=3x10-2s-1

B: ε=3x10-2s-1

C: ε=3x10-2s-1

D: ε=3x10-2s-1

B: ε= 4.5 x10-5s-1

A

A1 %

1 %

B C D

Fig. 3.21 (a) Compressive and (b) tensile stress–strain curves of Zr59Cu20Al10Ni8Ti3 bulk metallic glass (adapted from Zhang et al. 2003) (© 2003, Elsevier. Used with permission)

2500(a)

(b) (c) (d)

(e)

(f)

(g)2000

1500

1000

Stre

ss s

/M

Pa

Strain, ε / %

(a) Zr45Cu40AI7Pd5Nb3

(b) Zr50Cu35AI7Pd5Nb3

(c) Zr55Cu30AI7Pd5Nb3

(d) Zr60Cu25AI7Pd5Nb3

(e) Zr62Cu23AI7Pd5Nb3

(f) Zr65Cu20AI7Pd5Nb3

(g) Zr67Cu18AI7Pd5Nb3

strain rate:5X10-4s-1 5%

500

0

Fig. 3.20 Compressive stress–strain curves of Zr–Cu–Al–Pd–Nb bulk metallic glass (adapted from Zhu et al. 2012) (© 2012, Elsevier. Used with permission)


79

Given these unique mechanical characteristics, Fig. 3.25 shows the comparison of mechanical properties between conventional metals, plastics, and metallic glasses (Telford 2004).

With regard to applications, BMGs are widely used in sporting goods due to their high strength and elastic limit, especially in golf clubs, tennis rackets, baseball bats, skis, bicycle parts, etc. (Telford 2004). Their excellent formability at higher tem-peratures (can be micromolded) enable them to be used in micromachines such as micromotors and microgears (Wang et al. 2005b). Further, their high strength and large elastic strain limit make them suitable candidate for spring materials (automo-tive valve spring components) as an increase in these properties imply high resil-ience. As already mentioned, due to their superior magnetic properties, they are excellent materials for transformer cores, and their resistance to hydrogen embrittle-ment, high hydrogen permeability, high hydrogen storage capacity, and corrosion resistance increase their applicative potential as permeation membranes, storage, and as separator membranes. Other applications include optical mirror devices as metallic glasses can be more reflective due to the absence of grain boundaries (scatter light).

Fig. 3.22 Typical tensile fracture surface of (a) Zr41.25Ti13.75Ni10Cu12.5Be22.5 BMG at room temperature, with fracture angle of 56°. Note the presence of cleavage type vein fracture (as seen in b) along with some molten metal droplets due to the adiabatic behavior (adapted from Wang et al. 2005a) (© 2005, Elsevier. Used with permission)


80

Fig. 3.23 Fracture surface of metallic glasses after compression test. The images are representa-tive images showing the fracture surface and outer surface of the specimens in: (a, b) Zr62Cu23Al7Pd5Nb3 BMG and (c, d) Zr65Cu20Al7Pd5Nb3 BMG, respectively (adapted from Zhu et al. 2012) (© 2012, Elsevier. Used with permission)

Another potential application of metallic glasses is in the field of biomedical implants, (e.g., Ti-BMG, has elastic modulus similar to bone) due to their excellent corrosion resistance and biocompatibility.

3.4 Limitations

Although the above-mentioned highlights their superior properties and a variety of prospective applications, the limitation in the wider use of amorphous alloys/metal-lic glasses arises due to their dimensions (limited to few tens of millimeters), their metastable nature (structural and thermal stability) which in turn could affect the mechanical behavior. Also, the absence of strain hardening in BMG (at room tem-peratures) results in catastrophic brittle failure, and poses a severe constraint in being used as a structural component.


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1800

1600

1400

1200

1000

800

600

400

200

0.00

0.00

20040060080060

120180240300

020406080

100120

(d)a

b

dc

(b)(c)

(a)

(g)(f)

(e)

(j)(i)

(h)

0.4 0.8 1.2 1.6 2.6 2.4 2.8

0.02 0.04Strain rate, ε/s-1

True strain, ε

Strain rate, ε/s-1

Elo

ngat

ion.

δ/%

Tru

e st

ress

, σ/

MPa

Tru

e st

ress

, σ/

MPa

0.06 0.08

1000

100

10

m=1.0

m=0.4

m=0.1

10-3 10-2 10-1

676 K656 K

0.10

676 K

(a) 1.52 x1 0-3s-1 (b) 3.79 x1 0-3s-1

(c) 6.06 x1 0-3s-1 (d) 7.58 x1 0-3s-1

(e) 1.14 x1 0-2s-1 (f) 1.52 x1 0-2s-1

(g) 3.03 x1 0-2s-1

(h) 4.55 x1 0-2s-1 (i) 6.06 x1 0-2s-1

(j) 9.09 x1 0-2s-1

Fig. 3.24 (a) True stress–true strain curves of Zr41.25Ti13.75Ni10Cu12.5Be22.5 metallic glass tested under tension at strain rates varying between 1.52 × 10−3 and 9.09 × 10−2 s−1 and at a temperature of 676 K, in the supercooled liquid region, (b) photograph of the specimen exhibiting superplasticity, (c) variation of elongation with stain rates in Zr41.25Ti13.75Ni10Cu12.5Be22.5 metallic glass tested at temperature in the supercooled liquid region and (d) strain rate sensitivity indicates Newtonian viscous flow behavior at specific strain rates (adapted from Wang et al. 2005b) (© 2005, Elsevier. Used with permission)

2500

2000

1500

500

0

1000

SilicaWood

Steels

Titaniumalloys

Glassy alloys

Polymers

1Elastic limit (%)

Str

engt

h (M

Pa)

2 3

Fig. 3.25 Comparison of mechanical properties between conventional metals, polymers, and metallic glasses. Note the large elastic limit in the glassy alloys, coupled with very high strength (adapted from Telford 2004) (© 2004, Elsevier. Used with permission)

3.4 Limitations

82

An effective way to utilize the superior properties of amorphous metals/metallic glasses is to incorporate them as reinforcements in metal matrices (i.e., amorphous/metallic glass reinforced metal matrix composites, particularly light metal matrices such as Al and Mg). Such an effort would result in new hybrid materials with advan-tages of both light metal matrix and amorphous metals.

References

Argon AS (1979) Plastic deformation in metallic glasses. Acta Metall 27:47–58Argon A, Kuo H (1979) Plastic flow in a disordered bubble raft (an analog of a metallic glass).

Mater Sci Eng 39:101–109Basu J, Ranganathan S (2003) Bulk metallic glasses: a new class of engineering materials. Sadhana

28:783–798Bletry M, Guyot P, Blandin JJ, Soubeyroux JL (2006) Free volume model: high-temperature defor-

mation of a Zr-based bulk metallic glass. Acta Mater 54:1257–1263. doi:10.1016/j.actamat.2005.10.054

Chrobak A, Haneczok G, Kwapuliński P et al (2006) Soft magnetic properties of the Fe72Co10Nb6B12 amorphous alloy. J Alloys Compd 423:77–80. doi:10.1016/j.jallcom.2005.12.050

Chu J, Chiang C, Mahalingam T, Nieh T (2003) Plastic flow and tensile ductility of a bulk amor-phous Zr55Al10Cu30Ni5 alloy at 700 K. Scr Mater 49:435–440. doi:10.1016/S1359-6462(03)00302-6

De Hey P, Sietsma J, Van Den Beukel A (1997) Creation of free volume in amorphous Pd40Ni40P20 during high. Mater Sci Eng A Struct Mater 228:336–340

Drehman AJ, Greer AL, Turnbull D (1982) Bulk formation of a metallic glass: Pd40Ni40P20. Appl Phys Lett 41:716. doi:10.1063/1.93645

Duan G, Wiest A, Lind ML et al (2008) Lightweight Ti-based bulk metallic glasses excluding late transition metals. Scr Mater 58:465–468. doi:10.1016/j.scriptamat.2007.10.040

Duwez P (1967) Metastable phases obtained by rapid quenching from the liquid state. Prog Solid State Chem 3:377–400

Falk ML, Langer JS (1998) Dynamics of viscoplastic deformation in amorphous solids. Phys Rev E 57:7192–7205

Fleury E, Jayaraj J, Kim YC et al (2006) Fe-based amorphous alloys as bipolar plates for PEM fuel cell. J Power Sources 159:34–37. doi:10.1016/j.jpowsour.2006.04.119

Hara S, Hatakeyama N, Itoh N et al (2002) Hydrogen permeation through palladium-coated amor-phous Zr-M-Ni (M=Ti, Hf) alloy membranes. Desalination 144:115–120

Hofmann D, Duan G, Johnson W (2006) TEM study of structural evolution in a copper mold cast Cu46Zr54 metallic glass. Scr Mater 54:1117–1122

Huang XM, Chang CT, Chang ZY et al (2008) Formation of bulk metallic glasses in the Fe–M–Y–B (M=transition metal) system. J Alloys Compd 460:708–713. doi:10.1016/j.jallcom.2007.09.063

Inoue A (1998) Bulk amorphous alloys: preparation and fundamental characteristics, vol 4. Trans Tech Publications, Zürich

Inoue A, Zhang T, Masumoto T (1989) Al–La-Ni amorphous alloys with a wide supercooled liquid region. Mater Trans 30:965–972

Jayalakshmi S, Vasantha VS, Fleury E, Gupta M (2012) Characteristics of Ni–Nb-based metallic amorphous alloys for hydrogen-related energy applications. Appl Energy 90:94–99. doi:10.1016/j.apenergy.2011.01.040

Johnson WL (1999) Bulk glass-forming metallic alloys: science and technology. MRS Bull 24:42–56. doi:10.1557/S0883769400053252


http://dx.doi.org/10.1016/j.actamat.2005.10.054



http://dx.doi.org/10.1016/S1359-6462(03)00302-6

http://dx.doi.org/10.1016/S1359-6462(03)00302-6

http://dx.doi.org/10.1063/1.93645


http://dx.doi.org/10.1016/j.jpowsour.2006.04.119



http://dx.doi.org/10.1016/j.apenergy.2011.01.040

http://dx.doi.org/10.1557/S0883769400053252

83

Kane SN, Tripathi S, Coisson M et al (2012) Microstructure and magnetic properties of (Fe100−

xCox)84.5Nb5B8.5P2 alloys. J Alloys Compd 536:S337–S341. doi:10.1016/j.jallcom.2011.11.110Kato H, Kawamura Y, Inoue A, Chen HS (1998) Newtonian to non-Newtonian master flow curves

of a bulk glass alloy Pd40Ni10Cu30P20. Appl Phys Lett 73:3665. doi:10.1063/1.122856Langer J (2006) Shear-transformation-zone theory of deformation in metallic glasses. Scr Mater

54:375–379. doi:10.1016/j.scriptamat.2005.10.005Laws KJ, Gun B, Ferry M (2008) Mechanical stability of Ca65Mg15Zn20 bulk metallic glass during

deformation in the supercooled liquid region. Mater Sci Eng A 480:198–204. doi:10.1016/j.msea.2007.07.034

Lu J, Ravichandran G, Johnson WL (2003) Deformation behavior of the Zr41.2Ti13.8Cu12.5Ni10Be22.5 bulk metallic glass over a wide range of strain-rates and temperatures. Acta Mater 51:3429–3443. doi:10.1016/S1359-6454(03)00164-2

Miller M, Liaw P (2008) Bulk metallic glasses—an overview. Springer, New YorkNagase T, Yokoyama A, Umakoshi Y (2010) In situ TEM observation of the glass-to-liquid transi-

tion of metallic glass in Fe–Zr–B–Cu alloy. Scr Mater 63:1020–1023. doi:10.1016/j.scriptamat.2010.07.037

Patterson J, Jones D (1979) The effect of homogeneous deformation on the crystallisation of a metallic glass. Script Metall 13:947–949

Samwer K, Fecht HJ, Johnson WL (1994) Amorphization in metallic systems. Top Appl Phys 72:5–64

Sheng HW, Luo WK, Alamgir FM et al (2006) Atomic packing and short-to-medium-range order in metallic glasses. Nature 439:419–425. doi:10.1038/nature04421

Sowjanya M, Kishen Kumar Reddy T (2014) Cooling wheel features and amorphous ribbon for-mation during planar flow melt spinning process. J Mater Process Technol 214:1861–1870. doi:10.1016/j.jmatprotec.2014.04.004

Spaepen F (1976) Mechanism for steady state inhomogeneous flow in metallic glasses. Acta Metall 25:407–415

Spaepen F (2006) Homogeneous flow of metallic glasses: a free volume perspective. Scr Mater 54:363–367. doi:10.1016/j.scriptamat.2005.09.046

Suryanarayana C (2001) Mechanical alloying and milling. Prog Mater Sci 46:1–184Suryanarayana C, Inoue A (2010) Bulk metallic glasses. CRC, Boca RatonTelford M (2004) The case for bulk metallic glass. Mater Today 7:36–43Turnbull D, Cohen MH (1961) Free-volume model of the amorphous phase: glass transition.

J Chem Phys 34:120. doi:10.1063/1.1731549Wang G, Shen J, Sun JF et al (2005a) Tensile fracture characteristics and deformation behavior of

a Zr-based bulk metallic glass at high temperatures. Intermetallics 13:642–648. doi:10.1016/j.intermet.2004.10.011

Wang G, Shen J, Sun JF et al (2005b) Superplasticity and superplastic forming ability of a Zr–Ti–Ni–Cu–Be bulk metallic glass in the supercooled liquid region. J Non Cryst Solids 351:209–217. doi:10.1016/j.jnoncrysol.2004.11.006

Xu D, Duan G, Johnson WL, Garland C (2004) Formation and properties of new Ni-based amor-phous alloys with critical casting thickness up to 5 mm. Acta Mater 52:3493–3497. doi:10.1016/j.actamat.2004.04.001

Yamaura S, Sakurai M, Hasegawa M et al (2005) Hydrogen permeation and structural features of melt-spun Ni–Nb–Zr amorphous alloys. Acta Mater 53:3703–3711. doi:10.1016/j.actamat.2005.04.023

Zhang Z, Eckert J, Schultz L (2003) Difference in compressive and tensile fracture mechanisms of Zr59Cu20Al10Ni8Ti3 bulk metallic glass. Acta Mater 51:1167–1179. doi:10.1016/S1359-6454(02)00521-9

Zhu S, Xie G, Qin F et al (2012) Ni- and Be-free Zr-based bulk metallic glasses with high glass- forming ability and unusual plasticity. J Mech Behav Biomed Mater 13:166–173. doi:10.1016/j.jmbbm.2012.04.011

References


http://dx.doi.org/10.1063/1.122856




http://dx.doi.org/10.1016/S1359-6454(03)00164-2




http://dx.doi.org/10.1038/nature04421

http://dx.doi.org/10.1016/j.jmatprotec.2014.04.004


http://dx.doi.org/10.1063/1.1731549

http://dx.doi.org/10.1016/j.intermet.2004.10.011

http://dx.doi.org/10.1016/j.intermet.2004.10.011

http://dx.doi.org/10.1016/j.jnoncrysol.2004.11.006



http://dx.doi.org/10.1016/S1359-6454(02)00521-9

http://dx.doi.org/10.1016/S1359-6454(02)00521-9

http://dx.doi.org/10.1016/j.jmbbm.2012.04.011


Chapter 4 Light Metal Matrix Composites with Amorphous Alloys/Bulk Metallic Glass Reinforcements (BMG)

Abstract In this chapter, the various synthesizing methods of metallic amorphous materials/bulk metallic glasses (BMG) into light metal matrices are presented, and is a new topic of research. The currently existing literature on the properties of vari-ous light metal matrix—amorphous/BMG reinforced composite systems are dis-cussed, particularly in terms of improvement in mechanical properties.

Keywords Light metal matrices • Reinforcement amorphous materials/bulk metallic glasses (BMG) • Processing methods • Selection of matrix reinforcement • Structural properties • Mechanical properties

4.1 Introduction

Reinforcing amorphous alloy/metallic glass in metal matrix is a relatively new research concept. As mentioned in Chap. 2 , the limitation of conventional ceramic reinforced metal matrix composites is the low ductility and drastic reduction in toughness. Such reduction in ductility is due to the inherent brittle nature of the ceramic reinforcement (low fracture strain) and the poor interfacial characteristics between the reinforcement/matrix. The difference in the type of bonding of the metal matrix and the ceramic particle, and the large difference in the thermal coef-fi cient of expansion between the metal matrix and ceramic reinforcement, gives rise to irregular stress fi elds across the matrix and the interface, resulting in the brittle nature of the composite. Similarly, as mentioned in Chap. 3 , amorphous alloys/metallic glasses, exhibit superior strength properties (~2 GPa) and large elastic strain limit (~2 %). However, the lack of strain hardening and the constraint of not being able to prepare materials of larger dimensions limit their use to very selective applications. In this regard, reinforcing metal matrices with amorphous/metallic glass reinforcement is a viable option to utilize the advantages of both the metal as well as that of the amorphous materials. This will particularly be signifi cant in light metal matrices such as Al and Mg which have relatively low melting point, consid-ering the metastable nature of the amorphous/glassy materials. Further, the rein-forcement/matrix bonding will be predominantly metallic in nature thereby reducing

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the negative effects experienced in conventional ceramic reinforced composites. With this viewpoint, recently research efforts have been made to synthesize amor-phous/glass reinforced light metal matrix composites and to understand its struc-tural/mechanical characteristics. In the following section, the synthesis, selection of matrix/reinforcement, and properties of light metal matrix composites reinforced with amorphous alloy/metallic glasses are presented.

4.2 Synthesis, Matrix Reinforcement Selection, and Properties

4.2.1 Liquid State Processing: Infi ltration Method

Lee et al. ( 2004 ) was the fi rst to synthesize light metal matrix reinforced with amor-phous/glassy materials. They employed infi ltration casting technique to produce the metallic glass reinforced Al-alloy matrix composite. As the matrix, Al-alloy was selected due to its ductile nature. The matrix materials, Al–6.5Si–0.25Mg (wt%) alloy (A356 alloy), exhibit excellent castability and can ensure the fabrication of sound, defect- free composites (macroscale casting defects such as shrinkage and porosity) during the infi ltration process. As the infi ltration process is a liquid state process, and that the melting point of Al-alloys are <660 °C, it is critical to select the proper amorphous/metallic glass reinforcement, considering the metastable nature of the glassy materials. Hence, the amorphous alloy selected should have excellent thermal stability against crystallization. In the work, Lee et al. ( 2004 ) selected Ni 39.2 Nb 20.6 Ta 40.2 (wt%) alloy as it has very high thermal stability, with a crystalliza-tion onset temperature of 994 K (721 °C), which is higher than the liquidus tem-perature of the A356 alloy (886 K, i.e., 613 °C). In fact Ni-based amorphous systems have the highest thermal stability and Ni–Nb–Ta system exhibit the best thermal properties (Lee et al. 2003 ). Hence, it can be expected that the amorphous nature of the Ni–Nb–Ta alloy will be retained, without being affected by the infi ltration pro-cess (i.e., amorphous-to-crystal transformation would not occur).

In the work by Lee et al. ( 2004 ), Ni–Nb–Ta ternary alloys of the above- mentioned composition were prepared by arc melting Ni (99.99 wt%), Nb (99.9 wt%), and Ta (99.99 wt%) under an Ar gas atmosphere. From the obtained ingots, amorphous ribbon specimens (thickness: ~30 μm, width: ~1 mm) were produced by melt spin-ning process (see Chap. 3 for arc-melting and melt-spinning methods). As the light metal-amorphous/BMG composite production is by infi ltration method, preform of amorphous reinforcements was produced, by cold-pressing of the ribbons (~3 g in weight) at a pressure of 16 MPa into cylindrical preforms with diameter: 9 mm and height: 15 mm. The amorphous preform was then placed in quartz tube of inner diameter 10 mm, upon which the A356 Al-alloy ingots were placed. The A356 alloy along with the preform was heated in vacuum (10 −3 Torr). Upon the melting of the alloy, the molten metal was pressure infi ltrated (using Ar gas at 53 kPa) and held at

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953 ± 5 K for 600 s, to produce the hybrid amorphous alloy/metal matrix composite. Figure 4.1 shows the schematic of the composite production method (Lee et al. 2004 ). The volume fraction of the amorphous reinforcement used was 20 %. The structural, thermal, and mechanical properties of the composite so produced were then investigated.

Structural analysis of the prepared composite was studied using XRD in com-parison with that of the amorphous ribbon and the base Al-alloy (Fig. 4.2 ) (Lee et al. 2004 ). It showed retention of amorphous structure of the reinforcements even after infi ltration processing, proving that infi ltration processing is one effective route to make amorphous materials/BMG reinforced light metal matrix composites, pro-vided that the amorphous material has high thermal stability. Optical microscopy studies showed homogeneous distribution of the amorphous reinforcements and absence of macroscale defects, again indicating that the process methodology and the process parameters are suitable to obtain a defect-free composite. Also, no new additional phases were observed highlighting the thermal and structural stability of the reinforcement. Further, thermal analysis by DSC (Fig. 4.3 ) (Lee et al. 2004 ) indicated the presence of one endothermic and two exothermic peaks in the com-posite. The endothermic peak (835–897 K) indicated the melting range of the base alloy, whereas the exothermic peak showed the amorphous-to-crystallization trans-formation temperature. A decrease in temperature of crystallization was found in the amorphous phase of the composite, when compared to the amorphous ribbon, and was attributed to the relaxation effect due to the infi ltration process.

Fig. 4.1 Schematic of the infi ltration process employed to produce A356 Al-alloy reinforced with Ni–Nb–Ta amorphous reinforcements ( adapted from Lee et al. 2004 ) ( © 2004, Elsevier. Used with permission )


Fig. 4.2 Structural investigation by XRD on the A356/Ni–Nb–Ta amorphous materials reinforced composite, showing retention of amorphous structure in the composite ( adapted from Lee et al. 2004 ) ( © 2004, Elsevier. Used with permission )

Fig. 4.3 Thermal investigation by DSC on the A356/Ni–Nb–Ta amorphous materials reinforced com-posite, showing relaxation of the reinforced amorphous ribbon due to the temperature involved in the infi ltration process, indicated by the lowering of the onset of temperature for crystallization (as seen in the exothermic peak at A ) ( adapted from Lee et al. 2004 ) ( © 2004, Elsevier. Used with permission )

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With regard to the mechanical properties, microhardness values were found to increase due to amorphous reinforcement addition (81 H v in matrix of the composite and 751 H v on the ribbon in the composite). Compression tests on the composites showed signifi cant improvement in yield and ultimate strengths (~26 % improve-ment in yield strength), as shown in Fig. 4.4 (Lee et al. 2004 ).

4.2.2 Solid State Processing: Powder Metallurgy-Based Methods

As discussed above, the crystallization temperature of the amorphous alloy/metallic glass is very critical in selecting the method during composite preparation. Most of the metallic glasses (except for Ni-based systems) have crystallization temperature lower than the melting point of Al- and Mg-alloys. During solid state processing (sintering methods), the sintering temperature/time plays a very important role.

Eckert et al. ( 2008 ) prepared pure Al reinforced with Ni 60 Nb 40 (at.%) amorphous particle reinforced composite by sintering methods. Mechanical alloying technique (refer Chap. 3 ) was used to prepare the Ni 60 Nb 40 amorphous powders. After powder production, pure Al was reinforced with Ni 60 Nb 40 amorphous powders of 30 % vol-ume fraction and was either sintered (conventional sintering in a furnace), hot pressed or hot extruded. In order to avoid crystallization of the amorphous rein-forcement during processing, the temperature used during the process was ~823 K,

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Fig. 4.4 Compressive stress–strain curve of Ni–Nb–Ta amorphous ribbon reinforced A356 alloy composite showing improvement in mechanical properties ( adapted from Lee et al. 2004 ) ( © 2004, Elsevier. Used with permission )


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which is ~50 K less than the crystallization onset temperature of the amorphous powder. The structural properties of the as-received powder and that of the compos-ites showed no amorphous-to-crystalline transformation of the reinforcement, and that only the peaks of crystalline Al were present. Hardness and compression tests showed signifi cant enhancement in mechanical properties, which was attributed to the inherent high strength of the Ni 60 Nb 40 alloy (strength ~2 GPa and hardness ~800 H v ) and are shown in Table 4.1 .

The effect of amorphous/glassy reinforcements on Mg-matrix was investigated by Dudina et al. ( 2009 ). In this work, AZ31 Mg-alloy was selected as a base mate-rial. The AZ91 ingots prepared by induction melting under argon were melt spun into ribbons of thickness 30 μm. The reinforcement material, i.e., Vitraloy 6 of com-position Zr 57 Nb 5 Cu 15.4 Ni 12.6 Al 10 glassy alloy was also in the form of ribbons. The glassy ribbons were cut and milled under argon in a low-speed vibratory mill for 15 h to produce glassy powders. In order to make the composite, the melt spun cut Mg-alloy ribbons were reinforced with 15 % volume fraction of the milled Zr-glassy powder and were again subjected to milling in the vibratory mill at similar condi-tions. The composite powder so prepared was then sintered in vacuum using high frequency induction heating (maximum sintering temperature: 713 K, heating rate: 1.6 K/s) at a pressure of 50 MPa, for a duration of 120 s. It should be noted that in this work, the sintering temperature was selected so that it is within the supercooled liquid region of the Vitraloy 6 (i.e., glassy reinforcement) and close to the solidus temperature of the AZ91 matrix alloy. The composites thus prepared were examined for their structural and mechanical properties. X-ray diffraction pattern of the com-posite showed retention of the amorphous structure and crystalline phases corre-sponding to Mg and Mg 17 Al 12 alone were identifi ed (Fig. 4.5 ) (Dudina et al. 2009 ). It highlighted the effi ciency of the processing method, i.e., high frequency induction heating under pressure, which is a rapid sintering process, and can prevent amorphous- to-crystalline transformation. Microstructure of the composite as observed in SEM also showed uniform distribution of the glassy reinforcement par-ticles with no pores or reaction layers at the interface, indicating sound sintered product. With the samples being 4 mm long with a cross-section of 2 × 2 mm 2 , microhardness and compression tests were conducted and a signifi cant improve-ment in the strength properties are observed. The mechanical properties are tabu-lated in Table 4.2 (Dudina et al. 2009 ).

Table 4.1 Mechanical properties of Ni 60 Nb 40 (at.%) amorphous particles reinforced pure Al-matrix composites prepared by sintering, hot pressing, and hot extrusion at 823 K

Processing Young’s modulus (GPa) Compressive yield strength (MPa)

Al-30 % V f Ni 60 Nb 40 (at.%) reinforced composites Furnace sintered 65.2 94 Hot pressed 67.3 106 Hot extruded 72.2 134

The data was obtained from compression tests (Eckert et al. 2008 )


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Fig. 4.5 XRD pattern of AZ91 Mg-alloy reinforced with Zr-base metallic glass produced by high frequency induction heating under pressure. Note the presence of amorphous halo at 2 θ = 37° and the crystalline peaks of Mg and the intermetallic phase (Mg 17 Al 12 ) ( adapted from Dudina et al. 2009 ) ( © 2009, Elsevier. Used with permission )

Table 4.2 Mechanical properties of AZ91 Mg-alloy reinforced with 15 % V f of Vitraloy 6 (i.e., Zr 57 Nb 5 Cu 15.4 Ni 12.6 Al 10 glass (at.%)) and A520 Al-alloy reinforced with 15 % V f of Cu 54 Zr 36 Ti 10 (at. %), produced by high-frequency induction heating under pressure

Material Hardness ( H v)

Compressive yield strength (MPa)

Compressive fracture strength (MPa)

Compressive deformation (%)

Pure Al 682 143 404 21.0 AZ91-15 % V f Zr 57 Nb 5 Cu 15.4 Ni 12.6 Al 10 glass reinforced composites

123 325 542 10.5

A520 Al-alloy – 190 – – A520-15 % V f Cu 54 Zr 36 Ti 10 glass reinforced composites

– 580 840 14

6061 Al-alloy – 180 – – 6061-15 % V f [(Fe 0.5 Co 0.5 ) 75 B 20 Si 5 ] 96 Nb 4 glass particle reinforced composites

– 167 570 13

The data shown here were obtained from compression tests. Note that in A520 alloy and 6061 alloy matrices, there is no fracture (even for >20 % strain) and hence the ultimate strength and fracture strain values are not provided (Dudina et al. 2009 , 2010 ; Aljerf et al. 2012 )


92

In a similar kind of work, Dudina et al. ( 2010 ) reinforced Al-alloy A520 with Cu-based Cu 54 Zr 36 Ti 10 (at.%) glassy alloy (of 15 % V f ) by using high frequency induction sintering under pressure. The metallic glass ribbons prepared by arc melt-ing/melt spinning were powdered in a vibratory mill for 24 h and mixed with the cut ribbons of Al-matrix (procedure similar to that in ref. (Dudina et al. 2009 )). The alloy-reinforcement mixture was milled for 8 h and processed using high frequency induction sintering under pressure (50 MPa) and time 120 s, at a temperature of 720 K. The sintering temperature was selected after considering the T g (715 K) and T x (760 K) of the Cu 54 Zr 36 Ti 10 (at.%) amorphous ribbon (Dudina et al. 2010 ). It should be noted that the selected temperature was within the supercooled liquid region of the glassy alloy and that the Al-matrix would be in a soft state at that sin-tering temperature (close to the solidus temperature). Structural analysis using XRD indicated no additional phases and transformation of the amorphous phase (Fig. 4.6 ). Further, uniform distribution of the amorphous phase with no preferred orientation is recorded. The compressive properties showed a signifi cant improvement in strength with slight reduction in the fracture strain (Fig. 4.7 ). The strengthening in the composite was attributed to the load transfer effect (high load carrying capacity of the amorphous reinforcement), dislocation strengthening effect due to the prepa-ration method, Orowan strengthening, grain refi nement (due to melt spinning and ball-milling of the matrix alloy) and solid solution strengthening due to possible diffusion of Cu from the reinforcement into the Al-matrix (Dudina et al. 2010 ). Table 4.2 lists the compressive properties of the Cu-BMG reinforced Al-composite.

The high frequency induction sintering under pressure was also used by Aljerf et al. ( 2012 ) to reinforce Al6061 alloy with [(Fe 0.5 Co 0.5 ) 75 B 20 Si 5 ] 96 Nb 4 glassy parti-cles, except that the process parameters, i.e., the milling time, sintering temperature, and time were different. Also, the milling of the glassy ribbons was done using ceramic balls (for 6 h) as the glassy ribbons were harder than the stainless steel

Fig. 4.6 XRD pattern of A520 Al-alloy reinforced with Cu 54 Zr 36 Ti 10 (at.%) glassy alloy produced by high frequency induction heating under pressure. Note the presence of amorphous halo at 2 θ = 41° ( adapted from Dudina et al. 2010 ) ( © 2010, Elsevier. Used with permission )


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balls. Similarly, considering the T g and T x of the alloy ( T g : 821 K and T x : 861 K), the induction sintering temperature was 828 K, consolidated at 70 MPa. The DSC curves (Fig. 4.8a ) and XRD patterns (Fig. 4.8b ) show the retention of amorphous structure in the composites and are also compared with the unreinforced alloy and the glassy material. The stress–strain behavior of the composite (Fig. 4.9 ) indicates excellent improvement in strength properties. The mechanical property data are listed in Table 4.2 (Aljerf et al. 2012 ).

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Fig. 4.7 Compressive stress–strain curve of A520 Al-alloy reinforced with 15 % V f of Cu 54 Zr 36 Ti 10 (at. %) glassy alloy showing improvement in mechanical properties, with good ductility ( adapted from Dudina et al. 2010 ) ( © 2010, Elsevier. Used with permission )

Fig. 4.8 ( a ) DSC curves and ( b ) XRD patterns showing the retention of amorphous structure in the [(Fe 0.5 Co 0.5 ) 75 B 20 Si 5 ] 96 Nb 4 glassy particles reinforced Al 6061 composites. A comparison has been made with the unreinforced alloy and the glassy material ( adapted from Aljerf et al. 2012 ) ( © 2012, Elsevier. Used with permission )


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The synthesis, structure, compressive properties, and the deformation behavior of Al-matrix reinforced with high volume fractions of Zr 57 Ti 8 Nb 2.5 Cu 13.9 Ni 11.1 Al 7.5 glassy powder (40 and 60 % volume fraction), prepared using powder metallurgy was studied by Scudino et al. ( 2009 ). In this investigation, Zr 57 Ti 8 Nb 2.5 Cu 13.9 Ni 11.1 Al 7.5 glassy powder was prepared by mechanically alloying elemental powder mixtures of the required composition in a planetary ball-mill with hardened steel balls and vials in an argon atmosphere, at a ball-to-powder ration of 13:1 and a milling time of 120 h and at 150 rpm. In order to prepare the composites, pure Al was reinforced with 40 and 60 % volume fraction of the Zr 57 Ti 8 Nb 2.5 Cu 13.9 Ni 11.1 Al 7.5 glassy powder. The Al-glassy powder mixture was consolidated by hot pressing followed by hot extrusion (under argon atmosphere and an extrusion ratio of 6:1) at a temperature of 673 K and pressure of 500 MPa. Microstructure and compression properties were investigated, with the dimensions of the compression test sample being 8 mm length and 4 mm diameter. It should be noted that the glassy powder has a glass transition temperature of 668 K and two crystallization temperatures with onset values being 716 K and 757 K, respectively. The sintering temperature was hence slightly lower than the T g of the metallic glass reinforcement. The XRD study indicated the reten-tion of amorphous structure and the presence of only pure Al-crystalline peaks in the composite, indicating no occurrence of crystallization. Scanning electron micro-scope images (Fig. 4.10 ) (Scudino et al. 2009 ) showed uniform distribution of the reinforcement particles and no porosity (Fig. 4.10a ). Grouping of particles (Fig. 4.10b ) was observed at several locations due to the high volume fraction. The compressive stress–strain curve of the composites is shown in Fig. 4.11 , and is also compared with that of the melt spun Zr 57 Ti 8 Nb 2.5 Cu 13.9 Ni 11.1 Al 7.5 ribbon (Scudino et al. 2009 ). In comparison to pure Al and the composite with 40 % V f of glassy reinforcement, the stress–strain behavior of the 60 % V f composite indicate work softening behavior. Further, while the strength of the composite is lower than the strength of the metallic glass, the large plastic deformation in the composite is attributed to the ductile nature of the Al-matrix, which is absent in the fully

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amorphous material. Also, the elastoplastic deformation of the composites was modeled using the self-consistent effective medium approximation (SECMA) methods. The shear log model was used for theoretical estimation of yield strength and was compared with that of the experimental values (Scudino et al. 2009 ).

Jayalakshmi et al. ( 2013 , 2014 ) prepared pure Al- and Mg-based composites reinforced with Ni 60 Nb 40 (at.%) amorphous particles with varying volume fractions. Similar to the unconventional sintering method employed in ref. Dudina et al. ( 2009 , 2010 ), Aljerf et al. ( 2012 ), they used bidirectional microwave sintering method to sinter the composites. Amongst the powder metallurgy-based techniques, microwave- assisted rapid sintering has several advantages (Gupta and Eugene 2007 ). The tech-nique employs two-directional heating to sinter materials, through a combined action of microwaves and microwave-coupled external heating source. By this method, sintering temperatures close to melting point of light metals such as Al, Mg can be achieved in a shorter period of time that results in dense products with relatively larger dimensions. However, as the sintering time is very short (10–15 min max), the high temperature during sintering would not affect the amorphous state of

Fig. 4.10 SEM images showing ( a ) uniform distribution and ( b ) grouping of particles due to high volume fraction, in Zr 57 Ti 8 Nb 2.5 Cu 13.9 Ni 11.1 Al 7.5 glassy powder reinforced Al-matrix composite ( adapted from Scudino et al. 2009 ) ( © 2009, Elsevier. Used with permission )


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the reinforcing materials and in addition would ensure effi cient bonding between the reinforcement and the matrix. In the investigation, fi rstly, Ni 60 Nb 40 (at.%) amor-phous alloy powder was prepared by mechanically alloying powder mixtures of elemental Ni and Nb metals. The powder mixture was milled at room temperature in air for 87 h, using a planetary ball-mill with a ball-to-powder ratio of 3:1 and milling speed of 200 rpm. For composite production, the reinforcement particles of specifi c volume fraction are mixed with pure Al and Mg matrices, respectively. The powders were blended for a duration of 1 h and then cold compacted (consolidated) at 450 MPa for 1 min. The compacted billets were of the dimension 36 mm in diameter and 50 mm in height. The compacted billets were then placed in a microwave heat-ing setup, which can provide bidirectional heating (Gupta and Eugene 2007 ), and were microwave sintered at 100 % power level for 12 min and 30 s, so as to reach a sintering temperature of 550 °C (based on a prior calibration). The sintered billet was then soaked at 400 °C for 1 h, after which it was hot extruded at 350 °C at an applied pressure of 500 MPa. An extrusion ratio of 20.25:1 was used and the result-ing composites were of 8 mm in diameter and suffi ciently long to investigate the tensile properties (gauge length 25 mm, diameter: 5 mm). Indeed, the report on the tensile behavior of amorphous alloy/glass reinforced composites was the fi rst of its kind. Structural, electrical, and mechanical properties of these composites were investigated.

In Al-Ni 60 Nb 40 (at.%) amorphous particles reinforced composite of volume frac-tions 5 %, 15 %, and 25 %, respectively (Jayalakshmi et al. 2013 ), structural analy-ses revealed uniform distribution of reinforcement, absence of interfacial products,

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Fig. 4.11 Compressive stress–strain curve of Zr 57 Ti 8 Nb 2.5 Cu 13.9 Ni 11.1 Al 7.5 glassy powder rein-forced Al-matrix composite showing improvement in mechanical properties. The stress–strain curve of the glassy ribbon is also referred to show the superior properties of the reinforcement ( adapted from Scudino et al. 2009 ) ( © 2009, Elsevier. Used with permission )


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structural dilatation and change in aspect ratio of reinforcement at 25 % V f , and retention of amorphous structure of reinforcement at all V f . Figures 4.12a–d and 4.13 show the distribution of the reinforcements and the XRD pattern of the composites (Jayalakshmi et al. 2013 ). The retention of amorphous structure even at the highest volume fraction was evident. The change in aspect ratio and resulting structural dilatation of the 25 % V f composite was attributed to the local stress variations and temperature gradients within the composite arising in situ during hot extrusion, due to reduced interparticle spacing. With regard to electrical resistivity, the composites showed higher resistivity values when compared to pure Al, due to the disordered structure of the amorphous phase. However, when the resistivity value of the amor-phous Ni–Nb alloy was considered, the resistivity is reduced to larger extent (Jayalakshmi et al. 2013 ). The mechanical properties of the prepared composites are listed in Tables 4.3 and 4.4 , and the stress–strain curve under both tension and com-pression are given in Fig. 4.14 (Jayalakshmi et al. 2013 ). The microhardness value increase by 130 % in the 25 % V f composite when compared to pure Al. With regard to the compressive yield strength, all the composites show ~45–100 % enhancement in strength, and that even at 50 % strain, the composites do not fracture indicating

Fig. 4.12 FESEM images showing uniform distribution of amorphous reinforcement particles in pure Al-Ni 60 Nb 40 (at.%) amorphous particles reinforced composite with, ( a ) 5% V p , ( b ) 15 % V p , and ( c ) 25% V p . ( d ) Representative image showing particle/matrix interface free of any interfacial products (Note V p is particle volume fraction, V f ) ( adapted from Jayalakshmi et al. 2013 ) ( © 2013, Elsevier. Used with permission )


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Fig. 4.13 XRD patterns showing the retention of amorphous structure in the pure Al-Ni 60 Nb 40 (at.%) amorphous particles reinforced composite with varying volume fraction of reinforcements ( adapted from Jayalakshmi et al. 2013 ) ( © 2013, Elsevier. Used with permission )

Table 4.3 Compressive properties of Al-Ni 60 Nb 40 (at.%) amorphous particle reinforced composites

Material Microhardness ( H v )

Compressive yield strength (MPa)

Ultimate compressive strength (MPa)

Failure strain (%)

Pure Al 54 80 245 >50 Al-5 % V f Ni 60 Nb 40 (at.%) composite

75 114 300 >50

Al-15 % V f Ni 60 Nb 40 (at.%) composite

103 125 333 >50


125 155 375 >50

Note that pure Al as well as the composites did not fracture even for strains >50 % (Jayalakshmi et al. 2013 )

Table 4.4 Tensile properties of Al-Ni 60 Nb 40 (at.%) amorphous particle reinforced composites

Material Tensile yield strength (MPa)

Ultimate tensile strength (MPa) Failure strain (%)

Pure Al 65 75 25.8 Al-5 % V f Ni 60 Nb 40 (at.%) composite

50 80 16.8


75 85 18


102 120 9.5

Note that this is the fi rst of the works in which the tensile behavior of these novel hybrid materials have been conducted and presented (Jayalakshmi et al. 2013 )


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good compressive ductility (Fig. 4.14a ). Figure 4.15 compares the compressive properties of composites produced by microwave sintering and those prepared by other methods, (as explained in the preceding paragraphs), and are found to be supe-rior/compatible with those data. Under tensile loading (Fig. 4.14b ), the yield and ultimate strengths do not increase with volume fraction, but seem to depend on a critical volume fraction after which strength improvements occur (~15–60 % increase when compared to pure Al). The ductility reduction is not drastic, which is usually observed in ceramic reinforced composites, and much higher than the amor-phous ribbon, which shows no/little plastic deformation. The tensile fracture surface of the composite Fig. 4.16a, b was characterized by ductile features, good interfacial bonding, and no cracking reinforcement (except for particle cracks/breakage in the 25 % V f composite) (Jayalakshmi et al. 2013 ).

In Mg-Ni 60 Nb 40 (at.%) amorphous particles reinforced composite of volume frac-tions 3 %, 5 %, and 10 %, respectively, the composites exhibited refi nement in grain size (Table 4.5 ), uniform distribution of reinforcement at low volume fraction, and agglomeration at high volume fractions (Fig. 4.17a, b ), clear reinforcement/matrix interface (Fig. 4.17c ), and retention of the amorphous structure at all volume fractions (Fig. 4.18a ) (Jayalakshmi et al. 2014 ). Further, from the XRD results it was inferred that the reinforcement change the crystal orientation, i.e., in the composites, the basal

3400

3200

3000

2300Compression Data of NI-base BMG [4]

Tensile Data of NI-baseAm or ph ous Ribbon [23]

Compression Tests8 x 10-4s-1 8 x 10-4s-1

Tensile Tests

Pure AI Pure AI

2600

500

400

300

200

100

0

0 00

100

200

25% VP

Pure AIAI + 5% Vp

AI + 15% Vp

AI + 25% Vp

Pure AIAI + 5% Vp

AI + 15% Vp

AI + 25% Vp

25% VP

15% VP

5% VP

5% VP

15% VP

1200

1400

5 10 15 20 25 3010 20

Engineering Strain, % Engineering Strain, %

30 40 50

e : e :∑ ∑

Fig. 4.14 Engineering stress–strain curve of pure Al and Al-Ni 60 Nb 40 (at.%) particle reinforced composites under: ( a ) compression and ( b ) tensile loading conditions. A comparison with that of a Ni-base BMG/amorphous alloy ribbon is also made ( adapted from Jayalakshmi et al. 2013 ) ( © 2013, Elsevier. Used with permission )


100

Fig. 4.15 The variation of strength ratio (compressive yield strength) with volume fraction in Al-Ni 60 Nb 40 (at.%) particle reinforced composites. The values are also compared with those of amorphous glassy particle reinforced composites existing in literature (discussed in the previous paragraphs) ( adapted from Jayalakshmi et al. 2013 ) ( © 2013, Elsevier. Used with permission )

Fig. 4.16 Fracture surface of Al-Ni 60 Nb 40 (at.%) particle reinforced composites showing ( a ) good interface bonding and ductile features at lower volume fractions and ( b ) ductile features, particle cracking and typical vein pattern fracture characteristics of the amorphous phase at the highest volume fraction (25 % V f ) ( adapted from Jayalakshmi et al. 2013 ) ( © 2013, Elsevier. Used with permission )


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planes were not entirely aligned parallel to the extrusion direction; rather they were oriented at an angle indicating multiple crystallographic orientations (Fig. 4.18b ) (Jayalakshmi et al. 2014 ). The hardness and compression test properties listed in Table 4.5 showed a remarkable increase in properties with increasing V f , owing to the inherent high hardness, strength, and elastic strain limit of the amorphous reinforce-ment. The stress–strain curve is given in Fig. 4.19 . The enhanced properties highlight that when compared to amorphous reinforcements, conventional ceramic reinforce-ments, amorphous reinforcements are better alternatives. Further, the retention of amorphous structure in both Al and Mg composites prove that the microwave sinter-ing method followed by hot extrusion was effective in synthesizing MMCs with amorphous alloy particle reinforcements (Jayalakshmi et al. 2013 , 2014 ).

In a recent work, Zheng et al. ( 2014 ) prepared 2024 Al-alloy reinforced with Fe-BMG particles using powder metallurgy method. Sintering by induction method was employed. Initially the 2024 alloy powder and the Fe 73 Nb 5 Ge 2 P 10 C 6 B 4 glassy powders were prepared using gas atomization and water atomization, respectively. The prepared powders were then milled together using steel balls, with a ball-to- powder ratio of 10:1, at 480 rpm, under argon atmosphere with stearic acid process control agent. The composite powders were consolidated under vacuum in a stain-less steel die of inner diameter 20 mm, using induction heating sintering under a pressure of 400 MPa. Sintering temperature was 823 K with a holding time of

Table 4.5 Grain size, microhardness, and compressive properties of Mg-Ni 60 Nb 40 (at.%) amorphous particle reinforced composites (Jayalakshmi et al. 2014 )

Material Grain size (μm)

Microhardness ( H v )



Failure strain (%)

Pure Mg 22.0 ± 6 43 70 265 16.2 Mg-3 % V f Ni 60 Nb 40 (at.%) composite

9.4 ± 4 62 85 283 17.6

Mg-5 % V f Ni 60 Nb 40 (at.%) composite

11.0 ± 6 84 130 320 18.4

Mg-10 % V f Ni 60 Nb 40 (at.%) composite

12.5 ± 7 95 90 322 17.2

Fig. 4.17 SEM images of Mg-Ni 60 Nb 40 (at.%) particle reinforced composites showing ( a ) uniform distribution at low volume fraction and ( b ) areas of reinforcement agglomeration at the highest volume fraction used. ( c ) The interface at all volume fractions are clear and are free of any reaction products ( adapted from Jayalakshmi et al. 2014 ) ( © 2014, Elsevier. Used with permission )


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Fig. 4.18 XRD pattern of Mg-Ni 60 Nb 40 (at.%) particle reinforced composites showing retention of amorphous structure at all volume fractions and ( b ) change in crystal orientation due to amorphous reinforcement addition ( adapted from Jayalakshmi et al. 2014 ) ( © 2014, Elsevier. Used with permission )


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30 min. The sintered billet was hot extruded at 823 K at an extrusion ratio of 10:1. Structural analysis indicated that nanostructure was attained in the Al-2024 matrix due to mechanical alloying. TEM analysis indicated a clear interface between the BMG particle and the nanostructured Al-matrix (Fig. 4.20 ) (Zheng et al. 2014 ). The compression properties of the produced composites were conducted and the stress–strain curve is shown in Fig. 4.21 (Zheng et al. 2014 ). The composites exhibited

Fig. 4.19 Engineering stress–strain curve of pure Mg and Mg-Ni 60 Nb 40 (at.%) particle reinforced composites under compression showing improvement in strength properties ( adapted from Jayalakshmi et al. 2014 ) ( © 2014, Elsevier. Used with permission )

Fig. 4.20 TEM image showing the clear interface in the composite containing 2024 Al-alloy matrix reinforced with Fe 73 Nb 5 Ge 2 P 10 C 6 B 4 glassy particles ( adapted from Zheng et al. 2014 ) ( © 2014, Elsevier. Used with permission )


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yield and fracture strength of 403 and 660 MPa, respectively, with a fracture strain of ~12 %, resulting in enhanced toughness. The superior properties were attributed to the nanostructure of Al-matrix and uniform distribution of the amorphous rein-forcement particles.

4.2.3 Other Methods/Systems

The above-mentioned solid state-processed amorphous alloy/metallic glass- reinforced light metal composites were prepared by sintering methods. Fujii et al. ( 2011 ) prepared pure Al reinforced with Fe-BMG using friction stir processing, which is a modifi cation of the friction stir welding process (refer to Chap. 2 ). As the friction stir welding temperature of the Al-alloys is in the range of 300–450 °C, amorphous/BMG reinforcement with T g higher than that temperature was selected by the authors. For this reason, Fe-based glassy particles of composition Fe 72 B 14.4 Si 9.6 Nb 4 were used. Pure Al-plates of dimensions 300 × 70 × 5 mm 3 was used as the base matrix. In friction stir processing, the tools usually consist of a shoulder (larger diameter of ~15 mm) and probe (smaller diameter of ~6 mm; length ~4.3 mm; and 10° recessed shoulder surface). A screw-type probe was used with the tool tilted by 3° during the process. For the dispersion of the reinforcements, a gap

Fig. 4.21 Engineering stress–strain curve of 2024 Al-alloy matrix reinforced with Fe 73 Nb 5 Ge 2 P 10 C 6 B 4 glassy particles under compression loading condition, that is also compared with the behavior of as-atomized 2024 alloy and milled 2024 alloy ( adapted from Zheng et al. 2014 ) ( © 2014, Elsevier. Used with permission )


http://dx.doi.org/10.1007/978-3-319-15016-1_2

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was intentionally made by placing a shim of 2 mm thickness between the two plates. The reinforcement particles were fi lled into the gap and initially FSP was performed without probe so as to prevent fl ying of particles and seal the particle in the surface. Next, FSP performed with the probe for single or multiple passes through the length of the fi lled area in the base matrix. Structural and mechanical studies were con-ducted and it was revealed that the dispersed Fe-based metallic glass particle reduced the coarsening of Al-grains. However, although the hardness improved, the dispersion had little effect on the hardness. With a change in the number and/or the speed of pass, the reaction between pure Al and Fe-based metallic glass particles increased. This resulted in the formation of Al 13 Fe 4 precipitates that improved the hardness in the stir zone.

Other works include the synthesis of nanocrystalline aluminum matrix com-posites using hot extrusion of cryomilled 5083 Al-alloy reinforced with Al 85 Ni 10 La 5 amorphous alloy powder (Zhang et al. 2006 ). In this work, the com-pression yield strength of the as-extruded composite with volume fractions 10 % and 20 % was 813 and 906 MPa, respectively. In a similar work, Al-matrix was reinforced with 20 and 40 % volume fraction of nanocrystalline Al 70 Ti 20 Ni 10 particles. Mechanical properties indicated a signifi cant improvement in com-pressive properties, with fracture strain varying between 28 and 43 % (Scudino et al. 2010 ).

4.3 Conclusions

As observed in most of these studies, interfacial properties (or rather the absence of interfacial reaction products) have contributed much towards the improvement of composite properties, particularly when compared to the conventional metal matrix composites. The inherent superior mechanical properties of the amorphous/glassy materials result in the superior performance of the composites. In essence, the amorphous/glass reinforced light metal matrix composites should be designed so as to compensate for the disadvantages faced by conventional MMCs and at the same time utilize the superior properties of the glassy materials. A synergistic effect of superior properties of the matrix and reinforcement is expected to give rise to novel and advanced hybrid materials.

References

Aljerf M, Georgarakis K, Louzguine-Luzgin D et al (2012) Strong and light metal matrix compos-ites with metallic glass particulate reinforcement. Mater Sci Eng A 532:325–330. doi: 10.1016/j.msea.2011.10.098

Dudina DV, Georgarakis K, Li Y et al (2009) A magnesium alloy matrix composite reinforced with metallic glass. Compos Sci Technol 69:2734–2736. doi: 10.1016/j.compscitech.2009.08.001

4.3 Conclusions




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Dudina DV, Georgarakis K, Aljerf M et al (2010) Cu-based metallic glass particle additions to signifi cantly improve overall compressive properties of an Al alloy. Compos Part A Appl Sci Manuf 41:1551–1557. doi: 10.1016/j.compositesa.2010.07.004

Eckert J, Calin M, Yu P et al (2008) Al-based alloys containing amorphous and nanostructured phases. Rev Adv Mater Sci 18:169–172

Fujii H, Sun Y, Inada K et al (2011) Fabrication of Fe-based metallic glass particle reinforced Al-based composite materials by friction stir processing. Mater Trans 52:1634–1640. doi: 10.2320/matertrans.M2011094

Gupta M, Eugene WWL (2007) Microwaves and metals. Wiley, Hoboken Jayalakshmi S, Gupta S, Sankaranarayanan S et al (2013) Structural and mechanical properties of

Ni60Nb40 amorphous alloy particle reinforced Al-based composites produced by microwave- assisted rapid sintering. Mater Sci Eng A 581:119–127. doi: 10.1016/j.msea.2013.05.072

Jayalakshmi S, Sahu S, Sankaranarayanan S et al (2014) Development of novel Mg–Ni60Nb40 amorphous particle reinforced composites with enhanced hardness and compressive response. Mater Des 53:849–855. doi: 10.1016/j.matdes.2013.07.022

Lee M, Kim JH, Park JS et al (2004) Fabrication of Ni–Nb–Ta metallic glass reinforced Al-based alloy matrix composites by infi ltration casting process. Scr Mater 50:1367–1371

Lee M, Bae D, Kim W, Kim D (2003) Ni-based refractory bulk amorphous alloys with high ther-mal stability. Mater Trans 44:2084–2087

Scudino S, Liu G, Prashanth KG et al (2009) Mechanical properties of Al-based metal matrix composites reinforced with Zr-based glassy particles produced by powder metallurgy. Acta Mater 57:2029–2039. doi: 10.1016/j.actamat.2009.01.010

Scudino S, Ali F, Surreddi K et al (2010) Al-based metal matrix composites reinforced with nano-crystalline Al-Ti-Ni particles. J Phys Conf Ser 240:012154

Zhang Z, Han BQ, Witkin D et al (2006) Synthesis of nanocrystalline aluminum matrix compos-ites reinforced with in situ devitrifi ed Al–Ni–La amorphous particles. Scr Mater 54:869–874. doi: 10.1016/j.scriptamat.2005.11.003

Zheng R, Yang H, Liu T et al (2014) Microstructure and mechanical properties of aluminum alloy matrix composites reinforced with Fe-based metallic glass particles. Mater Des 53:512–518. doi: 10.1016/j.matdes.2013.07.048


http://dx.doi.org/10.1016/j.compositesa.2010.07.004

http://dx.doi.org/10.2320/matertrans.M2011094








Chapter 5 Future Work

Abstract In this chapter, the future research and applicative prospects of amor-phous/BMG reinforced light-metal matrix composites are highlighted.

Keywords Amorphous/metallic glass reinforced LMMCs • Research prospects • Applications

5.1 Future Research and Applicative Prospects

The discussions (in the preceding chapters) on amorphous alloy/metallic glass rein-forced light-metal matrix composites indicate that although the research topic is at its infant stages, there is a lot of potential to be explored in the fi eld. Some of them include modifi cation of processing methods and identifi cation of new methods, investigation of other properties, and preparation of the composites in other forms/products. A multitude of light-metal matrices and reinforcements can be investi-gated considering the different compositions that can be formulated in a single amorphous/glassy system. Another important aspect is the devitrifi cation (inten-tionally devitrifi ed by thermal/stress application) of the amorphous/glassy rein-forcement phase and to study its effect on the properties and interfacial characteristics. The study would tend to explore the atomic diffusion characteristics between the matrix and the reinforcement and the ensuing deformation behavior.

With regard to applications, most of the applications pertaining to conventional MMCs can be substituted by amorphous/glass reinforced composites. Special atten-tion needs to be given to applications wherein wear resistance is essential. While conventional MMCs have high wear resistance, they counter abrade the contacting materials due to their high hardness/brittleness of the ceramic reinforcement. Amorphous/glassy material reinforced composites (hard but metallic in nature) would reduce abrasion of the counterface. Also to be noted is the property of low stiffness and high strength of the metallic glass that provides them with very high resilience, i.e., the ability to store elastic strain energy and release it, a key property for ballistic applications. Also, the stand alone characteristic/advantage of glassy

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material is its high temperature deformation behavior where it can undergo superplastic deformation. This property needs to be utilized, especially in Al-amorphous/glassy reinforced composites, to make components of small dimen-sions (microgears, micromotors, etc.). Prospective applications would also be in the fi eld of corrosion resistance, as most of the conventional MMCs show poor corro-sion resistance. As is known, the corrosion resistance of amorphous/BMG materials is very high due to the absence of grains/grain boundaries in the amorphous struc-ture. How does the advantage of high corrosion resistance of BMG translate to, when used as reinforcement in light-metal matrices should be explored.

5 Future Work

109© S. Jayalakshmi and M. Gupta 2015 S. Jayalakshmi, M. Gupta, Metallic Amorphous Alloy Reinforcements in Light Metal Matrices, SpringerBriefs in Materials, DOI 10.1007/978-3-319-15016-1

A Activation energy , 77 Al , 1 Al-composites , 36

A356 alloy , 34 Al 2080 alloy , 37 Al–AlN , 35 Al–B 4 C , 36 2024, 6060, 5754 Al-matrices with 45 % V f

AlN , 36 Al–Mg–BN , 36 Al–Mg–Si alloy , 38 Al–Si , 38 Al–SiC , 36 Al–Si–Mg–Zn , 38 Al–7Si–10 % V f SiC , 35 Al–TiC , 38

Amorphous/glassy material , 85 amorphous/glassy state-to-nanocrystalline/

crystalline state transformation , 60 arc melting , 69 critical cooling rate ( R c ) , 69 crystallization temperatures ( T x-onset and

T x-peak ) , 62 devitrifi ed structure , 72 diffused halo pattern , 60 driving force , 60 free-energy , 60 free-volume , 60 freezing temperature , 62 glass forming ability (GFA) , 65 high pressure die-casting , 69 medium-range order (MRO) , 73 melt spinning , 65

Newtonian viscous fl ow behaviour , 62 non-equilibrium path , 60 quenching , 70 reduced glass transition temperature , 65 ribbons, wires, powders , 77 squeeze casting , 70 suction casting , 70 super-cooled liquid region , 62 superplasticity , 62 T m melting , 62 undercooling , 65 viscosity , 62 volume , 62

Amorphous reinforcements Fe 72 B 14.4 Si 9.6 Nb 4 glassy alloy , 104 Fe 73 Nb 5 Ge 2 P 10 C 6 B 4 glassy powders , 101 Ni 60 Nb 40 (at.%) amorphous particle

reinforced composite , 89, 95 Ni 39.2 Nb 20.6 Ta 40.2 (wt.%) alloy , 86 Vitraloy 6 , 90 Zr 57 Ti 8 Nb 2.5 Cu 13.9 Ni 11.1 Al 7.5 , 94

Amorphous structure , 94 glass transition temperature ( T g ) , 62 shear transformation zone (STZ) , 74

Aspect ratio , 97 Atomic clusters , 79 Atomic structure , 60, 73

B Ball-milling

amorphization , 71 ball-to-powder ratio , 22, 94,

96, 101

Index

110

Ball-milling (cont.) cold welding, fl attening, fracturing and

re-welding , 22 reactive mixing , 21

C Crystalline , 59

crystal lattice , 59 equilibrium conditions , 60 nearest-neighbors , 73 thermodynamically stable , 60

D Deformation mechanism

grain boundary migration , 13 grain boundary sliding , 13 homogenous , 62 inhomogenous , 62 shear band formation , 62 slip , 13 twinning , 13

Dislocation, high dislocation density , 12

E Energy , 1

F Fracture mechanisms , 33

decohesion of particles , 34 fi ber breakage , 33, 43 fi ber cracking , 33 fi ber pull-out , 34 particle clusters , 34 particle fracture , 34 void formation , 34

Fuel effi ciency , 1

G Glass forming ability

deep eutectic , 66 multi-component system , 66 negative heat of mixing , 66 radius ratio , 66

H Heterogeneous nucleation , 69 Hot extrusion , 94, 96, 101, 105 Hybrid materials , 82, 105

I Induction melting , 69 Infi ltration techniques

infi ltration under pressure , 3, 15, 17, 35, 86

preform , 3, 15, 17, 36, 39, 86 pressure-less infi ltration , 15

Interface chemical reactions , 12 dislocations , 12 intermetallics , 38 local stress concentrators , 12 mechanical bonding , 11 segregation , 16, 27 thermal stresses , 2, 7, 12, 32 wettability , 11, 35, 38, 53

Inter-particle spacing , 97

L Light metal matrix reinforced with

amorphous/glassy materials Al6061 alloy with

[(Fe 0.5 Co 0.5 ) 75 B 20 Si 5 ] 96 Nb 4 glassy particles , 92

Al–6.5Si–0.25Mg (wt.%) alloy (A356 alloy) , 86

A520 with Cu-based Cu 54 Zr 36 Ti 10 (at.%) glassy alloy , 92

Mg–Ni 60 Nb 40 (at.%) amorphous particles reinforced composite , 96

structural dilatation , 97 Light metals , 1

automobile and aerospace , 2 Liquid-state processing , 3, 13, 86

centrifugal casting , 14 disintegrated melt deposition , 18 infi ltration techniques , 15 process parameters , 17 sand casting, permanent mould casting,

pressure die casting/squeeze casting , 14 stir casting , 13 ultrasonic assisted casting , 13

Liquidus temperature , 86 Long-range atomic order , 59, 62

M Master alloy , 35, 69 Matrix , 1–3, 8 Matrix/reinforcement interface , 8, 14, 32 Mechanical alloying, dispersion , 39 Metallic amorphous alloys , 4, 59 Metallic materials , 4, 53, 59

Index

111

Metal matrix composites (MMCs) low ductility , 3, 4, 53, 85 matrix-particle interface/bonding , 3, 18,

22, 36, 38, 53, 85 reinforcements , 2 residual stress fi eld , 10 strengthening mechanisms , 5, 7,

31, 62 Mg , 1

basal slip , 40 casting , 40 wrought , 40

Mg-composites AE44 (Mg–4.0Al–4.1RE–0.3Mn) , 43 Al 2 O 3 short fi ber , 46 AM100 alloy , 46 AZ91/Al 18 B 4 O 33 w , 45 AZ91/SiCw , 45 carbon fi bers , 47 Mg 2 B 2 O 5 whisker-reinforced

AZ91D , 44 Mg-hydroxyapatite , 45 Mg–Li–Al 2 Y , 44 Mg–SiC particles and mullite fi bers , 43 Mg–Zn–Ca/SiCp , 45 TiC/AZ91D , 44 ZC63 alloy , 46 ZK60A/(SiCw+B 4 Cp) , 45 ZK51/SiCw , 44

Microstructure , 2 texture evolution , 40

Microwave sintering, bi-directional hybrid microwave-assisted rapid sintering , 24

MMCs/LMMCs processing , 2

N Nano-composite

Al-7075 alloy matrix , 39 Al-MWCNT , 39 Al–2024 nano-Al 2 O 3 , 38 Al–SiC , 39 Al–3.0 wt.% Mg , 38 AZ31B–1.5 % V f Al 2 O 3 , 47 carbon nanotubes , 47 heat treating , 38 Mg–Al 2 O 3 , 47 Mg–(2–4)Al–1SiC with 2 % nano-SiC , 48

Nanocrystalline aluminum matrix composites , 105

Nanocrystalline reinforcements, Al 70 Ti 20 Ni 10 particles , 105

Non-crystalline , 59, 60 random-network structure , 59

O Ordered structure of atoms , 59 Other processes

accumulative roll bonding , 4, 28, 30 friction stir process , 28 severe plastic deformation , 72

P Phase transitions , 73 Powder metallurgy

ball-milling , 21, 60, 70, 92 blending/mixing , 21 cold-pressing , 86 cryomilling , 23 grinding medium , 70 hot/vacuum hot pressing , 21 mechanical alloying , 24 microwave sintering , 24, 95, 99, 101 rapid sintering , 90, 95 spark plasma sintering , 22, 25, 39

Properties , 2 AES , 36 brittle , 13 CBED , 36 coeffi cient of thermal expansion , 10 compression , 43 cracking and failure , 13 differential scanning calorimeter , 62 ductile , 34 ductility , 8 EDS , 36 elastic modulus , 8 elastic strain limit , 53, 79, 85, 101 electrical resistivity , 97 fi ne grained structure , 18 hardness , 3, 8, 35, 38, 39, 90, 101, 105 high strength , 8 localized plastic deformation , 12 mechanical properties , 7, 9, 10, 22, 23,

25, 34–52, 59, 68, 72–80, 85, 87, 90, 96, 105

porosity , 14, 16, 18, 19, 25–27, 30, 38, 86, 94

SADP , 36 SEM , 36 solidifi cation rate , 17, 18, 21, 67–69 TEM , 36 tensile strength , 8, 35–39, 44, 45, 49 tension-compression yield asymmetry , 40 toughness , 3, 8, 85, 104 XRD , 36, 60, 72, 87, 92, 94, 97, 99 yield strength , 4, 8, 31, 32, 38, 43, 47, 48,

89, 95, 97, 105

Index

112

R Rapid solidifi cation processes , 67 Reinforcements

agglomerate , 14 ceramics , 2 continuous reinforcements and

discontinuous , 8 distribution , 13 fi bers/whiskers/particles , 8 micron-sized , 3 nano-sized , 4 particles , 10 short fi bers , 17 type , 13 whiskers , 9

S Semi-solid processes , 3, 25

compocasting , 27 rheo-processing , 27 thixo-process , 26

Severe plastic deformation, equi-channel angular pressing , 29

Solidifi cation rate , 67 Solid-state processes , 3, 21

powder metallurgy , 21

Splat cooled , 67 Strain hardening , 80 Strain rate , 77 Strengthening mechanisms

enhanced dislocation density , 32 Hall–Petch effect (grain refi nement) , 31 load bearing effect , 32 Orowan strengthening , 32 strength prediction , 32

Structural properties , 53, 77, 85, 90 positron annihilation spectroscopy , 73 TEM/HRTEM , 73 3D atom probe , 73 XAFS , 73

T Thermal properties

endothermic peak , 87 exothermic peak , 87 isothermal , 73

Time–temperature–transformation (TTT) diagram , 65

V Viscoplastic deformation , 77

Index

Documents

Metallic Amorphous Alloy Reinforcements in Light Metal Matrices