Fabrication of fe tib2 nanocomposite with use of high energy milling followed by in situ reaction synthesis and sintering

Fabrication of Fe TiB2 Nanocomposite with use of High Energy Milling Followed by in situ Reaction Synthesis and Sintering by Huynh Xuan Khoa Thesis submitted to the Graduate School, University of Ulsa.

Fabrication of Fe-TiB2 Nanocomposite with use of High-Energy Milling

Followed by in-situ Reaction Synthesis and Sintering

by Huynh Xuan Khoa

Thesis submitted to the Graduate School, University of Ulsan

in partial fulfillment of the requirement for the degree of

Doctor of Philosophy

In Materials Science and Engineering

APPROVED:

Prof Byoung-Kee Kim, Chairman Prof Young-Soon Kwon Prof Jin-Chun Kim PhD Ji-Hoon Yoo Prof Ji-Soon Kim, supervisor

December, 2014 Ulsan, Republic of Korea

Abstract

Metal matrix composites reinforced with nano-particles are very promising materials which is suitable for a large number of applications Fe-based composites reinforced by TiB2 particulates have attracted much attention due to its excellent

mechanical properties as well as low coefficient of thermal expansion In-situ formation results in the clean particle-matrix interfaces with higher interfacial strength, finer reinforcement size and better particle-size distribution Hence, the in-situ technique is the optimal choice for the synthesis of nanocomposite

In this study, Fe-TiB2 nanocomposite was in-situ fabricated from titanium

hydride (TiH2) and iron boride (FeB) powders by high-energy ball milling and

subsequent heat-treatment High-energy ball milling was chosen for mechanical activation as an effective method to achieve the desired results subsequently The

consumption during milling and used for discussion on the powder characteristics and the subsequent reaction behavior About 20% of the input energy were transferred into the material at the milling speed of 500 rpm and 33% at 700 rpm

By increasing the milling energy, distribution of starting powders was gradually homogeneous and reduced their size to a nanoscale Moreover, the thermal behaviors such as decomposition of TiH2 and the formation reaction of TiB2 from

Ti and FeB were lowered Obviously, Fe-TiB2 nanocomposite powders after

reaction synthesis showed more homogeneous microstructure for powder mixture milled with higher specific energy Microstructure was characterized by smaller 5

nm TiB2 particulate homogeneous distributed in Fe matrix

Understanding the reaction mechanism helps controlling the affecting factors to achieve the best results Phase change was analyzed by X-ray diffraction and phase distribution was observed by electron microscopy during reaction synthesis of powder mixture milled with various milling conditions in order to explore the formation mechanism of TiB2 particles in Fe matrix The result

indicated that titanium reacts with boron at the interface of Ti and FeB by gradual diffusion reaction, forming TiB2 particles, reducing the amount of boron and

induced phase transition of FeB to Fe2B The process ended when whole the Ti

phase transfered into TiB2 phase and Fe matrix formed from the position of Fe2B

left The reaction rate strongly depended on the size and distribution of FeB particles With the finer FeB, the more homogeneous microstructure of Fe-TiB2

composite powder formed, involving nano TiB2 particles distributed in the Fe

matrix

Most refractory reinforced - metal composite are used for wear resistance parts and cutting tools, so sintering is always next stage of the manufacturing process of materials A part of this study intended to examine the consolidation of nanocomposite The sintering process was performed by both pressureless (PLS) and pressure (SPS) sintering techniques The main effecting factors of sintering time and temperature were investigated The result showed that microstructure and properties of the composites strongly affected by sintering time and temperature With the dominant advantage of low sintering temperature and short sintering time, the SPSed-samples retained nano-size TiB2 particles and obtained very high

density and hardness The PLSed-samples showed sub-micrometer TiB2 particles,

but the hardness obtained also high, equivalent to some WC-Co systems

Acknowledgements

It is my great pleasure that I am in a position to express my deep gratitude

to some persons without their help the present research work could have not taken final shape Foremost, I express my cordial gratitude, thanks and regards to

my supervisor Professor Ji-Soon Kim I am extremely grateful for his sincere help, valuable suggestions and constant encouragement and unique guidance during the four years of my Ph.D course in the University of Ulsan

I would like to express my sincere gratitude to Korean students, Mr G P Ahn, B H Lee, Y H Lee, Sang-W Bae, Sun-W Bae, W J Kim and J Y Joe for their many valuable suggestions, encouragement and assistance during my research work

I am expressing my gratitude to the Korea Institute of Ceramic Engineering and Technology for helping in measurement of mechanical properties A special thanks to Mr Sang-Ha Park of Deagu Machinery Institute of Components & Materials for the use of their equipment and advice I am also thankful to all staffs

of UOU Research Facilities Center for their enthusiasm in the investigation of my specimens

I would like to thank my committee members: Prof Ji-Soon Kim (Advisor), Prof Young-Soon Kwon, Prof Byoung-Kee Kim, Prof Jin-Chun Kim and PhD Ji-Hoon Yoo for their support and advice in developing this document

Finally, I am grateful to my wife and my daughter for their constant inspiration to carry out the research work

Date:

University of Ulsan,

Ulsan city, Republish of Korea Huynh Xuan Khoa

Table of Contents Abstract .I Acknowledgements III List of Figures IV List of Tables VII

Chapter 1 Introduction 1

Reference 5

Chapter 2 Theoretical Background 7

2.1 Metal matrix nano-composites (MMnCs) 7

2.1.1 Processing techniques for metal matrix nano-composites (MMnCs) 9

2.1.2 Strengthening in particulate reinforced metal matrix composites 17

2.1.3 Previous works on production process of Fe-TiB2 composite 19

2.2 Mechanical activation by high-energy ball milling 26

2.2.1 Mechanical activation 26

2.2.2 High-energy milling equipments 27

2.2.3 Planetary high-energy ball milling and processing variables 27

2.2.4 Energetic of mechanical activation process 30

2.2.5 Effect of mechanical activation on properties of solid 38

2.3 Synthesis reaction mechanisms and reactions in the solid state 40

2.4 Sintering process 46

2.4.1 Presureless sintering 46

2.4.2 Spark plasma sintering – Outstanding method for sintering MMnCs 47

Reference 54

Chapter 3 Experimental procedure 58

3.1 Materials 58

3.2 High-energy ball milling Process 59

3.3 Heat-treatment process 62

3.4 Sintering of nanocomposite powder 63

3.4.1 Pressureless sintering 63

3.4.2 Spark plasma sintering 65

3.5 Characterization 65

3.5.1 Particle size analysis 67

3.5.2 XRD analysis 67

3.5.3 SEM and TEM analysis 67

3.5.5 Density and hardness measurement 68

3.5.6 Wear test, and transverse rupture strength test (TRS) 69

Reference 71

Chapter 4 Energetics of high-energy ball milling process 72

4.1 Indirect approach – Calculation of milling energy by collision model 73

4.1.1 Calculation of milling energy by collision model 73

4.1.2 Measuring the power consumption and comparison with measurement 79

4.2 Direct measurement of total energy during high-energy ball milling of FeB and TiH2 powder mixture 81

4.3 Summary 82

Reference 84

Chapter 5 High-energy ball milling process of initial FeB-TiH2 powder mixture 85

5.1 The state of the powder mixture during High-energy ball milling 85

5.2 Effect of milling energy on mixing homogeneity and size of powder 85

5.3 Effect of Milling Energy on Reaction Behavior of Powder mixture 90

5.4 Sumamry 92

Reference 93

Chapter 6 Fabrication of Fe-40 wt% TiB2 nanocomposite powder from FeB and TiH2 powders - Powder synthesis and formation behavior of TiB2 particulates in Fe-matrix during reaction synthesis 94

6.1 Reaction synthesis of milled powders by heat treatment 94

6.1.1 Shape and Particle Size of Fe- TiB2 nanocomposite powder 94

6.1.2 Phase analysis of Fe- TiB2 nanocomposite powder 94

6.1.3 Microstructure of Fe- TiB2 nanocomposite powder 98

6.2 Formation behavior of TiB2 particulates in the Fe - matrix during reaction synthesis 101

6.2.1 Phase change during reaction synthesis 101

6.2.2 Composite Microstructure and Analysis 103

6.2.3 Discussion 108

6.3 Summary 112

Reference 113

Chapter 7 Combination of Synthesis and Sintering for Consolidation of Fe-TiB2 Nanocomposite from FeB and TiH2 114

7.1 Sintering Behaviors (sintering conditions vs shrinking) 114

7.2 Phases change during sintering 119

7.3 Microstructure evolution during sintering 121

7.4 Relation of density and hardness 125

7.5 Transverse rupture strength (TRS) and abrasive wear of the Fe-TiB2 composites 127

7.5.1 Transverse rupture strength 127

7.5.2 Abrasive wear 130

7.6 Summary 138

Reference 138

Chapter 8 Conclusions 140

8.1 Energetics of high-energy ball milling 140

8.2 High-energy ball milling process of initial FeB-TiH2 powder mixture 141

8.3 Powder synthesis and formation behavior of TiB2 particulates in Fe-matrix during reaction synthesis 142

8.4 Combination of Synthesis and Sintering for Consolidation of Fe-TiB2 Nanocomposite from FeB and TiH2 143

List of Figures Fig 2.1 Scheme of planetary disk with movement in a counter direction; Rp and Wp

- revolution radius and speed, Rv and Wv - rotation radius and speed, Hv -

pot height 29

Fig 2.2 A snapshot of the ball motion inside the planetary ball mill with set of forces on a ball from video introducing Pulverisette 6 classic line – FRITSCH GmbH 29

Fig 2.3 High-energy ball milling parameters of a planetary mill and cylindrical vial 30

Fig 2.4 The various possible energy dissipation processes that can occur when mechanical energy is input into a solid 31

Fig 2.5 Energy dissipated per hit versus the rotation speed of the planetary ball mill (Fritsch “Pulverisette 5” (left) and AGO-2 (right)) 33

Fig 2.6 Milling map showing the Pe-nv relationship with typical planetary mill rotation speeds 35

Fig 2.7 The milling map of MoSi2 at different conditions 36

Fig 2.8 Scheme of the planetary mill using for the electrical power measurements 37

Fig 2.9 Electrical power absorption measured vs the rotation speed of the mill for the Fe-Zr powder system (a), and experimental (full triangles) and calculated (full lines) power absorption for the Fe-Zr (99balls, db = 10, 40g powder, BPR:10) and the Ti-Al (70 balls db= 8 + 13 balls db = 10, 20g powder, BPR: 9.3) experiments(b) 38

Fig 2.10 Two solids will react only when in close proximity 43

Fig 2.11 Formation of phase boundary 44

Fig 2.12 A “hopping” mechanism in the reaction A + B → AB 44

Fig 2.13 SPS system configuration 53

Fig 2.14 DC pulse current flow through the particles 53

Fig 3.1 Morphology of starting powders; a –TiH2, b - FeB powders and c - powder mixture of FeB-TiH2 (2:1.1) 60

Fig 3.2 Structure of planetary high energy ball mill (AGO-2) 61

Fig 3.3 Schematic illustration of the AGO-2 mill with the electric power meter (3) 61

Fig 3.4 General view of the tube furnace (up), alumina boat was tied by W wire (middle) and illustration of alumina boats in the furnace (below) 64

Fig 3.5 General view of SPS apparatus (up) and Parameters of graphite die and punches (below) 66

Fig 3.6 (a) Schematic of pin-on-disk wear test system according to the standardized method ASTM G99 – 95a F is the normal force on the pin, d is the pin or ball diameter, D is the disk diameter, R is the wear track radius, and w is the rotation velocity of the disk (b) Wear tester and (c) pin holder (H) 69

Fig 3.7 Illustration for TRS test 70

Fig 4.1 Structure modeling of an AGO-2 planetary mill 74

Fig 4.2 Hindering factor vs nv = Nb/Nb,tot for different ball diameters 76

Fig 4.3 Kinetic energy released per hit as a function of the rotation speed of the mill and the mass of the ball 77

Fig 4.4 The results of powder calculation for WC ball 5mm 78

Fig 4.5 Electrical power absorption measured vs the rotation, nv = 0.1 80

Fig 4.6 Electrical power absorption measured vs the rotation, nv = 0.2 80

Fig 4.7 Electrical power absorption calculated (full lines with triangles and circles) and measured (full triangles and circles) vs the rotation speed of the mill for the FeB, TiH2 powder system 81

Fig 4.8 Change in milling energy with rotational speed and milling time 83

Fig 5.1 XRD pattern of the FeB-TiH2 (2:1.1) powder mixture after high-energy ball milling under various milling conditions 86

Fig 5.2 SEM-SEI images of FeB-TiH2 powder mixtures milled at different milling energy levels 87

Fig 5.3 Morphology and the mapping images of Fe and Ti for various milling time 88

Fig 5.4 Results of particle size analysis of starting powder mixture and the powders milled at 700rpm/15, 60 & 180min 89

Fig 5.5 SEM-BSE images show cross-section of as-milled powders (from a to f) and TEM image (g) shows the distribution of powder mixture of (f) powder 91

Fig 5.6 DSC curves of FeB-TiH2 (2:1.1) powder mixtures milled under various conditions 92

Fig 6.1 SEM-SEI images of the Fe-TiB2 composite powder after synthesis 96

Fig 6.2 XRD-patterns of the Fe-TiB2 composite powder after synthesis 97

Fig 6.3 SEM-BSE images show the microstructure of Fe-TiB2 composite powder after synthesis 99

Fig 6.4 TEM images show microstructure of the composite powder (a) and TiB2 size and shape (b) 100

Fig 6.5 XRD-pattern of the FeB-TiH2 powder mixture after milling 700rpm/180min 101

Fig 6.6 XRD-pattern of the FeB-TiH2 powder mixture after milling 700rpm/60min 102

Fig 6.7 XRD-patterns of the FeB-TiH2 powder mixture after milling 500rpm/120min 102

Fig 6.8 SEM-BSE images show microstructure of Fe-TiB2 composite powder after milling 500rpm/120min and heat treatment at 900oC/ 00 to 120min 104

Fig 6.9 EDS-line scanning on Fig 6.8 105

Fig 6.10 Microstructure of Fe-TiB2 composite powder after milling 700r/60min 106

Fig 6.11 SEM-BSE images show microstructure of Fe-TiB2 composite powder after milling 700rpm/180min and heat treatment 750oC/ 00, 05, 15 & 30min 107

Fig 6.12 DSC curves of the as-milled powders 110

Fig 6.13 Illustrating the formation of microstructure of the Fe-TiB2 nanocomposite 111

Fig 7.1 Shrinkage of samples sintered at the various sintering condition by PLS technique 115

Fig 7.2 Shrinkage of sample sintered at 11500C was recorded continuously by SPS 116

Fig 7.3 The densification rate of sample sintered at 11500C vs relative density 116

Fig 7.4 Broken surfaces of sample sintered at before and after densification rate peaks 117

Fig 7.5 Broken surfaces of sample enlarged from the square in Fig 7.4 118

Fig 7.6 XRD pattern of some samples sintering by PLS 120

Fig 7.7 XRD pattern of some samples sintering by SPS 120

Fig 7.8 Microstructure evolution during sintering at 1350oC by PLS (left: x 3,000, right: x 10,000) 122

Fig 7.9 Microstructure evolution during sintering at 1400oC by PLS (left: x 3,000, right: x 10,000) 123

Fig 7.10 Microstructure of some samples sintering at various sintering condition by SPS 124

Fig 7.11 Density and Hardness of some sample sintered by PLS 125

Fig 7.12 Density and Hardness of some sample sintered at various sintering condition by SPS 126

Fig 7.13 Displacement - load relation of Fe – 40wt% TiB2 composite fabricated by PLS (left) and SPS (right) during TRS test (5 specimens: 1, 2, 3, 4, 5) 128

Fig 7.14 Fractured surfaces of Fe – 40wt% TiB2 composite fabricated by PLS (above) and SPS (below) after TRS test 129

Fig 7.15 Variation of volume loss with sliding distance for all specimens 131

Fig 7.16 Volume loss for all specimens tested under an applied load of 18N and 23N 132

Fig 7.17 Wear rate (volume loss per unit sliding distance) for all specimens tested under an applied load of 18N and 23N and a sliding velocity of 1.4 m/s 132

Fig 7.18 Wear rate versus hardness for all specimens tested under an 134

Fig 7.19 SEM-SEI micrograph showing the worn surface of all specimens tested 136

List of Tables

Table 2.1 Phase transformation (single phase) 43

Table 2.2 Types of solid state heterogeneous reactions 43

Table 3.1 Chemical composition of FeB powder 59

Table 3.2 Particle size analysis of starting powders (volume distribution) 60

Table 4.1 Parameters of the AGO-2 and relation between them 74

Table 4.2 Summary of calculating character factor for AGO-2 mill 75

Table 4.3 Number of ball at different degree of filling 76

Table 4.4 Summary of calculating relative impact velocity and kinetic energy of one ball 77

Table 4.5 The energy transferred to the material for different milling time 83

Table 5.1 EDS analysis on cross-section surface 92

Table 6.1 Particle size analysis of composite powders (volume distribution) 97

Table 6.2 EDS analysis of cross-section surface 100

Table 6.3 EDS analysis of cross-section surface 105

Table 6.4 Thermal properties of as-milled powders with different activation time by DSC 110

Table 7.1 The list of sintering conditions vs density and hardness 126

Table 7.2 TRS tested results of Fe – 40wt% TiB2 composite fabricated by PLS 128

Table 7.3 TRS tested results of Fe – 40wt% TiB2 composite fabricated by SPS 129

Table 7.4 Data of wear test for Fe - 40wt% TiB2 and the comparative materials 130

Table 7.5 Summarizing properties of Fe-40wt% TiB2 composite and comparing with WC-Co system 137

Chapter 1

Introduction

TiB2 is one of the excellent reinforcements for steel due to its high melting point

(2980°C), high hardness (3400 kg/mm2), high elastic modulus, good corrosion resistance,

and chemical inertness [1] Hence, Fe-TiB2 composite has attracted much attention in

many applications to tools, dies and wear-resistant parts for last several decades

In situ fabrication technique is known to be the most suitable for composite materials with fine and homogeneous distribution of particulate reinforcements It includes usually the synthesis of dispersoid phase within a matrix by a chemical reaction between the constituent elements of starting materials It has advantages like fine particulate size of reinforcement, homogeneous distribution in the matrix and clean particulate-matrix interface with improved wettability and high interfacial strength (improved mechanical properties) [2, 3]

Synthesis methods of TiB2 can be summarized as follows: laser cladding [4, 5],

plasma transferred arc (PTA) [6, 7], aluminothermic reduction [8], and self-propagating high temperature synthesis (SHS) [9, 10] According to our search, a research on fabrication of Fe-TiB2 composite from FeB and TiH2 starting powders has not been

reported yet FeB and TiH2 powder materials have advantage in cost and availability, and

additionally, they are both brittle and therefore suitable to exploit the advantage of energy ball milling process such as reducing the particle size, homogeneous mixing of constituents and mechanical activation for further reaction, etc [11]

high-Mechanical activation by high-energy ball milling is defined as an increase in reaction ability due to stable changes in a solid structure It is based on the application of mechanical impact upon the treated material that initiates a number of physical and chemical effects in it During milling process besides particle size reduction a number of different types of structural and electronic defects arise, internal energy and specific surface area increase, the materials become more chemically reactive; the energy stored

in the milled materials ensures that the subsequent solid state reaction occurs under more favorable energetic condition The improvement can be the decrease of the reaction temperature and associated enthalpy value [12-14], creating nanostructured powders [15], obtaining a higher relative density [16], etc

Planetary ball mill is known a representative high-energy ball mill used most frequently in research laboratories Milling parameters like rotation speed, milling time, ball-to-powder ratio, ball media, etc can be varied to establish various milling conditions which affect powder characteristics and subsequent processes Therefore, it is important

to know the relationship between processing variables and milling energy, and thereby also the behavior of energy transfer to powder during milling Milling energy can be calculated by the mathematical model [17, 18] or direct measurement [19, 20] In this study the electrical energy consumption during the milling process was measured as suggested by Magini et al [21], where the energy transferred to the material was studied

by measuring the energy input into the mill with empty and filled vials (all the other conditions remained the same) using a high precision electrical power meter

This work aims at investigating the feasibility of fabricating nanoscale TiB2

-reinforced Fe matrix composite by a combination of high-energy ball milling and subsequent heat treatment, using titanium hydride (TiH2) and ferroboron (FeB) powders

as starting materials We expected that this process can provide us with a novel way for rapid, simple and cost-effective in situ synthesis method

Motivation:

It has been not found any publication related to fabrication of Fe-TiB2 composite

from FeB and TiH2 initial powders The FeB and TiH2 materials are both cheap and

available The FeB and TiH2 materials are both brittle, so it is easy to reduce their

particle size by milling TiB2 is in situ created in Fe matrix, so the composite is improved

wettability and better particle-size distribution Fabrication process is simple including mechanical activation (MA) and heat treatment (HT)

Purpose of this study:

The present work aimed at investigating the feasibility of fabricating TiB2

-reinforced Fe matrix composite by a combination of mechanical activation (MA) and heat treatment (HT) using titanium hydride (TiH2) and ferroboron (FeB) as initial materials It is expected that the route is able to provide a novel process for rapid, simple and cost-effective synthesis method

Objectives:

Fabricating Fe - TiB2 (40 wt%) matrix composite in two form of powder and bulk

The composite in the form of bulk is utilized for cutting tools; composite powder is used

or enhancing low wear resistance such as Stainless Steel and High Strength Low Alloy Steel (HSLA-Steel) that have been popularly utilized in Powder Metallurgy We have expected that the complete reaction between FeB and TiH2 powders to form a Fe-TiB2

end product without intermediate phase of Fe-Ti-B Fe - TiB2 composite gains a

homogeneous microstructure TiB2 particle is fine and homogeneous distributed in Fe

matrix The interface of TiB2 and Fe matrix is clean Sintered composite obtains high

density and hardness, high wear resistance and transverse rupture strength

The rest of the thesis is organized as follows:

Chapter 2 reviews background which most directly concerned and important for this study An overview of Particulate Reinforced Metal Matrix Composites was performed on processing technique, strengthening mechanism and previous works on production process of Fe-TiB2 composite was summarized The concept of mechanical

activation, reaction synthesis and sintering process are also introduced in this chapter

Chapter 3 describes experimental procedures used for composite fabrication, the various characterization methods used for confirming the phase, observing the morphology, microstructure, thermal behavior The test of mechanical properties also presents in this section Chapter 4 is intended for enneretics of high-energy ball milling process Chapter 5 describes the mechanically activated process of the initial FeB-TiH2

powder mixture Chapter 6 presents the fabrication of Fe-40 wt% TiB2 nanocomposite

powder from FeB and TiH2 powders The results of powder synthesis by high-energy ball milling and heat treatment was presented and discussed This chapter also discusses the formation behavior of TiB2 particle in Fe-matrix during reaction synthesis The

thermodynamic consideration of Fe-Ti-B system, effect of milling energy on reaction rate and the formation mechanism of TiB2 and typical structure of nanocomposite are

respectively presented and discussed In Chapter 7, the combination of synthesis and sintering for consolidation of Fe-TiB2 composite will be summarized in detail The

investigation of mechanical properties of Fe-TiB2 composite bulks such as density,

hardness, wear resistance and transverse rupture strength (TRS) are major results to be presented and discussed in this chapter Finally, conclusions drawn from this study and the future directions of this study are summarized in Chapter 8

[3] Casati, R and M Vedani, Metal Matrix Composites Reinforced by

Nano-Particles—A Review Metals, 2014 4(1): p 65-83

[4] Du, B., et al., Laser cladding of in situ TiB2/Fe composite coating on steel Applied

Surface Science, 2008 254(20): p 6489-6494

[5] Du, B., et al., In situ synthesis of TiB2/Fe composite coating by laser cladding

Materials Letters, 2008 62(4–5): p 689-691

[6] Darabara, M., G.D Papadimitriou, and L Bourithis, Production of Fe–B–TiB2

metal matrix composites on steel surface Surface and Coatings Technology,

2006 201(6): p 3518-3523

[7] Xibao, W., W Xiaofeng, and S Zhongquan, The composite Fe–Ti–B–C Coatings

by PTA powder surfacing process Surface and Coatings Technology, 2005 192(2–3): p 257-262

[8] Anal, A., T.K Bandyopadhyay, and K Das, Synthesis and characterization of

TiB2-reinforced iron-based composites Journal of Materials Processing Technology, 2006 172(1): p 70-76

[9] Degnan, C.C and P.H Shipway, The incorporation of self-propagating,

high-temperature synthesis-formed Fe-TiB2 into ferrous melts Metallurgical and Materials Transactions A, 2002 33(9): p 2973-2983

[10] Lepakova, O.K., L.G Raskolenko, and Y.M Maksimov, The mechanism of phase

and structure formation of the Ti-B-Fe system in a combustion wave Combustion, Explosion and Shock Waves, 2000 36(5): p 575-581

[11] Baláž, P., High-energy milling, in Mechanochemistry in Nanoscience and

Minerals Engineering 2008, Springer Berlin Heidelberg p 103-132

[12] Berbenni, V., A Marini, and G Bruni, Effect of mechanical activation on the

preparation of SrTiO3 and Sr2TiO4 ceramics from the solid state system SrCO3–TiO2 Journal of Alloys and Compounds, 2001 329(1–2): p 230-238

[13] Pourghahramani, P and E Forssberg, Effects of mechanical activation on the

reduction behavior of hematite concentrate International Journal of Mineral Processing, 2007 82(2): p 96-105

[14] Singh, S., et al., Effect of mechanical activation on synthesis of ultrafine Si3N4–

MoSi2 in situ composites Materials Science and Engineering: A, 2004 382(1–2):

p 321-327

[15] Gras, C., et al., Mechanical activation effect on the self-sustaining combustion

reaction in the Mo–Si system Journal of Alloys and Compounds, 2001 314(1–2):

p 240-250

[16] Lee, S.E., et al., Effects of mechanical activation on the sintering and dielectric

properties of oxide-derived PZT Acta Materialia, 1999 47(9): p 2633-2639 [17] Burgio, N., et al., Mechanical alloying of the Fe−Zr system Correlation between

input energy and end products Il Nuovo Cimento D, 1991 13(4): p 459-476

[18] Murty, B.S., M Mohan Rao, and S Ranganathan, Milling maps and amorphization

during mechanical alloying Acta Metallurgica et Materialia, 1995 43(6): p 2443-2450

[19] Iasonna, A and M Magini, Power measurements during mechanical milling An

experimental way to investigate the energy transfer phenomena Acta Materialia,

1996 44(3): p 1109-1117

[20] Magini, M., et al., Power measurements during mechanical milling—II The case of

“single path cumulative” solid state reaction Acta Materialia, 1998 46(8): p 2841-2850

[21] M., M and I A., Energy Transfer in Mechanical Alloying Materials Transactions,

JIM, 1995 36(2): p 123-133

Chapter 2

Theoretical Background

2.1 Metal matrix nano-composites (MMnCs)

Metal matrix composites (MMCs) are essentially metallic materials reinforced with a second phase Even though any material can practically be used as the second phase, ceramic oxides, nitrides and carbides are most common This class of material attempts to combine the ductility and toughness of the metal matrix with higher modulus, wear resistance, and thermal stability of the ceramics to achieve attractive features superior to the unreinforced metal

Metal matrix composites (MMCs) reinforced with nano-particles, also called Metal Matrix nano-Composites (MMnCs), are being investigated worldwide in recent years, owing to their promising properties suitable for a large number of functional and structural applications The reduced size of the reinforcement phase down to the nano-scale is such that the interaction of particles with dislocations becomes of significant importance and, when added to other strengthening effects typically found in conventional MMCs, results in a remarkable improvement of mechanical properties

The main issue to be faced in the production of MMnCs is the low wettability of ceramic nano-particles with the molten metal matrix, which do not allow the production

of MMnCs by conventional casting processes Small powder aggregates are in fact prone

to form clusters, losing their capability to be homogeneously dispersed throughout the matrix for an optimal exploitation of the strengthening potential For this reason, several alternative methods have been proposed in order to overcome this problem

The production methods can be categorized into two major groups: ex situ and in situ The first synthesis route consists of adding nano-reinforcements to the liquid or powdered metal, while in situ processes refer to those methods leading to the generation

of ceramic nano-compounds by reaction during processing, for example by using reactive gases Several methods have been developed for ex situ synthesis of MMnCs In particular, different powder metallurgy techniques were successfully employed Moreover, ultrasound-assisted casting plays a particularly promising role for its high potential productivity Alternative methods are listed and discussed in a following section The methods used for the characterization of MMNCs are the same of those used for conventional MMCs and alloys Of course, the downsizing of the reinforcement implies the use of higher resolution techniques for characterization of morphology and local chemistry of the constituents

In the literature, different kinds of matrix metals have been coupled with several types of nanometric phases Ceramic compounds (SiC, Al2O3, etc.), intermetallic

materials and carbon allotropes were used to reinforce Al, Mg, Cu and other metals and alloys Particular importance is assigned to carbon nanotubes (CNT), which are characterized by very high strength, stiffness and electrical conductivity These properties confer higher mechanical strength while improving electrical and thermal properties of the base material Moreover, MMnCs revealed to be able to improve other interesting engineering properties, such as damping capacity, wear resistance and creep behavior A review of fabrication methods for MMnCs has been covered by Casati, R and M Vedani [1]

2.1.1 Processing techniques for metal matrix nano-composites (MMnCs)

The desire to manufacture composites with reproducible microstructures and predictable properties has led to the development of a variety of MMCs processing techniques Reviews of fabrication methods for MMCs have been covered by Ibrahim [2] and Tjong [3] Important considerations in materials selection are the characteristics of the matrix, the reinforcement, and the interface formed between them Although the final use of the composite dictates the choice of matrix, density, ductility, thermal conductivity, and fracture toughness are generally considered For the reinforcement, thermal stability, tensile strength, coefficient of thermal expansion, chemical compatibility with the matrix, modulus, density, and cost are some of the significant considerations The metal-reinforcement interface is perhaps the most difficult to control Ideally, the two materials must be chemically compatible for a clean and a strongly bonded interface Since ceramic particles are only partially wetted by molten metals, surface treatments must be used to enhance bonding of the two dissimilar materials Coating of ceramic powders with thin metallic films has been used successfully in many instances to improve bonding between matrix and reinforcement materials

Processing of particulate reinforced composites is performed by a variety of techniques, but can generally be grouped into solid state and liquid metal processing Powder metallurgy is the most common solid state processing technique It involves blending metallic and reinforcement powders according to required weight fractions and then consolidating the blended powders into semi-finished or finished products Usually the matrix alloy is in pre-alloyed form rather than elemental powders Some initial pre-blending steps such as sieving and de-agglomeration of powders are usually essential to obtain a uniform distribution of powders Heat treatment of some metallic powders may, however, be necessary to reduce oxide layers metallic surfaces Due to differences in the

densities of the metal and ceramic materials, segregation of the powders can occur during blending High energy ball milling also known mechanical alloying originally developed for oxide dispersion strengthening of nickel based alloys, may be employed at the blending stage to enhance dispersion of fine reinforcement particles in the metal matrix Mechanical alloying also enhances mechanical bonding between the matrix and reinforcement Consolidation procedures for blended composite powders include uniaxial pressing followed by sintering (effective for small samples), cold isostatic pressing followed by sintering, hot isostatic pressing and, hot extrusion In liquid state processing, reinforcing particulates are directly incorporated in the molten metal Examples of liquid state processing techniques are melting infiltration, slurry casting, and spray deposition

For the large-scale production of metal matrix nanocomposites, the main problem

to face is the low wettability of ceramic nano-particles, which does not allow the preparation of MMnCs by conventional casting processes since the result would be an inhomogeneous distribution of particles within the matrix The high surface energy leads

to the formation of clusters of nanoparticles, which are not effective in hindering the movement of dislocations and can hardly generate a physical-chemical bond to the matrix, thus reducing significantly the strengthening capability of nanoparticles Several unconventional production methods have been studied by researchers in order to overcome the wettability issue, either by formation of the reinforcement by in situ reaction or by the ex situ addition of the ceramic reinforcement by specific techniques Hereafter, the most studied and successful methods are described by classifying them into liquid, semisolid and solid processes

Liquid Processes

For composites prepared by the conventional liquid metallurgy route, severe aggregation of nanoparticles frequently occurs even when mechanical stirring is applied before casting This is due to poor wettability and high viscosity generated in the molten metal owing to high surface-to-volume ratio of the nano-sized ceramic particles The density of nanoparticles does not play an important role in the production process of nanocomposites Such small particles are supposed to float on the top of the molten bath, even if their density is relatively higher than that of the liquid matrix This issue was indeed of paramount importance in micron-sized particle reinforced composites, but it is felt that in nano-reinforced materials, other effects such as those induced by extensive surface tension play a much more important role

High-intensity ultrasonic waves revealed to be useful in this context since they produce acoustic transient cavitation effects, which lead to collapsing of micro-bubbles The transient cavitation would thus produce an implosive impact, strong enough to break the nanoparticle clusters and to uniformly disperse them in the liquid metal According to this technique, a good dispersion of 2% vol of SiC nano-particles (d < 30 nm) in aluminum alloy 356 was achieved An improvement of 20% in hardness over the unreinforced alloy was achieved Lan et al [4] produced nano-sized SiC/AZ91D Mg alloy composites through the same method A fairly good dispersion of the particles was achieved despite some small clusters still existed into the matrix Owing to general improvement of the dispersion, the 5 wt.% SiC reinforced composite led to a microhardness increase of 75% For A356 alloy based nano-composites produced by ultrasonic assisted casting, γ-Al2O3 revealed a better nucleation catalyzer than α-Al2O3

probably due to its lower lattice mismatch with the metal matrix Other tests were also conducted in the same research with TiC and SiC of different sizes Tensile tests performed on AZ91D alloy and on the same material reinforced by 1 wt.% of nano-AlN

produced by ultrasound-assisted casting, revealed an increase of yield strength in MMnCs at room temperature of 44% and of 21% at 200 °C when compared to the unreinforced AZ91D alloy For the same materials, a decrease of fracture strain at room temperature was achieved while an enhanced ductility was measured at 200 °C Improved ductility was detected by Wang et al [5] even at room temperature The yield strength (YS), ultimate tensile strength (UTS) and fracture elongation of an AZ91 alloy were 104 MPa, 174 MPa and 3.6%, respectively, whereas the corresponding values for the AZ91 alloy reinforced by 0.5 wt.% of 50 nm SiC were: 124 MPa, 216 MPa and 6.6%, respectively

The addition of 1.5 wt.% SiC to Mg–4Zn alloy obtained by an ultrasonic cavitation-based solidification process led to an increase of RT ductility of more than twice as well as to improved YS and UTS A reduction of grain size was also observed

by the same authors in reinforced sample (150 μm vs 60 μm), which increased the castability of the alloy This behavior was supposed to be related to an improved casting quality, since the resulting finer grain size of the composite can improve melt feeding characteristic minimizing porosity, shrinkage and enhancing hot-tearing resistance

In situ MMnCs have been successfully prepared by liquid metallurgy processes [6-8] 50 nm-TiB2-reinforced copper-matrix composites were produced by adding B2O3,

C and Ti in a Cu–Ti melt The composites exhibited significantly improved mechanical properties In particular, the YS of Cu and Cu/TiB2 was 298.7 MPa and 509.6 MPa,

respectively Al/TiB2 nanocomposites were also synthesized by an in situ method, by adding a mixture of potassium hexafluorotitanate (K2TiF6) and potassium

tetrafluoroborate (KBF4) salts in an Al melt under argon atmosphere

Disintegrated Melt Deposition (DMD) is a further liquid metallurgy process successfully employed for nano-composite production Alumina nanoparticles have been

well dispersed in Al–Mg alloys by heating the metal in argon atmosphere and adding the ceramic particles by means of a vibratory feeder The melt was stirred and poured, then disintegrated with argon gas jets and deposited onto a metallic substrate Finally, the MMnCs were extruded to reduce porosity down to very low levels and to achieve a good dispersion of the particles

Selective laser melting (LSM) was also used to produce Ti-based composites reinforced by nanoparticles Powders were milled by high-energy ball milling and then melted by laser beam under protective atmosphere Through this method, a unique microstructure very different from the initial microstructure of the reinforcement was achieved A proper decrease in volumetric energy density led to the development of TIC whisker and of uniformly dispersed nano-lamellar TiC starting from dendritic TiC The same research confirmed that well dispersed nano-particles induce improved mechanical and wear properties to the Ti matrix

Melt stirring, high-pressure die casting [9] and arc-discharge plasma [10] method were also used to produce AZ91/CNT composites and in situ Al/AlN MMnCs, respectively Finally, it was highlighted that the main problem to be faced in the production of CNT-MMnCs by the liquid metallurgy method is the interaction of the nanotubes with the liquid metal In fact, the process may cause damage to CNTs or formation of chemical reaction products at the CNT/metal interface Therefore, this synthesis route is mainly indicated for composite matrices having low-melting temperatures and reduced reactivity with the reinforcement phase The problem of low wettability of CNTs can be partially overcome by coating CNT with metal layers (for example Ni) The field of surface modification appears as quite promising and it is open

to innovation for attenuating the drawbacks on wettability and the tendency to clustering

of nanoparticles

Semi-Solid Processes

Only few works are available in literature about this topic even if this method has been widely applied for micrometer-size particle-reinforced MMCs, and it would be extremely interesting for large-scale production [11, 12] A356/Al2O3 MMnCs were

produced by using a combination of rheocasting and squeeze casting techniques Rheocasting is a semi-solid phase process, which has several advantages: it is performed

at lower temperatures than those conventionally employed in foundry practice, resulting

in reduced thermochemical degradation of the reinforcement surface Moreover, the material shows thixotropic behavior typical of stir cast alloys and production can be performed by conventional foundry methods During rheocasting, the pre-heated nanoparticles are added in the semi-solid slurry while it is vigorously agitated in order to achieve a homogenous particle distribution Then the slurry is squeezed using a hydraulic press Mg alloy AZ91 ingots reinforced by nano-SiC particles were produced by semisolid stirring-assisted ultrasonic vibration After homogenization treatment and extrusion, the SiC reinforcement featured a fairly good dispersion, although bands of accumulated nanoparticles were present and their amount could be reduced by increasing the extrusion temperature

An innovative method named semi-solid casting (SSC) was proposed Zinc alloy AC43A reinforced by 30 nm β-SiC was used for samples preparation by SSC The SSC experiments were carried out by pouring ultrasonicated molten MMNC material (450°C) from a graphite crucible into a steel injection device, which was preheated to 400 °C Liquid MMnC was cooled down to 386 °C achieving less than 30% of solid fraction Then, the injection sleeve was inverted and placed on top of a steel mold The plunger was activated and the semi-solid material was injected into the mold The produced

samples showed strength properties comparable to those by ultrasound-assisted casting but with improved ductility

Solid Processes

Several solid methods were studied and developed for preparing MMnCs [13-15]

In particular, different powder metallurgy techniques were successfully employed in this respect Some papers focus on mechanical alloying which a powder metallurgy technique is consisting in repeated cold welding, fracturing and re-welding of powder particles in a high-energy ball mill This technique is of fundamental importance since it allows achieving a better dispersion of nano-powder into the composite by breaking up the ceramic clusters It can also be exploited for the formation of alloys by diffusion mechanisms starting from pure metals, and to produce performs by in situ reaction of nano-reinforcements Therefore, mechanical alloying, which cannot be separated from the opportunity of breaking up of the nano-ceramic clusters, is a value-added option offered by this particular processing route

It has been proved that the presence of nanoparticles can accelerate the milling process (stimulating plastic deformation, cold welding, and particle fragmentation) and grain refinement mechanism Process control agent (PCA) has a strong influence on the morphological evolution of powders during ball milling The addition of 1.5% stearic acid as PCA prevents cold welding of Al particles during ball milling and leads to an increase of hardness of the hot-compacted samples Speed and time of milling, mass of balls and powder, and ball diameter also contributes to final hardness development In particular, a pronounced decrease in energy transfer from the balls to the powder was found by raising the amount of balls

High-energy ball milling proved to be a suitable technique for production of in situ MMnCs [16-18] Al–TiN composite was prepared by milling elemental Al and Ti powders with ring-type organic compound pyrazine in benzene solution Mg 5 wt.% Al alloy in situ reinforced with TiH2 was also prepared by mechanical alloying of elemental

powder of Mg, Al and Ti, using polyethylene-glycol to provide hydrogen for the formation of TiH2 and to prevent excessive cold welding during ball milling After

attritioning, the powders were cold isostatically pressed (CIP), extruded and thermal treated The mechano-chemically milled specimens showed very fine microstructure and good dispersion of fine reinforcements, a slight increase in YS and ductility was observed Iron-wustite (Fe–FeO) nanocomposites were also produced by mechano-chemical processing starting from Fe and Fe2O3 powder with different mole ratios These

materials showed a ferromagnetic-like behavior, which was interpreted according to spinel-like defect, clusters In another work, Mg-5Al-10.3Ti-4.7B (wt.%) powder mixture was ground using high-energy ball milling and subsequently extruded They observed the formation of non-equilibrium Ti3B4 phase in extruded samples The in situ

formation of TiB2 also investigated via chemical reaction among Al, TiO2 and B2O3 The powders were cold compacted into green compacts and sintered at different temperatures

By this method, 53% of increment in YS and UTS was achieved In situ TiB2 reinforced

Cu alloy composite was indeed achieved via argon atomization at 1400 °C followed by hot isostatic pressing (HIP) at 200 °C under 200 MPa pressure

Moreover, Cu-Al2O3 nanocomposites have been prepared by two chemical routes: through decomposition of Al(NO3)3 to Al2O3 by calcination of a paste of CuO–Al(NO3)3

followed by H2 reduction and sintering, or through hydrolysis of Al(NO3)3 solution

followed by calcination, reduction and sintering The latter method led to the formation

of finer Al2O3 (30 nm vs 50 nm) and promoted enhanced properties in terms of relative

density, microhardness and abrasive wear resistance [15] Submicron-sized titanium carbide was successfully sintered from the reaction of Ti salt (K2TiF6) and activated

carbon, by controlling the degree of reaction through the temperature and the amount of

C In this respect, it was observed that at low temperatures, formation of Al3Ti was predominant while at high temperatures (above 1000 °C), the intermetallic compound was not stable and TiC was preferentially formed [19]

2.1.2 Strengthening in particulate reinforced metal matrix composites

Discontinuous-reinforcement MMCs are all characterized by the presence of relatively small particles of reinforcing material spread uniformly throughout the metallic matrix The reinforcement can be in the form of a particulate, whisker or platelet

in range from several hundred microns to less than a micron in size They are usually some form of ceramic powder; SiC, Al2O3 and B4C are common reinforcements for

discontinuous MMCs The particles are normally randomly oriented in the matrix, with

no direction receiving preferential reinforcement The properties of the composite are therefore isotropic It should be noted that for reinforcements such as whiskers ad platelets, which are not the same size in all dimensions, some operation such as extrusion and rolling may cause some particle realignment This can introduce some anisotropy in the properties of the composite [20]

The primary similarity between particulate-reinforced composites is the isotropic nature of their properties The high stiffness and strength of the ceramic particulate augments the stiffness of the matrix The orientation of the particles is random, and so the reinforcement aids the matrix no matter in what direction the forces are applied Certain processes, however, can cause preferential alignment For example, rolling can

cause flat platelets to align themselves in parallel layers, like leaves in a book, making the stiffness higher in the two long dimensions of the sheet This requires a high degree

of material flow during the forming and doesn’t occur in more moderate treatments [21]

Another similarity is the dependence of the composite’s properties on the amount

of reinforcement present The modulus of elasticity of aluminum with no reinforcement

is about 70 GPa When enough SiC particulate is added to equal 15% of the overall volume, the modulus increases to 98 GPa The modulus increases monotonically as the reinforcement loading increases until, at 40 vol% it is 147.6 GPa, more than twice the stiffness of the unreinforced material

As the percentage of ceramic increases, the composite displays more and more properties of the ceramic One of these properties is improved stiffness, and another is decreased ductility As the volume of ceramic increases, the amount of matrix between the particles decreases The ability of the matrix to deform is the basis of the composite’s ductility The ductility of the unrienforcd matrix in the case of aluminum can be over 10%

as measured by the strain to failure of test samples As the percentage of reinforcement increases to 40vol%, the strain to failure decreases to less than 2% For comparison, typical graphite-aluminum continuous-filament composite containing 40vol% graphite has less than 1% strain to failure

As a result of the reduced ductility of heavily loaded composites, manufacturing processes requiring the large-scale movement of materials must be carried out at high temperatures and the material moved slowly, to allow the limited matrix present to conform Processes must be designed to keep the material subjected to compressive loads during its movement to avoid cracking

2.1.3 Previous works on production process of Fe-TiB2 composite

Titanium diboride has several attractive properties It has high hardness (> 3000HV), a high melting temperature (3225oC), low density (4.5 g/cm3) and good

thermal (96 W/m/K) and electrical (22x106 Ω cm) conductivity (ref Material Properties

of Titanium Diboride) Thus, steel-matrix composites, which incorporate TiB2 as the

reinforcing phase, have increased stiffness, hardness, and wear resistance, along with reduced coefficients of thermal expansion and only a moderate decrease in thermal conductivity properties

Synthesis methods of TiB2 can be summarized as follows: powder metallurgy,

casting, laser cladding, plasma transferred arc (PTA), aluminothermic reduction, propagating high temperature synthesis (SHS) and spark plasma sintering

self-Aaustenitic 304 stainless steel composites reinforced with various volume fractions of TiB2 particles by means of hot-iso-staic pressing (HIP) have been prepared

by S.C Tjong [22] For MMCs fabrication, the ceramic particle and stainless steel powders were mixed for 5 h, and these powder mixtures were filled into 304 stainless tube containers The samples were HIPed at 1180oC and 100 MPa for 1 h The tensile

properties of specimens investigated reveals that the additions of hard ceramic particles improve the mechanical strength of stainless steels at the expense of its ductility In general, the 0.2% offset yield strength of MMC tends to increase with increasing TiB2

volume content Sulima I et al [23] investigates the effect of the reinforcing ceramic particles on the mechanical and tribological properties and microstructure of the steel-TiB2 composites The austenitic AISI316L stainless steel reinforced with 10 vol.% and

20 vol.% TiB2 particles were produced using the high temperature-high pressure (HT-HP)

method The sintering process was carried out at a pressure of 7.0±0.2 GPa and temperature of 1200°C for 60 seconds The materials were characterized by very high

level of consolidation, which was equal to 96% for composites with 10 vol.% and 20 vol.% TiB2 particles The composites exhibited higher Young’s modulus, Vickers hardness and

compression strength when compared with conventionally austenitic AISI316L stainless steel The addition of 20 vol.% of TiB2 particles to steel caused significant reduction of

the values of friction coefficient The SEM studies of composites allowed to reveal TiB2

phase along grain boundaries In case of the composite with 20 vol.% TiB2, the

continuous layer of ceramic along the grain boundaries was observed Beside TiB2, TiC

have also proved their suitability in iron and steel matrices due to their high hardness, good wettability, low density, and their relatively chemical stability with the iron and steel matrices Hence, TiB2 and TiC particulates simultaneously reinforced in Fe matrix

have been interested Stainless steel matrix composites reinforced with TiB2 or TiC

particulates have been in situ produced through the reactive sintering of Ti, C and FeB [24] Stoichiometric amounts of Ti, C, FeB and 465 stainless steel powders were blended homogeneously using ball mill in two different compositions of 55 and 70 wt.% (TiB2+TiC) reinforcements After ball milling, the mixed powders were compacted into

pellets with 37.5 mm in diameter and 5 mm in height at a pressure of 600MPa The green compacts were sintered with pressure 1*10−2 Pa at 1400°C for 90 min The X-ray diffraction analysis ensures the completion of the reaction and the absence of any secondary reaction The microstructure of composites showed round-TiC and hexagonal-TiB2 particulate shapes The particle sizes of these reinforcements range from 0.2 to 8

μm The relative density and hardness of the composites gained 95.3% and 85.7 HRA for

55 wt.% (TiB2+TiC) reinforcements and 94.5% and 91.4 HRA for 70 wt.% (TiB2+TiC)

reinforcements

Liquid state fabrication of Fe-TiB2 composite involves incorporation of dispersed

phase into a molten matrix metal, followed by its solidification By utilizing the heat of

the liquid steel during casting to induce the SHS reaction of reactants, Q.C Jiang and his coworkers attempt to produce the locally reinforced steel matrix composites using an SHS-casting route, and have successfully fabricated the TiC or/and TiB2 particulates

locally reinforced steel matrix composites using Al–Fe–Ti– B, Al–Ti–B4C, Fe–Ti–B4C–

C and Ni–Ti–B4C systems [25-32] SHS-casting route includes dry-mixing of powder

blends under an argon atmosphere for a given time, and then powder mixtures were uniaxially pressed into cylindrical preforms under pressures ranging from 80 to 85MPa

to obtain densities of 70±5% theoretical density After being dried in a vacuum oven to remove any trace of moisture, the preforms were placed and fixed on the bottom of the sand mold Subsequently, the medium carbon steel melt prepared in a 5 kg medium-frequency induction furnace with a temperature of about 1500-1600oC was poured into

the mold to ignite the SHS reaction of the preforms After solidification and cooling, composite castings were removed and sectioned in the side position In another work [33], an Fe-TiB2 master alloy manufactured by an SHS process was introduced into a

medium carbon, low-alloy steel matrix using a liquid route Master alloys were incorporated into the molten steel by directly pouring the powders onto the top of the melt and stirring with a ceramic impeller Using the Fe-70 wt% TiB2 master alloy,

composites were produced which contained two different levels of TiB2 additions (referred to as composites C2.5TiB2 and C7.5TiB2) The use of SHS and casting routes

provides a promising process method for fabrication of steel matrix composites because

of its significantly inherent simplicity, potential cast-effectiveness in scale-up manufacturing, as well as near-net shape processing capability However, most of microstructure of the composites above show inhomogeneous distribution of TiC or/and TiB2 particulates TiC and TiB2 particulate size distributes in large range from several to

several tens micrometers Besides, some interphases also found in the composites by reaction of the steel melt with reactants or between reactants

Among the several techniques available to synthesize in situ particulate reinforced iron and steel matrix composites, it is well established that SHS has many attractive advantages such as high purity of product, low processing cost, and efficient in energy and time The mechanism of phase changes in the Ti-B-Fe system in a combustion wave for a mixture of Ti, B, and Fe powders and a FeB alloy-Ti mixture with the same proportion of elements is studied by O K Lepakova et al [34, 35] and [36] It is found that the mechanism of structure formation depends significantly on the type of contact between the initial components An X-ray phase and X-ray spectrum and structural microanalyses of the quenched layers of the specimens show that the first contact melts occurring in the combustion wave are ferroboron (the first type of mixture)

or ferrotitanium (the second type) melts In the first case, the calculated high-melting compound TiB2 forms as a result of the interaction between the two melts; in the second

case, it forms as a result of the interaction of the melt with solid ferroboron, which, in turn, determines the different type of microstructure of the final combustion products The highly disperse and more homogeneous structure of the product forms after combustion of the second-type mixture In another work, O K Lepakova and his coworkers [33] study phase composition and the microstructure of combustion products

of the three-component system Ti-B-Fe and the search for optimum technological processes for the manufacture of articles and coatings based on them Investigations were conducted for two types of powder mixtures; the first was a mixture of elemental powders Ti, B, Fe and the second was a mixture of ferroboron alloys, FeBn, with Ti

Cylindrical samples with a porosity of 35–40% were made from powder mixtures by cold pressing and reacted by SHS at constant pressure of 0.5MPa in inert argon The

ignition of the pressed samples was monitored with the help of a radiant heat color tungsten coil The results indicate that the increase in iron content in the reaction mixtures for both types of charge results in a decrease in the TiB2 particles size The

combustion products of the reaction of FeB alloys with Ti are characterized by a more finely dispersed structure in comparison with combustion products of mixtures of elemental Fe, B and Ti powders This is explained by the fact that maximum combustion temperatures of mixtures, consisting from elemental powders are essentially higher, than the combustion temperatures of mixtures in which alloys are used as reagents This work also investigated the effect of reactant size on the size and distribution of reinforcement The microstructure of combustion products of mixtures of Fe + Ti + 2B composition composed of finely dispersed starting powders (Ti≈15 μm; Fe ≈10 μm) The microstructure consists of TiB2 crystals with sharp edges with a size not more than 5 μm,

uniformly distributed in the matrix consisting mainly of iron and small amounts of titanium Increasing the average size of the titanium and iron particles in the starting mixtures leads to the formation of microareas with microstructures, different from equilibrium The products obtained directly from the SHS process are generally porous,

so this limits the application of SHS SHS plus mechanical pressing is an easy and economical way to fabricate cermets This approach combines SHS and dynamic consolidation to fabricate full-density composites rapidly in a single processing operation Immediately after SHS, when the sample is still red hot and soft, a quick and large mechanical force is applied to the sample Densification is achieved by forcing grain rearrangement and liquid phase capillary infiltration SHS plus quick pressing actually is called the SHS/QP technique Z.Y Fu et al prepared TiC–Ni and TiB2–Fe SHS/QP

technique [37, 38] Well-mixed samples were pressed uniaxially in a steel die into disks

of 5 cm diameter and about 1 cm height at 50% of theoretical density The disk samples