Design, synthesis, and properties of titanium metal matrix composites based on Ti-B-Fe System

The present study focuses on the design and synthesis of titanium metal matrix composites (MMCs) based on the Ti-B-Fe system in which the Fe-stablized beta titanium phase serves as the matrix phase and the titanium boride (TiB) compound serves as the reinforcement phase. One of the principal objectives of this work is to use ternary CALPHAD calculations to determine the phase fields such that the optimum processing temperatures for Electric-Field-Activated-Sintering (EFAS) can be found. The second objective is to use EFAS to rapidly synthesize the selected compositions. Additionally, microstructures, phases, flexural strength, and fracture toughness properties of selected composites, which have been processed by EFAS, have been investigated. Ti-B-Fe phase diagram showed the isothermal section was generated by three elements as Ti, B, and Fe, and the titanium-rich corner showed the phase field of each MMCs compositions and the ternary phase diagram also matched accurately with HT-XRD results. In general, it is found that the hardness of the composites increases with an increase in the amount of TiB phase. These hardness levels are 530, 680, and 1080 kg/mm2, for volume fraction of boride phases as 0.22, 0.27, and 0.79, respectively. The composite with 30 mol. % of B composition showed a very interesting microstructure in which a uniform distribution of TiB phase (Vf = 0.79) separated by β-Ti phase with both phases forming in-situ under EFAS. The composite exhibited good mechanical properties with good combination of hardness. The flexural strength of this composite was found to be in the range of 556-727 MPa and the fracture toughness shows a moderate value of 10.3 MPa√m. Fractographic analysis revealed that the strength and fracture toughness are limited by the brittle behavior of the β-Ti phase. Additional research is planned to introduce alloying elements for β-Ti phase to enhance ductility and toughness of the composites.

DESIGN, SYNTHESIS, AND PROPERTIES OF TITANIUM METAL MATRIX COMPOSITES BASED ON TI-B-FE SYSTEM by Ahmed Abdulaziz Degnah A dissertation submitted to the faculty of The University of Utah in partial fulfillment of the requirements for the degree of Doctor of Philosophy Department of Metallurgical Engineering The University of Utah December 2019 Copyright © Ahmed Abdulaziz Degnah 2019 All Rights Reserved The University of Utah Graduate School STATEMENT OF DISSERTATION APPROVAL Ahmed Abdulaziz Degnah The dissertation of has been approved by the following supervisory committee members: , Chair Ravi Chandran 9/12/2019 Date Approved , Member Sivaraman Guruswamy 9/12/2019 Date Approved , Member Krista Carlson 9/12/2019 Date Approved , Member Ashley Spear 9/12/2019 Date Approved , Member Anthony Sanders 9/12/2019 Date Approved and by Michael F. Simpson the Department/College/School of , Chair/Dean of Metallurgical Engineering and by David B. Kieda, Dean of The Graduate School. ABSTRACT The present study focuses on the design and synthesis of titanium metal matrix composites (MMCs) based on the Ti-B-Fe system in which the Fe-stablized beta titanium phase serves as the matrix phase and the titanium boride (TiB) compound serves as the reinforcement phase. One of the principal objectives of this work is to use ternary CALPHAD calculations to determine the phase fields such that the optimum processing temperatures for Electric-Field-Activated-Sintering (EFAS) can be found. The second objective is to use EFAS to rapidly synthesize the selected compositions. Additionally, microstructures, phases, flexural strength, and fracture toughness properties of selected composites, which have been processed by EFAS, have been investigated. Ti-B-Fe phase diagram showed the isothermal section was generated by three elements as Ti, B, and Fe, and the titanium-rich corner showed the phase field of each MMCs compositions and the ternary phase diagram also matched accurately with HT-XRD results. In general, it is found that the hardness of the composites increases with an increase in the amount of TiB phase. These hardness levels are 530, 680, and 1080 kg/mm2, for volume fraction of boride phases as 0.22, 0.27, and 0.79, respectively. The composite with 30 mol. % of B composition showed a very interesting microstructure in which a uniform distribution of TiB phase (Vf = 0.79) separated by β-Ti phase with both phases forming in-situ under EFAS. The composite exhibited good mechanical properties with good combination of hardness. The flexural strength of this composite was found to be in the range of 556-727 MPa and the fracture toughness shows a moderate value of 10.3 MPa√m. Fractographic analysis revealed that the strength and fracture toughness are limited by the brittle behavior of the β-Ti phase. Additional research is planned to introduce alloying elements for β-Ti phase to enhance ductility and toughness of the composites. iv Dedicated to my father Abdulaziz, who passed away peacefully, during my PhD program (08/01/2019) TABLE OF CONTENTS ABSTRACT ....................................................................................................................... iii LIST OF TABLES ............................................................................................................. ix PREFACE ........................................................................................................................... x Chapters 1. INTRODUCTION .......................................................................................................... 1 1.1 Titanium and its Alloys ........................................................................................ 1 1.2 Metal Matrix Composites (MMC) ....................................................................... 1 1.3 In-situ Composites ............................................................................................... 2 1.4 Electric-Field-Activated-Sintering (EFAS) ........................................................ 3 1.5 Ti-TiB Composites .............................................................................................. 3 1.6 Mechanical Properties of Ti-TiB Composites .................................................... 4 1.6.1 Elastic Modulus ...................................................................................... 4 1.6.2 Strength ................................................................................................... 5 1.6.3 Ductility .................................................................................................. 6 1.6.4 Creep ........................................................................................................ 6 2. LITERATURE REVIEW .............................................................................................. 7 2.1 Crystal Structure of Titanium ............................................................................. 7 2.2 Deformation of Pure Titanium ............................................................................. 7 2.3 Effect of Interstitials on the Mechanical Properties of Titanium ......................... 8 2.3.1 Effect of Interstitials on Titanium Phase Transformation ..................... 11 2.4 Titanium Alloys ................................................................................................. 12 2.4.1 Alloying of Titanium: Alpha and Beta Stabilizing Elements ................ 14 2.5 Metal Matrix Composites (MMCs) .................................................................. 17 2.5.1 Titanium Metal Matrix Composites (Ti-MMCs) ................................... 18 2.5.2 Methods of Reinforcement ................................................................... 19 2.6 Structure and Properties of Titanium Boride ..................................................... 20 2.6.1 Ti-B Phase Diagram .............................................................................. 22 2.6.2 Diffusion of Boron and Growth of TiB Whiskers ................................. 22 2.7 Processing Techniques for Ti-TiB Composites ................................................. 24 2.7.1 Electric-Field-Activated-Sintering (EFAS) Processing of Ti-TiB Composition .................................................................................................... 26 2.8 Microstructures of Ti-TiB Composites .............................................................. 27 2.9 Estimation of TiB Volume Fraction ................................................................. 29 2.10 Mechanical Properties of Discontinuous Reinforcements in MMC ................ 30 2.10.1 Strength and Stiffness of Ti-TiB Composites ...................................... 30 2.10.2 Flexure Strength and Fracture Toughness .......................................... 33 2.11 CALPHAD Approach for Alloy Design .......................................................... 35 3. MATERIAL DESIGN AND EXPERMENTAL PROCEDURE ................................ 66 3.1 Design of Composite Composition and Processing ........................................... 66 3.2 Phase Equilibria in Ti-Fe System ..................................................................... 67 3.3 Materials and Experimental Approach ............................................................. 68 3.4 Microstructure and Phase Characterization ...................................................... 69 3.5 Tensile Testing ................................................................................................... 70 3.6 Flexure Testing ................................................................................................. 70 3.7 Fracture Toughness Testing ............................................................................... 71 3.8 Estimation of TiB Volume Fraction ................................................................. 71 4. RESULTS AND DISCUSSION ................................................................................... 76 4.1 Ternary Ti-B-Fe System .................................................................................... 76 4.2 High Temperature X-ray Diffraction Study ....................................................... 77 4.3 Pseudo-binary Ti-B-Fe Phase Diagram with Constant Fe and Variation in B Content ..................................................................................................................... 78 4.3.1 Densification Behavior ......................................................................... 79 4.3.2 X-ray Diffraction Analysis ................................................................... 81 4.3.3 Microstructure ........................................................................................ 83 4.4 Pseudo-binary Ti-B-Fe Phase Diagram with Constant B and Variation in Fe Content ..................................................................................................................... 84 4.4.1 Densification Behavior ......................................................................... 84 4.4.2 X-ray Diffraction Analysis ................................................................... 85 4.4.3 Microstructure ........................................................................................ 85 4.5 Mechanical Properties of Ti-B-Fe System with Variation in B Content ........... 86 4.5.1 Hardness ................................................................................................. 86 4.5.2 Tensile Properties .................................................................................. 87 4.5.2.1 Fractography ............................................................................ 88 4.5.3 Flexure Strength ..................................................................................... 88 4.5.4 Fracture Toughness ................................................................................ 90 5. CONCLUSIONS......................................................................................................... 117 6. DESIGN, SYNTHESIS AND PROPERTIES OF TITANIUM BORIDE CERMET MATERIAL BASED ON TI-B-FE-MN SYSTEM ....................................................... 119 6.1 Introduction ...................................................................................................... 119 6.2 CALPHAD Approach ...................................................................................... 122 6.3 Experimental .................................................................................................... 123 6.4 Results and Discussion .................................................................................... 125 6.4.1 Processing ........................................................................................... 125 vii 6.4.2 Microstructure ..................................................................................... 126 6.4.3 Mechanical Properties .......................................................................... 127 6.5 Conclusions ...................................................................................................... 128 7. COMPUTATIONAL DESIGN, PHASE EQUILIBRIA, AND PROCESSING OF TITANIUM METAL MATRIX COMPOSITES IN TI-B-MO-FE-AL SYSTEM ....... 139 7.1 Introduction ...................................................................................................... 139 7.2 Design and Experimental Procedure ................................................................ 142 7.2.1 Design of MMC Composition and Processing ................................... 142 7.2.2 Materials and Experimental Method ................................................... 144 7.2.3 Volume Fraction Calculation ............................................................... 145 7.3 Results and Discussion .................................................................................... 146 7.3.1 EFAS Processing ................................................................................ 146 7.3.2 X-ray Diffraction Data ........................................................................ 147 7.3.3 Microstructure ...................................................................................... 147 7.4 Mechanical Properties ...................................................................................... 147 7.4.1 Hardness Evaluation ........................................................................... 147 7.4.2 Flexure Strength .................................................................................. 148 7.4.3 Fracture Toughness .............................................................................. 149 7.4.4 Tensile Properties................................................................................. 149 7.5 Conclusions ...................................................................................................... 150 REFERENCES ............................................................................................................... 174 viii LIST OF TABLES Tables 2.1 Slip and twinning systems in α-Ti ............................................................................. 65 3.1 Compositions of the synthesized MMCs .................................................................... 75 4.1 Total shrinkage in various MMCs in the Ti-B-Fe system ....................................... 112 4.2 Density and atomic/molecular weight of the elements/compounds ........................ 112 4.3 MMC compositions and densities ............................................................................ 112 4.4 The effect of composition on the d-spacing of (110)β-Ti and (200)β-Ti ...................... 113 4.5 Calculation of X-ray diffraction pattern for β-Ti phase ............................................ 113 4.6 Calculation of X-ray diffraction pattern for TiB phase ........................................... 113 4.7 Calculation of X-ray diffraction pattern for TiFe phase ........................................... 114 4.8 Volume fraction of phases in the MMCs .................................................................. 114 4.9 Shrinkage in various MMCs in the Ti-B-Fe system ................................................. 114 4.10 Experimental and theoretical densities of Ti-TiB MMCs with varying Fe content 115 4.11 Volume fraction of phases calculated using XRD data ......................................... 115 4.12 Tensile properties of composite Ti-10B-10Fe ........................................................ 116 7.1 Fracture toughness of Ti-10B-3.7Mo-4.3Fe-2.96Al v/s titanium alloys MMC ....... 173 7.2 Tensile properties of Ti-10B-3.7Mo-4.3Fe-2.96Al v/s titanium alloys MMC ......... 173 PREFACE This study focuses on the synthesis of novel metal matrix composites (MMCs) in Ti-B-Fe system along with an investigation of their mechanical properties including hardness, tensile strength, flexure strength, and fracture toughness. The major objectives of this study are: • To use the CALPHAD approach for designing the compositions of titanium boride reinforced titanium MMCs. This involves the determination of the high temperature phase equilibria, based on thermodynamic data for the Ti-B-Fe system. The merit of this approach is that desirable compositions can be rigorously identified for processing at a relatively low temperature. Thermo-Calc software facilitates the phase-field information, which helps in optimizing the processing conditions. This saves time and cost associated with conventional methods of empirical alloy synthesis. • To study the high temperature stability of phases of the MMCs, using high temperature X-ray diffraction studies of the MMCs. This will help to evaluate any high temperature phase transformations and will enable comparison with Thermo-Calc calculated phase fields to ensure consistency in the approach. • To synthesize fully dense MMCs in the Ti-B-Fe system by Electric-Field- Activated-Sintering (EFAS). This is to involve the variation of B content while Fe (beta stabilizer) content is maintained constant in one set and the variation of Fe content with the B content is maintained constant in the second set. • To optimize the ratio of B:Fe and to understand the relative effects of the two alloying elements on the stability of the beta titanium matrix. This ratio is also expected to control the morphology of the TiB reinforcement and its subsequent effects on mechanical properties. • To analyze the microstructures of the sintered MMCs and determine how they vary between the different compositions. The morphology and the stability of the phases are to be studied through optical as well as scanning electron microscopy and X-ray diffraction methods. These characterization methods will assist in determining the optimum composition for enhanced mechanical properties. • The final objective is to evaluate the mechanical properties of selected MMC compositions by various mechanical testing methods including Vickers hardness, tensile strength, and four-point bend test to evaluate the flexure strength and fracture toughness. The following hypotheses have been investigated for this research study:  Addition of B will promote volume fraction control.  Addition of Fe will achieve densification.  Addition of Mn will eliminate intermetallic formation of TiFe phase.  Addition of Al causes densification and eliminates intermetallic formation of TiFe phase.  Addition of Mo will eliminate intermetallic formation of TiFe phase.  Determine the optimum conditions of the EFAS using CALPHAD approach.  Determine the accuracy of the CALPHAD approach in the composites system by validating it, through the high temperature X-ray diffraction.  Using β-Ti as a metallic phase, to improve ductility, toughness, and investigating xi whether TiB can improve strength and hardness of the composite. Publications and Presentations 1. Jindal, V.; Sarda, A.; Degnah, A.; Ravi Chandran, K. S., Effect of iron & boron content on the Spark Plasma Sintering of Ti-B-Fe alloys. Advanced Powder Technology 2019, 2 (30), 423-427. 2. Ranjan, A.; Jindal, V.; Tyagi, R.; Degnah, A.; Ravi Chandran, K. S., Optimization of wear properties of TiB reinforced Ti composites containing 10-30 at% Iron. (Manuscript in preparation). 3. Jindal, V.; Sarda, A.; Degnah, A.; Ravi Chandran, K. S., Thermodynamic assessment of Ti-rich portion of the Ti-B-Fe system. (Manuscript in preparation). 4. Degnah, A.; Du, J.; Ravi Chandran, K. S., Design, synthesis and properties of titanium boride cermet material based on Ti-B-Fe-Mn system. (Manuscript in preparation). 5. Degnah, A.; Du, J.; Ravi Chandran, K. S., Computational design, phase equilibria, and processing of titanium metal matrix composites in Ti-B-Mo-Fe-Al system. (Manuscript in preparation). 6. Degnah, A.; Jindal, V.; Ravi Chandran, K. S. In Computational Thermodynamics Assisted Process Design of Ti-B-Fe System, Presented at Materials Science & Technology (MS&T), 2016. 7. Degnah, A.; Jindal, V.; Ravi Chandran, K. S. In Computational Phase Equilibria and Design of Metal Matrix Composites in Ti-B-Fe System, Presented at Materials Science & Technology (MS&T), 2017. 8. Degnah, A.; Jindal, V.; Ravi Chandran, K. S. In Computational Design, Phase Equilibria, and Processing of Titanium Metal Matrix Composites in Ti-B-Fe System, Presented at Powder Metallurgy and Additive Manufacturing of Titanium (PMTI19), 2019. Acknowledgements I would like to gratefully and sincerely thank my advisor Professor Ravi Chandran for his guidance, patience, understanding, and encouragement in all aspects of this scientific research. xii I would like to thank my committee members’ Professor Sivaraman Guruswamy, Professor Krista Carlson, Professor Ashley Spear and Professor Anthony Sanders for their insightful comments and suggestions. Many thanks to members my group for their technical assistance and motivation (Dr. Vikas Jindal, Jun Du, Alexander Lark, Husain Alnaser, and Somnaang Rou). I would like to extend my appreciation to my scholarship King Abdulaziz City for Science and Technology (KACST) for the financial support. Finally, I would like to acknowledge my family (Late Abdulaziz--father, Anisamother, Reem and Mram--sisters, Rami, Mahmmoud, and Ryan--brothers) for their love and support, providing the necessary strength, encouragement, and quiet patience during my graduate studies at The University of Utah. xiii CHAPTER 1 INTRODUCTION 1.1 Titanium and its Alloys Titanium (Ti) and its alloys generally have superior mechanical properties, relatively high strength to weight ratio, good biocompatibility, and high corrosion resistance. Titanium is lighter than steel and stronger than aluminum, which makes it desirable for many advanced engineering applications. As a result, Ti and its alloys are widely used in several applications, including in jet engines, marine equipment, automobile, chemical processing, and biomedical devices. 1.2 Metal Matrix Composites (MMC) Over the years, several studies [1] focused on improving the mechanical and physical properties of titanium alloys to make them viable for higher performance engineering applications. This motivated the development of metal matrix composites (MMC) based on titanium [1]. Here, the titanium matrix is reinforced with hard and stiff ceramic reinforcements such as particulates, short fibers, or continuous fibers. Titanium MMCs have attracted a significant research interest, since these composites can provide a combinations of high strength and stiffness arising from the ceramic phase reinforcement along with ductility and toughness inherited from the metal matrix. In-situ reinforcement 2 in MMCs is especially advantageous since it is associated with its lower cost and more flexibility in terms of processing than, for example, the continuous fiber composites. Some examples of ceramic reinforcements include SiC, Al2O3, TiB2, and TiB. Titanium MMCs can be produced by rapid solidification, arc melting, powder blending, and pressing [1, 2]. 1.3 In-situ Composites In-situ composites are composites containing multiple phases in the microstructure where the reinforcement can form within the matrix during the process of manufacturing. The in-situ composites are different from the ex-situ composites. In the ex-situ composites, the reinforcements are mixed with the matrix of the composite in a secondary processing step such as powder processing or melt infiltration. On the other hand, in in-situ composites, the phases are formed by a precipitation reaction. In the reaction, the reinforcement phase can form in different morphological forms. This can occur as continuous or discontinuous reinforcement. This reinforcement can be a ductile metal phase or a hard ceramic phase [3]. The in-situ process of forming a reinforcement phase can be done in a liquid state or in the solid state. In the liquid state, the precipitation reaction directionally solidified eutectics from under low undercooling whereas rapid solidification involves high undercooling because of larger driving force for nucleation and growth. Generally, the in-situ composites have several advantages compared to the ex-situ composites [3]. These advantages include: • Creep and fatigue resistance can be improved by having high strength reinforcements due to presences of small particles, which impede the movement of the dislocations in the matrix. 3 • In-situ reinforcements formed at high temperature are stable and do not dissolve/react with matrix like Ti-Al reinforced with TiB2. • Homogeneous distribution of the reinforcement can help to improve the mechanical properties as this does not facilitate the nucleation and propagation of cracks. 1.4 Electric-Field-Activated-Sintering (EFAS) Recently, Electric-Field-Activated-Sintering (EFAS) has emerged as a new technique to produce solid material in the form of densified composites. Composites can be processed in a short time by using the EFAS. The time for processing can be short due to the high heating rate generated from the high level of electric current passing through the die and the powder. The powders used in this technique can be conductive or nonconductive in nature. The nonconductive powders are heated by passing electric current through a conductive die. The conductive powders can accelerate the reaction sintering by generating local Joule heating and enhancing the diffusion process. The unique advantage of this technique is the rapid synthesis of the composite with nano grain size and uniform microstructure [4]. 1.5 Ti-TiB Composites In one of the first studies, Sahay et al. [5] showed that Ti-TiB in-situ composites can be made by reaction sintering during hot pressing. More recent work shows that MMCs based on Ti-B system can be synthesized using the EFAS approach [6, 7]. The powder mixture is sintered by heating the die or the powders by passing a current. The heating leads to solid state reaction and/or liquid phase formation facilitating diffusion of B from 4 TiB2 to Ti to form the TiB whiskers [6, 7]. This processing technique is cost and time efficient due to the rapid densification, and formation of uniform microstructure [8, 9]. The TiB can be formed in-situ in titanium matrix by taking advantage of a simple reaction: Ti + TiB2 → 2TiB [5]. The processing is done at a suitable temperature (~1350 °C) such that during the sintering process, the densification of the composites is also achieved. The volume fraction of TiB can be varied to modify the composite to acquire desirable mechanical properties. By adjusting the proportion of Ti and TiB2 powder before reaction sintering, either fully ceramic, or cermet, or MMC can be formed by in-situ reaction sintering [10]. When the volume fraction of TiB is in the range of 0.1 to 0.3, a MMC can be formed, if it is more than 0.98, a nano-ceramic can be formed. For or a volume fraction range of 0.3 to 0.98, cermets can be formed [10]. 1.6 Mechanical Properties of Ti-TiB Composites 1.6.1 Elastic Modulus In general, TiB phase is formed in the shape of whiskers (reinforcement), which can improve the mechanical properties of the composite. It is found that as the volume fraction of the TiB phase increases to 0.10, the stiffness can increase within 20-25 % in the Ti-TiB composite [11]. The stiffness even becomes higher as the volume fraction of the TiB phase increases. For example, Ti24.3Mo composite containing 0.34 volume fraction of TiB exhibits a high stiffness of 171 GPa compared to pure Ti, which shows a stiffness value of 110 GPa [11]. The elastic modulus of the Ti-TiB composite depends on the modulus of the TiB phase (whisker) [12]. The B structural arrangement as a chain along the TiB whisker may provide significant elastic anisotropy, possibly more than that 5 obtained with TiB2 phase as a reinforcement. 1.6.2 Strength In the Ti-TiB composite, the ultimate tensile strength varies between 673 MPa and 1820 MPa [11]. This value depends on the matrix composition and the reinforcement level in the Ti-TiB composite. These values are considered high compared to that for pure commercial titanium, which is 550 MPa. The TiB reinforcement can increase the ultimate tensile strength with increase in the TiB phase. For example, Ti-4.3Fe-7Mo-1.4Al.1.4V composite without TiB phase as a reinforcement exhibits an ultimate tensile strength as 1080 MPa, but for the composite with 0.30 volume fraction of TiB, the ultimate tensile strength is 1820 MPa [13]. The high value of the ultimate tensile strength can decrease the ductility level of the same composite. The ductility level without the TiB reinforcement can be as high at 17.5 %, but with 0.30 volume fraction of the TiB phase, the ductility level is about 1 % [13]. It is important to note that it is very difficult to standardize the Ti-TiB composite based on strength level as a function of the TiB volume fraction level, as the difference in the strengthening level of Ti-TiB composite depends on the aspect ratio of the TiB whisker and presence of nano-whisker Moreover, presence of the interstitials elements such as oxygen, carbon, and hydrogen and different alpha and beta stabilizer elements make it more difficult to predict systematically the strengthening of Ti-TiB composite. These interstitial elements can make changes in the matrix in the strengthening level, but their concentration levels are not often known. 6 1.6.3 Ductility Ti-TiB composites with high volume fraction of TiB exhibit low ductility with some even showing zero ductility due to the brittleness of TiB reinforcements. The ductility of unalloyed alpha titanium matrix is sensitive to the interstitial elements. These elements can diffuse into the titanium matrix during the processing and can affect the ductility level of the composite. The ductility level of the Ti-TiB composite can be enhanced by having beta titanium matrix even though in those composites with a high volume fraction of TiB phase, the ductility level can be increased by incorporating a beta titanium ductile matrix. For example, in Ti-24.3Mo composite with the TiB volume fraction of 0.34, the composite shows ductility level as 0.9 with ultimate tensile strength of 1100 MPa [11]. On the other hand, Ti-53Nb composite with 0.34 volume fraction of TiB shows a high ductility level of 1.65 % due to presence of beta titanium ductile matrix [11]. Composite Ti-53Nb exhibits lower elastic modulus and strength compared to composite Ti-24.3Mo, due to the low aspect ratio of the TiB whisker. 1.6.4 Creep One of the main goals in Ti-TiB composite is to improve the creep resistance of the titanium composite compared to the pure titanium, which is soft at high temperature. There are only a few titanium alloys having reasonable or a good creep resistance at 600 °C. It is of significance to compare the Ti-TiB composite with the titanium alloys at 600 °C. Significant reduction in creep rates can be achieved with the introduction of TiB reinforcements in comparison to that observed titanium alloys. CHAPTER 2 LITERATURE REVIEW 2.1 Crystal Structure of Titanium Titanium forms in two crystal structures: the HCP (α) phase and the BCC (β) phase, with the allotropic phase transformation from HCP to BCC occurring at 883 °C [14]. The low temperature Ti phase has (α-phase) the HCP structure, which is shown in Figure 2.1 (a). Titanium forms as BCC β-phase when it crystallizes from a high temperature liquid phase. The structure of BCC Ti is shown in Figure 2.1 (b) [15]. Design of titanium alloys often requires balancing the appropriate relative volume fraction of alpha and beta phases to achieve a good combination of mechanical properties. 2.2 Deformation of Pure Titanium Titanium generally deforms by several slip systems. Generally, the deformation of metals and alloys in the HCP crystal structure is controlled by c/a ratio of the unit cell. The c/a ratio is the ratio of lattice parameters, c and a, with a=a1=a2=a3 (Figure 2.1 (a)). The ideal value of the c/a ratio is 1.633 as derived from the hard sphere model [16]. But in reality, the c/a ratios for HCP metals deviate from the ideal value due to bonding and electronic structure factors. If the value of the c/a ratio for a HCP metal is less than the ideal value of 1.633 (e.g. Mg, Co, Zr etc.), then the metal has high deformability by slip 8 [16]. This is attributed to the lower interatomic spacing in the “c” direction, which facilitates the movement of extra half-plane of an edge dislocation [16]. On the other hand, when the ratio c/a is higher than the ideal ratio, it will be very difficult to deform by slip. For example, c/a ratio for zinc is 1.85 and for cadmium is 1.89 [16]. The lattice parameters of the unit cell of pure α-Ti are a = 0.295 nm and c = 0.466 nm with the c/a ratio being 1.59 [17]. This value is less than the ideal c/a ratio. Due to this, the alpha phase generally deforms by slip, leading to significant ductility at room temperature [17]. Titanium can be deformed up to a reduction of 90 % in thickness without cracking, due to low c/a ratio [18]. However, some twining may occur along with slip depending on test temperature and strain rate [19]. Beta phase, on the other hand, is ductile due to its BCC structure, along with the occurrence of a relatively high number of slip system. The ductility in pure titanium can be attained from a combination of twinning and slip [20, 21]. Table 2.1 presents the slip and twinning systems in α-Ti [22]. Figures 2.2 (a) and (b) show the slip system and twin planes, respectively, in alpha titanium [23]. 2.3 Effect of Interstitials on the Mechanical Properties of Titanium The properties of titanium and its alloys can be affected by the presence of the interstitial elements such as N, C, O, and H. These interstitial elements, even in small amounts, can increase the yield strength significantly. Correspondingly, they also decrease the ductility. It is important to limit the interstitial content in titanium to low levels, which are generally <0.3 wt. %. Commercially pure titanium (CP-titanium) is produced as bars, rods, and sheets. Figure 2.3 shows the effects of interstitial elements on the strength and ductility of CP-Ti 9 [24]. It is evident that oxygen has the greatest impact on the yield and the tensile strength. When the oxygen content increases from 0 to 0.4 wt. %, the yield strength increases from 170 MPa to 480 MPa. However, as oxygen increase from 0 to 0.4 in wt. %, the elongation decreases from 24 to 15 %. It is important to consider the role of oxygen content in the grade classification of CP-Ti. CP-Ti can be classified into four grades, with each grade defined by the upper and lower limit of oxygen. The maximum oxygen content specified are 0.18, 0.25, 0.35, and 0.4 wt. % for grade 1, grade 2, grade 3, and grade 4, respectively [25]. In titanium alloys, it is always recommended to restrict oxygen content to less than 0.2 wt. % in the alloy to achieve a reasonable combination of ductility and strength [26]. For example, for the Ti-6Al-4V alloy, the specified upper limit of oxygen is 0.13 wt. % [27]. The interstitial elements such as carbon and nitrogen get incorporated in titanium during the processing and heat treatment steps, hence they are usually present in significant amounts [24]. However, the upper limits of C and N levels are specified for each grade of Ti. For example, for Ti-6Al-4V alloy, the specified limits of C and N are 0.08 and 0.05 wt. % [28]. Hydrogen, although an interstitial at low concentrations, can lead to the formation of titanium hydride (TiH2) at high H concentrations. The source of the hydrogen can be from the extraction process, due to immersion in or exposure to organic solutions such as methanol. Figure 2.4 shows the binary phase diagram of Ti and H, which reveals that the hydrogen solubility in titanium is strongly dependent on temperature [29]. At room temperature the solubility of the hydrogen in titanium is limited to 1.1 at. % and the excess hydrogen is combined and precipitated as TiH2. Although the solubility of hydrogen is low 10 at relatively low temperatures, there is a greater extent of H solubility at higher temperature above 300 °C (eutectoid temperature). Figure 2.4 illustrates the eutectoid reaction at 300 °C between α-Ti and δ titanium hydride (TiH2). Hydride in titanium can precipitate in the form of needles during the slow cooling from the high temperature [30]. At high temperatures, >300 °C, an increased lead of β-phase stabilization occurs at high concentrations of hydrogen. This aspect is used in processing of titanium with hydrogen as a temporary alloying element. Hydrogen is used as a temporary alloying element to improve the mechanical processability of titanium alloys. This approach is known as thermohydrogen processing (THP). In THP, hydrogen is alloyed by heating titanium in H2 atmosphere to a high degree to form the β-phase. This β-phase is quite ductile, allowing a high degree of deformation during low temperature processing, such as cold rolling. After the completion of thermohydrogen processing, the hydrogen can be removed by controlled annealing in vacuum. In this approach, the microstructure of the titanium alloy can be refined due to phase transformation and this leads to improvement of the mechanical properties of the titanium alloy [31]. Hydrogen, when present in titanium and its alloys, can lead to embrittlement. Even at low concentrations, hydrogen has the most deleterious effect on the ductility and toughness of titanium alloys. Figure 2.5 shows that the impact toughness of titanium decreases with increasing amount of hydrogen [32]. The adverse effect on ductility and toughness is caused by the precipitation of TiH2. These precipitates can initiate brittle fracture by cleavage along the planes of precipitation and this can severely reduce the ductility in titanium and its alloys. Hence, hydrogen content is restricted to very low levels (typically<100 ppm) [32]. 11 2.3.1 Effect of Interstitials on Titanium Phase Transformation Alloying elements, especially interstitials, have a strong effect on the alpha-to-beta transformation in titanium. Figure 2.6 shows the effect of oxygen content on the α-transus and β-transus temperatures, on the basis of data obtained from Ti-O alloys with five different oxygen contents, 0.125, 0.25, 0.50, 0.75, and 1 wt. %. As the oxygen content is increased, both the α-transus and β-transus lines shift upwards, increasing the transition temperatures [33]. It is therefore evident that O acts as a stabilizer of α-phase in titanium. Figure 2.7 shows the Ti-C phase diagram constructed on the basis of five different compositions of Ti, with 0.125, 0.25, 0.50, 0.65, and 0.87 wt. % C, to determine the effect of carbon on phase transformation in titanium alloys [33]. The minimum carbon solubility is 0.3 wt. % in alpha titanium. When the addition of carbon is increased to 0.3 wt. %, there is secondary precipitation of titanium carbide (TiC). It is to be noted that this solubility limit of C does not change much at room-temperature due to slow diffusion kinetics. The increase in the carbon content, however, does not seem to increase the alpha-to-beta transformation temperature. In the composition containing 0.125 wt. % carbon, quenching from 900 °C resulted in the formation of α+β structures, which indicated the presence of a small α+β phase field at high temperatures. However, C as alloying element in Ti is not preferred because at concentrations >0.3 wt. %, TiC forms. TiC is brittle and reduces the ductility and toughness of titanium. Nitrogen is often picked up by titanium during melting and processing. Figure 2.8 illustrates the Ti-N phase diagram [33]. There is a high degree of solubility of N in Ti. The solubility limit is 0.3 wt. % from the more extensive phase diagram [34]. There is also a 12 slight upward shifting of α and β transus lines with increasing nitrogen content, as determined from the study [33] of quenched Ti-N alloys. In addition to the interstitial elements, substitutional alloying elements can increase or decrease the α→β transition temperature in titanium. These substitutional elements are either known as alpha or beta stabilizers [35]. In the manufacturing of titanium alloys, the relative balances of alpha and beta stabilizes elements are adjusted to design titanium alloys for various applications. 2.4 Titanium Alloys Alloying elements in titanium significantly affect the strength and ductility. Therefore, alloying is the principal route followed to design titanium alloys. It is generally desirable to have at least 10 % ductility in titanium alloys to ensure reliability in applications [36]. Most of the common titanium alloys contain two phases, which is the alpha phase and the beta phase. The ductility can be controlled by incorporating the ductile beta phase by suitable alloying of titanium. The beta titanium phase with the BCC structure is more ductile than the alpha titanium phase. The ductility can be increased by addition of beta stabilizer alloying elements to increase the amount of β phase. On the other hand, the alpha phase is much stronger and stiffer than the beta phase, and is reflected in the values of their elastic moduli. Alpha titanium exhibits an elastic modulus value of about 110 GPa, whereas beta titanium has the elastic modulus values of about 90 GPa [35]. Titanium alloys can be classified into four major categories based on the amount and type of alloying elements as shown in Figure 2.9 [37]. These four groups are classified 13 based on the amount of [Al]eq and [Mo]eq level. The four classifications of the alloys are described below: (I) α-Alloys Alpha titanium alloys are alloys that are made of nearly 100 % α phase. Alpha alloys mostly contain substitutional elements stabilizing α phase and some minimal amount of interstitial elements. Examples of substitutional elements include Al and Sn. Interstitial elements such as C, N, and O can be also allowed in the composition to contribute to strength, as long as the ductility is not adversely affected. The strength of the alloy increases with an increase in the level of α stabilizing and/or interstitial elements. This is mainly attributed to the solid solution strengthening effect brought by these elements. The mechanical properties of this alloy depend on the grain size, texture, and any residual dislocation density arising from mechanical working. The strength is enhanced by refining the grain size by mechanical working and annealing, during the rolling or forging operations [25]. (II) α+β Alloys α+β titanium alloys are made of both alpha and beta phases, which render the alloys a good combination of strength and ductility. The class of α+β alloys can be further classified into two subgroups depending on the amount of α or β stabilizers: α-rich alloys and β-rich alloys. The α-rich alloys exhibit high strength and low ductility (e.g. Ti-8Al1Mo-1V), as these alloys contain a relatively higher content of alpha stabilizers. The microstructure of a α-rich alloy consists of a very small amount (< a few %) of beta phase [25]. On the other hand, β-rich alloys contain a relatively high amount of beta stabilizers, which render the alloy with higher ductility and with high strength (e.g. Ti-6Al-4V alloy). 14 The amount of β-phase in these alloys is about 10 % by volume. Among the α+β alloys, Ti-6Al-4V is the most widely produced and used alloy in engineering applications [38]. (III) β-Alloys Beta titanium alloys are the most ductile and tough set of alloys among all titanium alloys. Beta alloys can be further classified into two subgroups: metastable and stable beta alloys. In metastable beta alloys, the beta phase is not fully stable at room temperature. The phase field of these alloys is located between the Ms curve and the intersection of the betatransus curve with the x-axis (Figure 2.10). The metastable alloys can be heat treated either by quenching and aging, or by slow cooling, from above the β transus temperature, with and without mechanical working. Recrystallization and transformation during heat treatment can be used to control the size and distribution of α and β phases. On the other hand, in stable beta alloys, the beta phase is completely stabilized down to room temperature by a high degree of alloying with β-stabilizers. The phase field of this class of alloys is located beyond the beta transus curve intersecting with the composition axis (Figure 2.10) [15]. These alloys can be cold-worked and strain hardened to a great extent to improve the mechanical properties. This is largely due to the absence of any phase transformation, due to the stable β phase at room temperature [15]. 2.4.1 Alloying of Titanium: Alpha and Beta Stabilizing Elements The α phase in titanium is stable at temperatures lower than the α-β transition temperature, which is 883 °C. In order to increase the strength of alpha titanium alloys, alloying elements (alpha stabilizers) are incorporated in the matrix. The alpha stabilizers can increase the α to β phase transition temperature beyond 883 °C [39]. 15 Figure 2.11 shows the effect of Al on the phase fields in Ti-Al system [39]. The phase diagram shows the variation of α/(α+ α2) and ( α+ α2)/ α2 solvus lines as a function of the Al content. Ti alloys containing more than 8 wt. % Al fall within the α+α2, indicating the formation of intermetallic phase of α2 (Ti3Al). When the Al amount is increased to about 12 wt. %, at temperature 575 °C, the complete formation of intermetallic phase α2 (Ti3Al) is formed. In addition, the Ti-Al phase diagram shows how the alpha and beta transus temperatures of titanium are affected by the addition of Al. Design of α-Ti alloys is generally based on the addition of aluminum to a maximum of about 8 wt. % due to the possibility of formation of intermetallic phase. The formation of intermetallic phase Ti3Al, which severely reduces the ductility of titanium alloys, generally occurs in the alloy compositions above 6-8 wt. % of Al. Therefore, aluminum content should be restricted to less than ~8 wt. % to avoid the formation of the intermetallic phase [39]. The most common and efficient alpha stabilizer element is Al [40]. The combined effect of α-stabilizing elements in wt. % is expressed as Al-equivalent, [Al]eq, in titanium alloys [17]: [Al]eq = [Al] + [Zr] 6 + [Sn] 3 + 10 [O] (2.1) Beta alloying elements in titanium stabilize the BCC phase, which is desirable for enhancing the ductility in titanium alloys. In order to increase the amount of beta phase titanium at room temperature, a high degree of beta stabilization is necessary [17]. The β stabilizing elements include Mo, Ta, Nb, W, V, Cr, Ni, Mn, Co, and Fe. These elements lower the α to β transition temperature from 883 °C and increase the beta phase stability at low temperatures. The most efficient beta stabilizing element is molybdenum. The 16 combined effect of beta alloying elements in wt. % is expressed as Mo-equivalent, [Mo]eq according to the equation presented below [17]: [Mo]eq = [Mo] + [Ta] 3 + [Nb] 3.6 + [W] 2.5 [V] + 1.5 + 1.25[Cr] + 1.25[Ni] + 1.7[Mn] + 1.7[Co] + 2.5[Fe] (2.2) Mo-equivalent plays an important role in designing beta titanium alloys with an appropriate amount of beta phase. Figure 2.10 shows the pseudo-binary diagram were the horizontal axis is represented in terms of the beta stabilizer content in titanium alloys. An increase in the [Mo]eq decreases all the phase boundaries in the diagram, which means that the extent of beta phase field increases with the [Mo]eq at a given temperature. In α+β alloys with a relatively low level of β-stabilizing elements, quenching from high temperature above the β-transus results in martensitic phase transformation. The martensite phases are metastable in nature. They can be classified into two types: HCP-form (α’) and orthorhombic form (α’’). These martensitic phases can be converted into α+β phase mixtures by the subsequent ageing heat treatment. The formation of α+β phase mixtures, from the martensitie phases, can occur by [35]: β → (quench) → [α′ /α′′ ] → (age) → [α + β] β → (cool) → α + β → (quench) → α + [α′ /α′′ ] + β → (age) → α + [α +β] + β (2.3) (2.4) In the first transformation path (2.3), there is the complete martensitic transformation. The martensite is subsequently converted to α+β by ageing. In the second path (2.4), only part of beta undergoes martensitic phase transformation. This martensite then is transformed into very fine α+β, whereas the remaining β forms coarse α+β 17 microstructure without undergoing martensitic phase transformation. Thus, there can be a dispersion of very fine α+β produced from martensite in a coarse α+β matrix [35]. There is also a ω-phase formation between 10-30 % of [Mo]eq composition during the cooling of a β-rich alloy from high temperature. The ω phase has an ordered crystal structure and promotes slip localization, during tensile deformation. This degrades the ductility of beta titanium alloys [35]. 2.5 Metal Matrix Composites (MMCs) MMCs are typically made of a low density metal (such as Al, Mg, and Ti) matrix reinforced with ceramic fibers or particulates (e.g. SiC, graphite). The reinforcement generally increases the strength, stiffness, and resistance to wear and high temperature deformation of the composite. However, some of the shortcomings of MMCs include low ductility and fracture toughness relative to the metal matrix. Additionally, the cost associated with the manufacturing of high-performance MMCs can be high because the manufacturing process is complex. For example, complicated processing steps like precoating fibers and pre-melting the matrix are required to enhance its wetting characteristics of the MMC [41, 42]. MMCs are currently used in some military and aerospace applications. MMCs are preferred for components in missiles, jet engines, aircrafts, and the space shuttle where high specific strength and stiffness levels are desired. The specific strength is defined as the strength divided by density, which is also known as the strength-to-weight ratio. The specific stiffness can defined as the Young's modulus divided by density, which is also known as the specific modulus. The specific strength and stiffness values are mostly useful 18 to compare the relative merits of the MMCs. One of earliest application is boron fiber reinforced Al composite for covers for a missile guided system [2]. In commercial applications (automotive industry in Japan), titanium-based MMCs have been used in the Toyota diesel piston engine due to its superior wear and heat resistance [2]. The type of reinforcement in a MMC is very important because it controls the mechanical properties, cost of fabrication, and the overall composite performance. There are two major types of reinforcements: continuous and discontinuous reinforcements. Discontinuous reinforcements include particulate or whisker reinforcements (e.g. SiC, B4C, TiC, and TiB2 etc.), whereas continuous reinforcements include fibers (continuous or woven fiber mats) [43]. 2.5.1 Titanium Metal Matrix Composites (Ti-MMCs) Titanium composites exhibit a very good balance of mechanical properties and low density, making them attractive for lightweight structural applications, especially those involving high temperatures [44]. There have been several studies [45, 46] aimed at improving the mechanical properties of the titanium composites. Early research studies of Ti MMCs were focused on embedding continuous fibers as the reinforcement in Ti [47]. Continuous fiber reinforcements in Ti make the composite expensive because it is difficult to implement mass production of continuous fiber composite. In addition, continuous fiber composites exhibit limited formability and are hence difficult to deform or shape to manufacture different components [48]. In recent years, research is driven towards the development of a novel composite manufacturing processes that allow the reinforcement to form in-situ in the metal matrix, 19 either by the exothermic reaction or by crystallization during solidification [49]. This approach can help to produce homogenous composites in complex shapes along with the incorporation of thermodynamically stable reinforcements [49]. Titanium diboride was considered in the past [50] as a reinforcement for titanium alloys. However, titanium diboride is not stable due to the reactivity with the titanium, and forms titanium mono boride (TiB). This is because titanium boride is an intermediate phase between titanium and titanium diboride as evident from Ti-B phase diagram [5]. 2.5.2 Methods of Reinforcement In general, reinforcements such as SiC, TiC, TiN, B4C, Al2O3 ZrB2, and TiB2 may be employed to enhance the stiffness, strength, and hardness of titanium, in order to make it competitive in engineering applications. However, most of these reinforcements are not stable in the titanium matrix at the processing temperature due to high reactivity of titanium with reinforcements. For aerospace applications, SiC fibers were considered as reinforcement, leading to early manufacturing trials with Ti-SiC MMCs [51]. However, the reaction between Ti and SiC, during processing, leads to the formation titanium silicide (Ti5Si3) along the interfaces between matrix and reinforcement [51-53]. This leads to poor mechanical properties, due to the embrittlement of fiber-matrix interface and debonding. The SiC reinforcements are thermodynamically unstable and hence not stable. However, a study [54] showed that TiC particles can be thermodynamically stable due to the existence of TiC at wide carbon range in the phase diagram. A study [55] showed that the TiC particles are well bonded to the titanium matrix. However, this study also revealed the disadvantage of TiC phase, which is the occurrence of C, as interstitials in Ti, affecting 20 the ductility of the matrix [56]. Moreover, in another study [46], the transmission electron microscopy (TEM) showed an annular region in the TiC phase. In the annular region, there was a depletion of carbon due to the carbon diffusion to the matrix, or the titanium diffusion into the particles [57]. This inter-diffusion affects the homogeneity and stability of the TiC phases, and hence TiC is not preferred as reinforcement, although it is relatively stable in Ti matrix, compared to the other reinforcements. Figures 2.12 (a) and (b) represent the microstructure of Ti matrix reinforced by SiC and TiC, respectively [56]. For the metal matrix composite, it is important to select a reinforcement that is compatible with matrix. Compatibility here means that, in addition to chemical compatibility, the matching of the thermal expansion coefficient with that of the matrix is desirable. This expansion matching reduces the residual stress in the composite that can arise during processing and/or deformation. Previous studies [49, 58] showed that titanium boride (TiB) phase offers a good matching of the thermal expansion with the titanium matrix which makes it very attractive from a processing and application standpoint. 2.6 Structure and Properties of Titanium Boride The presence of titanium boride (TiB), in the form of whiskers embedded in the ductile matrix of titanium, offers several benefits in terms of mechanical properties. Titanium monoboride phase is more attractive than titanium diboride due to the absence of the intermediate phase between the titanium and titanium monoboride [50]. The TiB reinforcement has been found to increase the hardness, yield strength, wear resistance, and fracture strength [9] of titanium alloys. 21 Titanium boride can be formed in-situ through reaction sintering of titanium and TiB2 powders. The reaction leads to the formation of TiB as whiskers [58]. The shape of the whiskers can be in the form of needles, cluster, or short agglomerated whiskers. Calculation of elastic modulus through Halpin-Tsai (HT) equation [59] shows that the elastic modulus increases with the volume fraction and aspect ratio of TiB whisker [12]. The TiB phase as a reinforcement in the titanium offers many advantages as indicated below:  The density of the TiB phase (4.56 g/cc) is very close to that of titanium (4.51 g/cc) [60], which is advantageous, because as the volume fraction of TiB phase increases, the density will not change much.  The TiB reinforcement and the Ti matrix share a clean interface, as observed from bright field TEM micrographs [61]. This is due to absence of an intermediate phase between Ti and TiB as seen in Ti-B phase diagram, which is presented in the next section.  The elastic modulus of TiB is ~370 GPa and the Vickers hardness is ~1600 − 1800 kg/mm2 [60].  The whisker shape can enhance the mechanical properties such as stiffness, strength, hardness, wear resistance, and creep resistance due to transfer of load from the matrix to the whisker. The Halpin-Tsai plot in Figure 2.13 shows the enhancement in stiffness of Ti due to Vf and aspect ratio.  The thermal expansion coefficient of TiB (8.6 × 10−6 /°C) is close to Ti (8.2 × 10−6 /°C) [60]. 22 2.6.1 Ti-B Phase Diagram Titanium boride phases can exist in three forms: TiB, TiB2, and Ti3B4 in the Ti-B system. The Ti-B phase diagram, given in Figure 2.14, shows that the boron content of the TiB phase ranges from 18 to 18.5 wt. %. The Ti3B4 occurs at a composition of 22.4 wt. % B and TiB2 exists at 30.1 to 31.1 wt. % of boron. These phases are formed by different invariant reactions as shown in the phase diagram. TiB phase can form by peritectoid reaction, TiB2 by congruent solidification, and Ti3B4 by peritectic reaction. 2.6.2 Diffusion of Boron and Growth of TiB Whiskers Figures 2.15 (a) and (b) show the crystal structures of both TiB2 and TiB phases, respectively [62, 63]. The crystal structure of the TiB is orthorhombic and that of TiB2 is HCP. The crystal structures of TiB and TiB2 are built from the same basic unit. The basic unit is a trigonal prism for both TiB and TiB2. Six Ti atoms form the triangular prism, occupying the corners, with the B atom located at the center of the prism. However, the major difference between the crystal structures of TiB and TiB2 is the packing arrangement of the trigonal prisms. The TiB2 phase is formed by the vertical stacking of the trigonal prisms. The arrangement of B in TiB2 is in the form of hexagonal coordination of B atoms. This leads to B layers stacked alternatively between Ti layers. In case of TiB, the B27 structure is formed by the horizontal stacking up of triangular prisms, joining each other along the rectangular faces of the neighboring prisms. The growth of TiB phase occurs by inter-diffusion of Ti and B atoms. There are several studies [64-66] that have investigated the growth of TiB from a TiB2 coated layer or from a TiB2 cladding layer. The thermodynamic driving force of the formation of TiB 23 is favorable and TiB whiskers can be formed upon the reaction of Ti and TiB2 [5, 58]. The necessary condition for the formation of TiB is that there should be sufficient time for B diffusion, and the average B concentration of the Ti-B alloy should be less than 18-18.5 wt. % in the reaction zone. The formation of TiB occurs by a simple reaction: Ti + TiB2 → 2TiB (2.5) Most of these studies [66, 67] investigating the formation of TiB were done using titanium as a substrate such that titanium boride layer forms on the substrate. For example, in the study done by Galvan et al. [67], the formation of TiB phase on Ti-6Al-4V substrate was enabled during the process of laser cladding. A powder mixture of TiB2/Ti was used to form a laser cladded layer. The TiB2/Ti powder was injected into laser beam. A stream of TiB2/Ti powder particles were injected around the Ti of the substrate to form TiB layer [67]. TiB whiskers formed in the subsurface region of the substrate. In addition, a study [68] demonstrated laser boronising technique to form boride coating on titanium substrate by using B powders. This study showed that it is possible to form TiB phase, without the formation of TiB2 phase, on the titanium substrate. The first study that showed that bulk Ti-TiB composite can be made by simple reaction is by Sahay et al. [5]. The study indicated that reaction sintering can be used to form Ti-TiB composites containing various volume fractions of TiB. The boron diffusion in the [010] TiB direction is much faster compared to other crystallographic directions due to the continuous B chain. Due to this reason, TiB tends to form a whisker morphology, with the whisker group in [010] direction. In addition, several studies [64, 69] have indicated that it is possible to form TiB whiskers with a certain orientation relationship with alpha-titanium. The orientation relationships between TiB and α-Ti are: 24 (100)TiBII(0001)α-Ti ; [010]TiBII[112̅0] α-Ti (2.6) The growth of the TiB whiskers takes place parallel to the axis of whiskers, where the rate of B diffusion is higher than the rate of Ti diffusion in the opposite direction [5, 64]. The diffusion coefficient of B in TiB phase has been reported to be about 45 times larger than the diffusion coefficient of B in TiB2 phase [64]. Moreover, the growth rate of TiB whiskers is also enhanced by the chain like arrangement of the B atoms and dense BB bonds in the [010] direction. The growth rate of these whiskers is higher in direction [010] compared to [101], [100], and [001] directions [69]. 2.7 Processing Techniques for Ti-TiB Composites Ti-TiB composites can be processed by different techniques such as rapid solidification processing [70], conventional ingot processing [71], and powder metallurgy techniques [49]. The study by Ranganath [72] was probably the first that showed that TiB and TiC can be formed inside the titanium matrix by combustion synthesis. Among the possible techniques, powder metallurgy is very attractive, due to the low cost of processing [73]. The powder metallurgy method can produce Ti-TiB composites with desirable microstructures, which are difficult to produce by other methods such as melting [58]. In early studies [5, 58], Ti-TiB composites were obtained by reaction sintering in a hot press. In this process, the TiB phase forms in Ti matrix by solid state reaction. The simple reaction sintering can occur and form the TiB whisker on the Ti matrix during densification of the composite itself. Optimization of the initial powder size, compositions, and process conditions (temperature and pressure) can lead to optimum composite microstructures. The 25 Ti-TiB composite processing in a hot press is relatively more time consuming due to the heating, holding, and cooling periods involved. The Electric-Field-Activated-Sintering (EFAS) is similar to the hot pressing where the powders are loaded inside the die and are then heated using a direct current under a certain ram pressure. The advantage of EFAS lies in direct Joule heating of Ti+TiB2 powders, and this is facilitated by the electrically conductive powder mixture. This enables rapid heating, reaction, and densification of the composite. The EFAS can provide rapid heating to the reaction sintering temperature. Because of Joule heating, the reaction sintering is possible at relatively low temperatures [6]. During the reaction sintering of Ti+TiB2 powder mixture, several reactions occur due to the exothermic nature of the overall reaction [58]. In the Ti-TiB composite formation, the three important reactions are: Ti + B → TiB (2.7) Ti + 2B → TiB2 (2.8) Ti + TiB2 → 2TiB (2.9) Figure 2.16 shows the Gibbs free energy change, as a function of temperature, for the three reactions [58]. All the reactions are thermodynamically favorable with negative ∆ 𝐺 values, but the second reaction (2.8) is thermodynamically most favorable. It is evident that Ti and TiB2 can react and form TiB, but the thermodynamic driving force of the reaction is relatively small due to a relatively small value of -∆ 𝐺. However, since ∆ 𝐺 is still negative, this means that TiB2 can be used as a source of B, and it can react with Ti to form TiB as long as the amount of B restricted to less than 18 to 18.5 wt. % [5]. 26 During sintering, the reaction can be completed quickly and densification can be rapid if a liquid phase forms. Liquid phases can lead to full densification of the composite by completely eliminating the pores [9, 6]. In a prior study [49], it was found that the addition of iron or molybdenum can lead to the formation of the liquid phase during sintering. 2.7.1 Electric-Field-Activated-Sintering (EFAS) Processing of Ti-TiB Composition EFAS is a relatively new sintering technique that can produce nano-crystalline materials with good a control over the extent of grain growth. In this technique, Joule’s heating (H = I2 Rt, where H is heat in Joules (J), I is current in Amperes (A), R is resistance in Ohms (Ω), and t is time in secs is used to create melting at particle contacts enhance the densification of the powder [74]. The EFAS has the advantage of densifying powders in nano-size at the same time avoiding grain growth, which can occur excessively in conventional sintering techniques methods. Several ceramics (e.g. MgAl2O4, Al2O3 etc.) that are of interest in thermoelectric, optical, magnetic, magnetoelectric, piezoelectric, or biomedical fields have been successfully made by EFAS. In EFAS technique, the Ti-TiB composites can be produced at relatively low temperatures (850-1350 °C), with full densification, and in a short time compared to hot pressing [6, 9]. It is important to note that by reducing the processing time, the cost of production can be reduced. Alternatively, higher heating rates can be achieved by passing relatively higher currents, which accelerates the formation of incipient melting zones at particle contacts [75]. 27 2.8 Microstructures of Ti-TiB Composites In Ti-TiB composites, the TiB phase is hard, and the Ti (α or β) phase is relatively more ductile and less stiff. In reaction sintering, the starting powder size composition and packing play important roles in affecting the microstructure of the composite [58]. The formation of TiB whiskers is affected by the packing density of the starting powders [58]. In the powder mixture, the TiB2 particles generally occupy the interstices between Ti particles due to their smaller size. During the growth of TiB whiskers by B diffusion, the adjacent whiskers growing from neighboring TiB2 particles can interfere with the growth of a TiB whisker [5]. The ratio of the bulk density of the powder to the true density of the material is defined as the packing density [58]. There are two major types of packing density used for effective packing of the powders: bimodal and trimodal packing. In the bimodal packing, the desired powder size ratio is about 20:1 or for an effective packing density of 85 %. In the bimodal packing, the TiB whiskers generally were found [58] to form around the beta grains due to the large powder size ratio. For the trimodal packing, the desirable powder particle size ratio is 49:7:1, which leads to a packing density of 92 % [58]. In the trimodal packing, TiB whiskers were found to grow in the matrix uniformly with high aspect ratio [58]. Figure 2.17 shows the effect of TiB2 powder arrangement in the microstructure of a typical Ti-TiB composite [5]. The diagrams depict how the whisker forms through the stages from (a) through (c), for the low volume fraction of TiB, and from (d) through (f), for the high volume fraction of TiB. The arrangement of TiB2 powder controls the morphology of the TiB whisker formed. In Figure 2.17 (a), the titanium particles are 28 partially surrounded by TiB2 powder particles, which are one-tenth of the size of the titanium particles. Specifically, the TiB2 particles will occupy the interstitial spaces between the titanium particles when their size is much small compare to the Ti particles. When the amount of TiB2 in small (sufficient to form ≤ 30 Vf of TiB), the TiB whisker can grow through a relatively large mean-free-path. This allows unhindered whisker growth, leading to high aspect ratio TiB whiskers, until the completion of the decomposition of the TiB2 or the unit is obstructed by a neighboring already grown whiskers. This explanation was provided [5] for the microstructure evolution in Ti-30TiB composite, with high aspect ratio of the TiB whiskers, as shown in Figure 2.18 (a). For the case of high volume fraction TiB composites, the schematics (d) through (f) illustrate the mechanism of TiB formation [5]. Here the titanium particles are completely surrounded by TiB2 particles. In this case, the mean-free-path is relatively less compared to the situation illustrated in (a) through (c). The TiB whiskers will grow simultaneously from many TiB2 particles and this growth will lead to whisker impeding each other early, because of the reduced interparticle distance. This makes the TiB whisker grow with a relatively low aspect ratio compared to the first case. Ti-TiB composites with high volume fraction of TiB whiskers; such as Ti-55TiB, Ti73TiB, and Ti-72TiB–8TiB2, form by stages shown in Figures 2.18 (b), (c) and (d). However, there is also the possibility of new growth of TiB whisker in its transverse direction when the TiB whisker growth in the axial direction is stopped by a neighboring TiB whisker. Since B diffusion in the transverse direction ([100], [101], and [001]) of TiB is much slower than that in the longitudinal direction ([010]) of TiB, the new whiskers may not form as one might expect [5]. 29 2.9 Estimation of TiB Volume Fraction For the Ti metal matrix composites, it is important to quantify the details of the microstructure and relate these details to the mechanical and physical properties of the composite. These details include the distribution of the phases, size, shape, volume fraction, etc. The volume fraction can be determined from the optical micrographs by stereology. This method is difficult to apply in the Ti-TiB composite due to the poor contrast between the phases (Ti and TiB). The X-ray method is the most accurate technique to determine the phase volume fraction in the Ti-TiB composite. The direct comparison method can provide a relatively more accurate for determination of phase volume fractions in composites [76]. The volume fraction of each phase in Ti-TiB composites can be calculated by the direct comparison method. It is important to select non-overlapping peaks to get an accurate value of the volume fraction. In this method, the volume fraction can be calculated by: Vf = R RTi ITiB Ti ITiB + RTiB ITi (2.10) where I is the integrated intensity of the target peak and R is a parameter that can be calculated by equation (2.11): R= |Fhkl |2 PL Vₒ (2.11) where Fhkl is the structure factor, p is a multiplicity factor, L is the Lorentz polarization factor, and Vₒ the volume of the unit cell. Figure 2.19 shows the X-ray diffraction patterns for several compositions of Ti-TiB composites as determined by Sahay et al. [5]. The X-ray pattern of pure Ti processed by hot pressing in the same way is also included. From the X-ray patterns, it is clear that only two phases Ti and TiB are presented in composites: Ti-30TiB, Ti-55TiB, and Ti-73 TiB. 30 As seen in the X-ray pattern when the TiB peaks [(200)TiB, (201)TiB, (210)TiB, (102)TiB, and (312)TiB] increase in intensely the peaks of Ti [(002)Ti, and (101)Ti], decrease. This indicates that the relative amount of TiB volume fraction increases with decreasing α-Ti phase volume fraction. In the composite with the high volume fraction TiB (Ti-86TiB6TiB2), TiB2 phase is present as shown by (001)TiB2, (100)TiB2, and (101)TiB2 peaks. The composite TiB-8TiB2 consists of only two phases: TiB and TiB2 [5]. The relative volume fraction of all the phases (Ti, TiB, and TiB2) were calculated using the integrated intensities by the direct comparison method [76]. The number before the TiB in the Ti-XTiB composition represents the value of the volume fraction of TiB. 2.10 Mechanical Properties of Discontinuous Reinforcements in MMC In MMCs, two types of discontinuous reinforcements can exist, that is, particulates or whiskers. The most common types of particulate reinforcements are SiC, B4C, TiC WC, and Al2O3. For the whiskers, the most common type is SiC and Al2O3. In general, either of these reinforcements can provide high specific stiffness, strength, hardness, and wear resistance. However, the strength and stiffness are enhanced at the expense of ductility and fracture toughness. The improvement in strength and stiffness largely comes from the load transfer from the matrix to the reinforcement [77]. 2.10.1 Strength and Stiffness of Ti-TiB Composites Godfrey et al. [78] investigated the elastic modulus for the Ti-TiB composites. This study focused mainly on low volume fraction TiB (<0.2) composites as these composites generally have better ductility than those with higher Vf of TiB. 31 Figure 2.20 shows the effect of the volume fraction of TiB on the elastic modulus of Ti-TiB composites as compiled by Godfrey et al. [78]. From the graph, it can be seen that as the volume fraction of the reinforcement increases, the Young's modulus of the composite increases. The rule-of-mixtures prediction of modulus is given by the dashed line. The graph also shows predictions from the Eshelby model, by varying the aspect ratio of the whisker. The Eshelby approach is a mathematical approximation used to determine the stresses occurring due to the misfit between the shapes of the constituent (matrix and reinforcement) in a composite, which allows the prediction of elastic properties of the composite for various volume fractions and aspect ratios of the reinforcement. This methodology uses the concept of “an equivalent homogenous inclusion” with an equivalent transformation strain so that it allows the stress field to be same as that of an actual inclusion; this technique is found to be accurate for predicting a wide range of composite properties [79]. The Eshelby approach [79] explains well the elastic behavior of the composites having TiB whisker with random orientation that are produced by powder metallurgy [79]. Figure 2.21 shows the tensile elastic modulus (obtained from tensile tests) and dynamic elastic modulus data plotted as a function of volume fraction of TiB within different Ti-TiB composites [12]. It is evident that as the volume fraction of TiB increases, the elastic modulus also increases. The graph illustrates the predictions from the HalpinTsai (HT) equation for two aspect ratios. The trends are calculated from the aspect ratio of the TiB whisker with respect to the shape of the whisker. The whiskers are considered as random oriented short fibers with an aspect ratio between 1 to 20. The graph shows reasonable agreement between the experimental results and the predictions by the HT 32 model. The dynamic elastic modulus data increase with increase in the volume fraction of the TiB in the Ti-TiB composite. Significant changes occur in the matrix due to the reinforcement in the MMC. Dislocations are generated in the matrix by the solutionizing and quenching process, due to the mismatch in the coefficient of thermal expansion between the matrix and the reinforcement. The strength of MMC can be increased by such dislocations. Additionally, dislocation generation can occur during deformation and this depends on particle volume fraction, particle size, and matrix strength [3]. Figure 2.22 shows the effect of volume fraction of TiB on the increase in the levels strength of the composites [78]. It is evident that an increase in TiB volume fraction can provide significant increase in the strength and stiffness of the Ti-TiB composite. Ductility and toughness of the MMC are generally reduced with increase in the volume fraction of the reinforcement, due to restriction of the space for dislocation movement, arising from the reduction in matrix volume fraction. Figure 2.23 shows the effect of volume fraction of the TiB on the ductility of Ti-TiB composites [78]. It can be seem that when the volume fraction of the TiB is about 10 %, the composite ductility is about 5-7 %. When the TiB volume fraction reaches ~30 %, the ductility is almost zero. This decrease in ductility is also made worse by the embrittling effects of oxygen and nitrogen in titanium on the MMC. Ductility of Ti-MMCs can be only increased by either increasing the intrinsic ductility of α-phase, or by introducing β-phase. 33 2.10.2 Flexure Strength and Fracture Toughness The flexural strength and ductility or toughness of titanium composites with TiB reinforcements have been studied by some investigations. In the study conducted by Feng et al. [80], the effect of sintering temperature and TiB volume fraction on flexural strength was evaluated by three-point bending. Figure 2.24, adapted from [80], shows the variation of relative density of the composite as a function of sintering temperature for Ti-TiB composites with 10 vol. % TiB. From the plot, it is seen that the relative density is only at 800 °C. On increasing the sintering temperature to 1000 °C, relative density exceeding 99 % was achieved. Further increase of sintering temperature to 1200 °C actually decreased the relative density substantially. According to the paper [80], debonding decreases the relative density of the composite, which is not theoretically possible if the sintering is conducted properly. According to the study [80], at an optimum temperature of 1000 °C, the in-situ TiB particles grow large enough to strengthen the titanium matrix and share the load effectively at this temperature. Figure 2.24 shows that the sintering temperature can affect density and volume fraction of TiB by two different mechanisms [80]. For density, an increase in sintering temperature should lead to better densification by enhanced plastic flow or liquid phase formation during sintering. For reaction to form TiB, an increase in temperature would be expected to increase the kinetics of TiB formation, thus increasing the Vf of TiB. While the data in Figure 2.24 are consistent with second mechanism, they are not consistent with the first mechanism. Figure 2.25, also taken from the same study [80], shows the effect of sintering temperature on volume fraction of TiB and on the flexural strength, although these two 34 factors can influence the flexural strength independently. The flexural strength reaches a maximum value of 1560 MPa at a sintering temperature of 1000 °C, which roughly corresponds to 10 vol. % TiB reinforcement. However, the data also show that further increase in the volume fraction of TiB results in a decrease in the flexural strength. The reason for this unexpected trend is not mentioned in the study, but this can be attributed to poor densifications. For the Ti-TiB composite, fracture toughness is an important parameter because plastic deformation and energy absorption, before fracture, will increase the reliability in application. The fracture toughness in such composites is often controlled by various factors like elastic modulus, flow stress, and the extent of plastic deformation associated with the metal matrix [81]. Figure 2.26 illustrates the variation of fracture toughness as a function of volume fraction (5, 10, 15, and 20 vol. %) of TiB for sintering temperature of 1000 °C [80]. It is seem that a maximum fracture toughness of ~12 MPa√m was attained after sintering at 1200 °C. However, this does not necessarily correspond to an increased formation (Vf) of TiB, because the same figure shows that the fracture toughness is decreased at high Vf levels of TiB. The study attributed this observation to the rapid growth of in-situ TiB needles due to the high temperature of sintering [80]. In addition, from the microstructure analysis of the composite sintered at 1200 °C, the occurrence of massive faceted TiB needles on the fracture surface was observed. 35 2.11 CALPHAD Approach for Alloy Design In traditional physical metallurgy approach, alloys are designed to form solid solutions or second phases for strengthening and toughening alloy. This alloy design approach is largely empirical, but enabled metallurgists to develop materials with superior mechanical properties. This approach benefited development of alloys with low and high temperature strength, resistance to creep, oxidation resistance, corrosion, and wear resistance. The solid solutions were designed based on a rigorous trial and error approach of synthesis so as to optimize the mechanical properties required for a particular engineering application. This solid solution strengthening does not affect the structureinsensitive properties much (e.g. E, G, K) but it enhances the structure-sensitive properties (e.g. strength, ductility, and toughness) [82]. Alloy properties and their behavior in applications are dictated by microstructure, which is controlled by the microstructure and phase details. Phase diagrams are very useful for predicting which phases can be expected for given conditions (temperature, pressure, overall composition of the alloy). Traditionally, the phase diagrams are constructed experimentally by determining the stability of a phase (minimum Gibbs free energy) at various temperatures [82]. However, using the thermodynamic properties (enthalpy, and entropy) of phases, the Gibbs free energy can be calculated theoretically; when this is done for various compositions and temperatures, it can lead to the constructions of a theoretical phase diagram. Such phase diagrams can be constructed using the CALPHAD method, which is the calculation of phase diagrams by computational methods. Recently, CALPHAD has emerged as one of the main computational method to determine the equilibrium phase diagram, especially for ternary compositions. The CALPHAD approach 36 uses a semi-empirical method using the thermodynamic data of multicomponent systems. This CALPHAD approach is powerful since it allows the construction of phase fields in complex multicomponent systems, not studied experimentally, by extrapolation of the thermodynamic properties. The most common programs following CALPHAD approach are MTDATA, FactSage, PANDAT, JMatPro, MatCalc, and Thermo-Calc. CALPHAD approach can save time and money by experimental work involved in constructing the phase diagrams. However, it should be emphasized that any phase field predicted by CALPHAD approach must be experimentally validated to ensure that it is reliable. The equilibrium Ti-Co phase diagram in Figure 2.27 indicates the presence of the three stable phases αTi, βTi, εCo, and αTi. Apart from these four phases, the phase diagram also shows the presence of TiCo2 (hexagonal), TiCo2 (cubic), TiCo (cubic), and Ti2Co (cubic) [83]. These phases are intermetallic and highly ordered. There is also occurrence of two eutectics and four peritectics in this system at 1020, 1170, 1058, 1235, 1190, and 1215 °C, respectively. This phase diagram is calculated by determining the phase field of each phase by evaluating the driving force (Gibbs free energy) behind formation of each phase using a cooling curve at different temperatures. CALPHAD determines the phase field of each composition with the minimization of Gibbs free energy at all possible temperatures and thus the equilibrium phase diagram is generated. In a previous study [9], CALPHAD approach was used to simulate the phase diagram of Ti-B-Fe-Mo system. The phase diagram was constructed by extrapolation of the all available experimental thermodynamic data derived from its three ternary systems of Ti-B-Fe, Ti-B-Mo, Ti-Fe-Mo, and B-Fe-Mo. In essence, CALPHAD approach was used for a comprehensive thermodynamic assessment using extrapolation of the thermodynamic 37 parameters like Gibbs free energy, entropy, enthalpy, activity of the constituents etc. of the binary systems into ternary systems, and finally to the quaternary system [9]. 38 Figure 2.1 Titanium crystal structure. (a) Alpha-titanium as HCP crystal structure at room temperature. (b) Beta-titanium as BCC crystal structure [15]. 39 Figure 2.2 Pure titanium. (a) Slip planes [15]. (b) Twin planes, four color planes are twin planes. The yellow and the light green indicate twin plans {112̅1} and the pink and the green refer to twin plans {101̅2}. The arrows indicate the corresponding twinning directions shown in Table 2.1 [23]. 40 Figure 2.3 Tensile properties of CP-titanium with varying oxygen, nitrogen, and carbon contents [24]. 41 Figure 2.4 Binary phase diagram of titanium and hydrogen [29]. 42 Figure 2.5 Effect of hydrogen on the impact strength of titanium [32]. 43 Figure 2.6 Effect of oxygen on the transformation temperature of titanium alloys [33]. 44 Figure 2.7 Effect of carbon on the transformation temperature of titanium alloys [33]. 45 Figure 2.8 Effect of nitrogen on the transformation temperature of titanium alloys [33]. 46 Figure 2.9 Titanium alloys classification based on phase stability [37]. 47 Figure 2.10 Effect of beta stabilizing elements on the titanium [15]. 48 Figure 2.11 Effect of aluminum on transition temperature of titanium of phase diagram [39]. 49 Figure 2.12 Microstructures of (a) Ti-SiC composite [41] (b) Ti-TiC composite [56]. 50 Figure 2.13 Elastic modulus as a function of TiB Vf for Ti-TiB composites. 51 Figure 2.14 Binary phase diagram of titanium and boron constructed by ThermoCalc alloys data base. 52 Figure 2.15 Crystal structure of (a) TiB2 and (b) TiB phases [62, 63]. 53 Figure 2.16 The formation of free energy in Ti compounds [58]. 54 Figure 2.17 The growth of TiB phase (whisker) during the sintering process (a) through (c) low volume fraction of TiB as a product and (d) through (f) process of high volume fraction of TiB [5]. 55 Figure 2.18 Composites (a) Ti–30TiB (b) Ti–55TiB (c) Ti–73TiB (d) Ti-72TiB– 8TiB2 [5]. 56 Figure 2.19 X-ray pattern of the Ti-TiB composite sintered with different temperatures and different volume fraction of TiB [5]. The numbers before TiB in the data caption indicate the vol. % of the phase. 57 Figure 2.20 Effect of the volume fraction of the TiB on the Ti-TiB composite on the elastic modulus [78]. 58 Figure 2.21 The elastic modulus vs. the volume fraction of the TiB in the Ti-TiB composites [12]. 59 Figure 2.22 Effect of TiB as volume fraction on the strength of the metal matrix composite [78]. 60 Figure 2.23 Effect of volume fraction of the TiB on the ductility of the metal matrix composite [78]. 61 Figure 2.24 Relative density variation with sintering temperature and volume fraction of TiB phase [80]. 62 Figure 2.25 Effect of volume fraction of TiB and sintering temperature on the flexural strength of composite [80]. 63 Figure 2.26 Effect of volume fraction of TiB and sintering temperature on fracture toughness of composite [80]. 64 Figure 2.27 Binary phase diagram of Ti and Co constructed by CALPHAD [83]. 65 Table 2.1 Slip and twinning systems in α-Ti Slip system Slip Plane & Direction Twinning system {0001̅} <112̅0> Basal plane, Basal direction {112̅1} <112̅6̅> {101̅0} <112̅0> Prism plane, Basal direction {101̅2} <101̅1̅> {101̅1} <112̅0> Pyramidal plane, Basal direction {112̅2} <112̅3̅> {112̅2} <112̅3> (c+a) slip CHAPTER 3 MATERIAL DESIGN AND EXPERMENTAL PROCEDURE 3.1 Design of Composite Composition and Processing Titanium metal matrix composites with high volume fraction TiB reinforcements are very challenging to design due to the lack of information about the right compositions that are processable and may provide good mechanical properties. The selection of correct process temperatures and pressures is important to achieve highly densified composite with homogenous distribution of the TiB phase. In this study, six compositions of Ti-MMCs were selected. Three compositions were selected based on the variation of B content by keeping the Fe content fixed to evaluate the effect of B on the microstructure, and another three compositions by varying the amount of Fe with constant B content to evaluate the effect of Fe on the microstructure. In the Ti-B-Fe system, Ti powder reacts with the TiB2 in the presence of Fe powder to produce TiB and β-Ti. As shown in the following reaction: Ti + TiB2 + Fe → 2TiB + Ti (β) (3.1) Fe promotes the formation of β-phase at room temperature, and the reaction sintering can be facilitated since Fe reduces the sintering temperature. The MMCs compositions of the synthesized materials processed by Electric-Field-Activated-Sintering (EFAS) are given in Table 3.1. 67 3.2 Phase Equilibria in Ti-Fe System In processing of Ti-B-Fe by EFAS, the formation of a liquid phase is desirable to accelerate the reaction sintering of Ti and TiB2 powders and improve the densification. For this purpose, Fe is desirable as alloying addition. The Ti-Fe experimental phase diagram (Figure 3.1) shows a low temperature eutectic reaction (~1085 °C), at the composition Ti with ~28 mol. % Fe. From a material design standpoint, a substantial amount of alloying by Fe, in the composites, is needed to achieve 30-50 vol. % of metallic β-Ti phase. Fe is a beta stabilizer and the objective is that the addition of Fe can stabilize a potentially ductile β-Ti phase in the composite. The experimental (ASTM Handbook) Ti-Fe phase diagram [84] and the CALPHAD calculated phase diagram are presented in Figures 3.1 and 3.2, respectively. The phase diagram calculated by using the CALPHAD approach is in good agreement with the experimental phase diagram. However, there are a few minor discrepancies between them. For example, the γ-Fe loop in the experimental Ti-Fe system is absent in the CALPHAD calculated phase diagram. The calculated phase diagram showed the five invariant reactions where the degree of freedom in the Gibbs phase rule is zero: (1) Liquid ↔ α-Fe+Fe2Ti at ~1289 °C. (2) Liquid ↔ Fe2Ti at ~1427 °C. (3) Liquid + Fe2Ti ↔ Fe2Ti + FeTi at ~1317 °C. (4) Liquid ↔ FeTi + β-Ti at ~1085 °C. (5) β-Ti ↔ FeTi + α-Ti at ~595 °C. The processing temperature of the composites was selected at 900 °C due to the absence of liquid phase, which allows solid state sintering. The Fe is chosen between 10- 68 30 mol. % during the sintering of the Ti-TiB composites owing to the fact that in this region (as seen from Figure 3.2), there is no formation of intermetallic phase of Fe2Ti. Although there is some amount of formation of TiFe > 20 mol. % Fe, this intermetallic TiFe is formed only in trace amounts since most of the Fe gets incorporated into the beta titanium, thereby stabilizing the matrix phase. 3.3 Materials and Experimental Approach The raw powder materials used in this study were (a) 99.95 % pure Ti (α-Ti) powders with an average particle size and O content (in wt. %) of 32 µm and 0.109 (supplied by Puris, Bruceton Mills, WV), respectively, (b) TiB2 powders with an average particle size of 14 µm with impurity content (in wt. %) of 0.50 O, 0.015 Zr, 0.50 C, 0.02 Fe, and 0.20 N (supplied by Momentive, Waterford, NY), and (c) 99.5 % pure iron powder with an average particle size of 6 µm and impurity content (in wt. %) of 0.20 O and 0.02 C (supplied by Thermo Fisher Scientific, Ward Hill, MA). The sizes of the powder were selected based on trimodal packing density calculations in a previous study [58], in order to achieve the highest packing density. The size ratio of powders used in this study was 32:14:6. The powders were blended in a ball mill containing titanium balls in order to have homogenous mixture. The ball milling containers were sealed with Ar to reduce oxygen pick up during milling. The ratio of the weights of powder to titanium ball was 15:1 in wt. %. The blending was done at 350 rpm for 24 hours. In this study, a 10-ton Electric-Field-Activated-Sintering (EFAS) (Model 25-10, Thermal Technologies, Santa Rosa, CA) unit was used to process bulk Ti-TiB composite with different volume fractions of TiB phase. The reaction sintering temperature in the 69 EFAS process was maintained in the range of 900-1100 °C. The process parameters for the sintering process included a constant pressure of 10 MPa, a heating rate of 50 °C/min, a holding time of 4 hours at the sintering temperature, and a constant cooling rate of 22.5 °C/min until the sample cooled to room temperature. These process parameters were kept constant for all the composition to evaluate the effect of B/Fe content variation on the sintering reaction and densification of the MMCs. Bulk densified samples of 60 mm in diameter and 14 mm in thickness were synthesized by EFAS. The final densities of the composites were measured by Archimedes principle. The hardness measurements were made using a Vickers Micro-hardness tester (Leco, M-400), with 1 kg load. At least twenty random indents were done on the surface and the sixteen most uniform indents were selected. 3.4 Microstructure and Phase Characterization The phases in the MMCs were identified by X-ray diffraction (XRD) analysis (Rigaku, Miniflex 600). XRD pattern was acquired with a scan rate of 0.5 degree/minute. The microstructures of the composites were investigated by an optical microscope (Leco, Olympus GX51) on MMCs. For optical microscopy, the samples were metallographically polished and etched by Kroll’s etchant, containing 96 ml H2O, 3 ml HNO3, and 1 ml HF. The fracture surfaces of some of the composites were analyzed by scanning electron microscope (SEM, (FEI Quanta 200, USA)) equipped with an energy-dispersive X-ray (EDS) detector. 70 3.5 Tensile Testing Bars for tensile specimens were cut by Electrical-Discharge-Machining (EDM). They were then machined to their final sizes according to the ASTM standard designation E8/E8M-15a for tensile testing specimens [85]. Tensile specimens had a diameter of 3.175 mm, with a total length of 50.80 mm and a gage length of 12.70 mm. The applied strain rate during tensile testing was 2 × 10−4 /s. The tensile test were performed using a Material Test System (MTS 810). During the tests, an extensometer was used to measure the strain. 3.6 Flexure Testing For the flexural strength measurement test, the specimen preparation involved cutting specimen blanks first by using EDM. The dimensions of these blanks were approximately 45 mm in total length, 4 mm in width, and 3 mm in thickness (3 × 4 × 45 mm3). Each specimen was polished in order to remove the affected layer left from the EDM by using SiC polishing paper starting from 320 grit and going down to 1200 grit, followed by final polishing with colloidal silica (0.08 µm). This procedure give a very high level of polishing. This polishing is intended to remove any scratches or pits from EDM so that failure will not initiate from such defects. The final dimension of the specimens were conformed to the ASTM standard following the C1161-13 designation [86]. Four-point flexural testing was conducted with a loading rate of 0.5 mm/min at room temperature using the polished samples, in MTS 810. 71 3.7 Fracture Toughness Testing For the fracture toughness test, single edge pre-cracked beam (SEPB) method was followed as per ASTM standard C1421-10 [87]. The specimen preparation procedure and the dimensions are the same as that of flexural strength testing. The pre-crack was induced on the specimens by using a bridge compression fixture. Three Vickers indents on the tension side served as the crack initiations. The indentations were made in an MTS 810 unit equipped with a Vickers indenter and with 100 N indentation load. The pre-crack length was observed with the aid of optical microscopy using dye that revealed the length the pre-crack on fracture surface. The length of the crack was determined as the average of three measurements measured at 25 %, 50 %, and 75 % of the thickness of the specimen. 3.8 Estimation of TiB Volume Fraction In a multiphase composite, the volume fraction of each phase can be determined using quantitative metallography, by applying the point-count method. However, this method requires a good contrast between the phases to get an accurate estimation of volume fraction. In the Ti-B-Fe system, four kinds of phases can be present such as Ti, TiB, TiB2, and TiFe. These phases have poor contrast relative to each other under the optical microscope, which makes the determination of the volume fraction of the phases difficult. Alternatively, the direct comparison method [76] can be applied to bulk multiphase materials to calculate the volume fraction of constituent phases. This method gives an accurate value of the volume fraction for Ti-TiB composites [5]. This method utilizes integrated intensity value from the X-ray pattern for quantitative estimation of volume fraction. In this method, the integrated intensity is first determined for the selected peaks 72 (non-overlapping peak) of Ti or TiB phases. The volume fraction of the TiB phase was calculated as: Vf = R RTi ITiB Ti ITiB + RTiB ITi (3.2) where I is the integrated intensity of the selected (hkl) peak and the R is an angulardependent parameter, which normalizes all the structure factor, multiplicity factor, and Lorentz polarization factor. The parameter R was calculated by the equation (3.3): R= |Fhkl |2 P L Vₒ (3.3) where Fhkl is the structure factor, p is multiplicity factor, L is Lorentz polarization factor, and Vₒ is the volume of the unit cell. By knowing the integrated intensities and R-values, the relative volume fraction of the β-Ti and TiB phases were calculated by this method. 73 Figure 3.1 Titanium iron binary phase diagram [84]. 74 Figure 3.2 Titanium iron binary phase diagram calculated from the thermodynamic data using Thermo-Calc. 75 Table 3.1 Compositions of the synthesized MMCs Composition (mol. %) Type Name Ti B Fe Ti-10B-10Fe 80 10 10 Ti-20B-10Fe 70 20 10 Ti-30B-10Fe 60 30 10 Ti-10B-10Fe 80 10 10 Ti-10B-20Fe 70 10 20 Ti-10B-30Fe 60 10 30 Ti-20B-20Fe 60 20 20 Constant Fe varying B Constant B varying Fe CHAPTER 4 RESULTS AND DISCUSSION 4.1 Ternary Ti-B-Fe System In EFAS processing, it is first required to fix the process temperature and this requires an understanding of the phase fields in Ti-B-Fe system. However, an experimental Ti-B-Fe phase diagram is not available. Constructing such a phase diagram for the composition range of interest here will be difficult and requires a more extensive effort, which is beyond the scope of this study. However, using CALPHAD approach, the phase fields in any ternary system can be calculated using thermodynamic data. The isothermal section of the Ti-Fe-B ternary phase diagram calculated by CALPHAD is shown in Figures 4.1 (a-c). This sections shows the phases present at 30, 700, and 900 °C. At 900 °C isothermal section, the compositions Ti-10B-10Fe and Ti-20B-10Fe existed within the phase field of β-Ti+TiB, which indicated the stability of two-phase (β-Ti and TiB) microstructures. On the other hand, Ti-30B-10Fe, Ti-20B-20Fe, Ti-10B-20Fe, and Ti-10B30Fe were within the three phase field (β-Ti+TiB+TiFe), indicating the stability of three different phases: β-Ti, TiB, and TiFe. It is to be noted that β-Ti and TiB are desirable phases here. TiFe is an undesirable intermetallic compound. 77 4.2 High Temperature X-ray Diffraction Study XRD data are helpful for the determination of phase boundaries (for example solvus lines), volume fraction of the phases, and the crystal structure of the phases. In order to investigate the phase fields of a system at a given temperature, usually, the alloys are equilibrated at that temperature and then quenched down to room temperature to avoid any phase transformation of the phases formed at high temperature. This technique may not be consistent or effective all times. Hence, in-situ XRD (or HT-XRD) is very useful in cases where the material at high temperature cannot be quenched for various reasons. However, this approach also requires judicious consideration of oxidation and melting, which may influence the X-ray results. In order to validate the CALPHAD calculations, it was necessary to perform HTXRD. The experiments here were performed on the MMCs at room temperature, 700 and 900 °C. The results are shown in the Figure 4.2 (a-f). The CALPHAD calculation of ternary section at 900 °C shows the presence of two stable equilibrium phases of β-Ti and TiB for the composites Ti-10B-10Fe and Ti-20B10Fe, which is consistent with the experimental HT-XRD pattern. In addition, for the composites Ti-30B-10Fe, Ti-10B-20Fe, Ti-10B-30Fe, and Ti-20B-20Fe, the CALPHAD calculations indicate the presence of three phases of β-Ti, TiFe, and TiB, which again is in agreement with experimental HT-XRD. The HT-XRD patterns are in good agreement with CALPHAD calculations since equilibrium conditions are generally easy to maintain at higher temperatures of 900 °C, but it becomes difficult at lower temperatures due to slow kinetics of transformation. However, at lower temperature of 700 °C and room temperature, CALPHAD calculations show the presence of α-Ti phase field, which does not match with 78 our X-ray pattern at this temperature. Due to slow kinetics of transformation, the phase of β-Ti already formed at 900 °C does not transform to α-Ti, as observed from the X-ray pattern at lower temperatures. This is even pronounced in the case of high-melting temperature elements/alloys just like the case of Ti-B-Fe system. In addition, from the HTXRD patterns, there was no evidence for any phase generation at 30, 700, and 900 °C, indicating that these MMCs were stable at these temperatures. 4.3 Pseudo-binary Ti-B-Fe Phase Diagram with Constant Fe and Variation in B Content CALPHAD calculations with three different contents of B (10, 20, and 30 mol. %) with constant Fe (10 mol. %) are performed to predict the optimal composition and processing temperature for attainment of sintering of Ti-TiB composites. Variation of the B content plays a crucial role in the evolution of phases in the Ti-B-Fe system, which can be simulated by CALPHAD calculations. Figure 4.3 presents the pseudo-binary phase diagram for the Ti-B-Fe system. This phase diagram was constructed with Fe fixed as 10 mol. % and B varied up to ~40 mol. %. The pseudo-binary phase diagram is helpful in determining the composition and the temperature for processing the MMCs in the Ti-B-Fe system. It can be seen that in this diagram, the compositions Ti-10B-10Fe and Ti-20B-10Fe lie within the phase field of TiB+β-Ti. This indicates the stability of two phases: the ductile matrix of β-Ti and the reinforcement of TiB in the region at 900 °C. On the other hand, the composition Ti-30B10Fe existed in a three phase field, indicating the stability of three phases: β-Ti, TiB, and TiFe at 900 °C. The pseudo-binary diagram also gave an idea of the extent of the β-Ti 79 phase field. The points shown by arrows in the diagram indicate the compositions and the chosen process temperatures. The sintering temperature was selected as 900 °C based on complete absence of liquid phase (below the liquids line), which facilitated sintering, without damaging the die. In order to ensure complete densification, some composition was needed to be densified at a high temperature (~1100 °C). The reason for this will be discussed later. 4.3.1 Densification Behavior Figures 4.4 (a) and (b) show the power input and the ram displacement of the EFAS recorded during the processing of composites at 900 °C. The power input behaviors of the three MMCs were largely similar due to the constant heating rate, holding time, and cooling rate. In all the MMCs, once target temperature (900 °C) was reached, the power input dropped along with a small fluctuation. This fluctuation may be due to the decrease in the resistance due to the increase in contact area between particles as a result of sintering. Once the sintering and densification attained completion, the power reached a steady-state due to no further variation in contact area resistance. There was, however, a slight variation in the steady-state power with composition, which was insignificant. The ram displacement was also recorded for each MMC composition during the EFAS process. As seen in Figure 4.4 (b), the ram displacement was rapid in the initial stage of process followed by a rapid densification of the material. For the three MMCs, the actual densification started around 500 °C after about 600 s and was completed after about 1080 s from the start of the process. The rapid ram displacement in the initial stages could be attributed thermodynamically to the formation of ternary eutectic liquid phase, which may 80 facilitate the reaction sintering of TiB and assist the densification of the composites. The eutectic reaction between Ti(L)+TiFe also occurred close to 1085 °C (see Figure 3.1), which led to a shrinkage in the volume in the Ti-Fe phase diagram [88]. However, for the Ti-Fe-B system, with an increase in B content, residual Ti should decrease. The shrinkage levels were calculated based on the ram displacement and are presented in Table 4.1 where ΔL changes with the ram displacement, as measured from the ram displacement data. The final height was measured by the micrometer, and initial height was calculated by adding the ΔL and final height. The experimental density value as well as the theoretical density were presented in Tables 4.2-4.3. The theoretical densities of the composite were calculated using the consideration of the reaction stoichiometry (Ti + TiB2 → 2TiB) to calculate the final weight of each composition based on initial amounts added in the sintering reaction, the density and molecular weight of each phases. Theoretical density of the composite is then calculated as the ratio of the total mass to total volume of the constituent phases (Ti, TiB, and β-Ti) present in the sintered composite. The experimental densities were calculated by the Archimedes principle based on the following equation: ρ = 1− ρw (Ww / Wa ) (4.1) where ρw is density of water = 1 g/cc, Wa is weight of sample in air, and Ww is weight of sample in water. The density measurements indicated that all the MMCs were nearly 100% dense. The densification was also confirmed by the examinations of microstructures under the optical microscope, which indicated the absence of the porosity in the processed MMCs. 81 4.3.2 X-ray Diffraction Analysis X-ray diffraction patterns of the different MMCs are presented in Figure 4.5. The X-ray patterns confirm the presence of two phases (β-Ti and TiB) in the MMCs Ti-10B10Fe and Ti-20B-10Fe and reveals three phases of β-Ti, TiB, and TiFe in Ti-30B-10Fe. It is seen that TiB2 peaks were absent in all the MMCs, which indicates that the sintering reaction reaches completion at 900 °C within 4 hours. Thus, it seems that all TiB2 particles were reacted to form the TiB phase. The patterns also reveal that the full β-stabilization was achieved since the β-Ti peak exhibited strong intense peak of (110) plane. There was also a complete absence of any α-Ti peak in all composites, which is expected due to slow kinetics of transformation from beta to alpha phase, as was previously observed in Ti-Fe binary phase diagram [84]. Fe is known as a strong beta-stabilizer element, which makes it difficult to form α-Ti at room temperature. There is no evidence of the formation of the Ti3B4 phase, which is reflected in the X-ray pattern by lack of reflections. The Ti3B4 is intermediate phase between TiB and TiB2. In a previous study [89], Ti3B4 was identified by peritectic reaction between TiB2 and liquid phases at 2200 °C and it was stable down to 1690 °C. There is no Ti3B4 present within XRD limit, resolution. The composite Ti-30B10Fe with the highest B content revealed the occurrence of intermetallic TiFe phase due to the diffusion of Fe in the residual Ti matrix. The XRD patterns indicate that as B content increased, the dominants peaks of TiB ((210), (102), and (112)) also increased and the dominants peak of β-Ti ((110)) decreased. It is also worthwhile to mention that as the boron content increased, β-Ti peaks (110) and (200) shifted to the right. Table 4.4 provides a comparison between the value of d-spacings from the X-ray pattern in the Joint Committee on Powder Diffraction Standards (JCPDS) 82 value. It can be seen that the β-Ti d-spacings of (110)β-Ti and (200)β-Ti decrease, which indicates that the volume of β-unit cell decrease with B content. The value of d-spacings for the characteristic peaks of β-Ti in JCPDS data were different from the experimental data as observed for all the MMCs Ti-10B-10Fe, Ti-20B10Fe, and Ti-30B-10Fe. The JCPDS data of β-Ti phase are from the high temperature XRD measurements of pure titanium above 880 °C where the β-Ti was stable. In the JCPDS data, the β phase was stabilized by temperature, whereas in this study, β-Ti was stabilized by alloying with Fe. The Fe is incorporated in the lattice of Ti, which changes the interplanar spacing of β-Ti phase [90]. This is reflected in the peak shift to the right, relative to the JCPDS data. It should be further noted that, as the B content increased, the residual Ti decreased and more Fe diffused into the titanium matrix. This further reduced the dspacing, resulting in the shifting of the diffraction peaks towards the right with increasing B content. The direct comparison method was used to determine the relative volume fraction of phases in the MMCs. The integrated intensities of selected peaks were determined for this purpose. For β-Ti, TiB, and TiFe phases, (110), (111), and (211) planes were selected, respectively. The integrated intensities determined from the XRD pattern are tabulated in Tables 4.5-4.7. The tables also present the values of structure factor, Lorentz-polarization factor, multiplicity factor, and unit cell volume for the phases considered. The calculated volume fractions of each phase are given in Table 4.8 for each MMC. 83 4.3.3 Microstructure The microstructures of the MMCs are presented in Figures 4.6 (a-f). Figures 4.6 (a) and (b) show the microstructures of the composite Ti-10B-10Fe at a low and a high magnification, respectively. The microstructures consist of two phases: β-Ti and TiB phase. The TiB whiskers, appearing as some elongated as well as clustered forms, are seen well-distributed throughout the β-Ti matrix. The morphology of TiB whiskers in this composite occurred in two forms. In the first form, it occurred as primary whisker, which was randomly distributed in the matrix with size thickness around 3 to 5 µm and length of 15 to 22 µm. In the second form, it occurred as clustered TiB whiskers, with a cluster size of about ~13 µm. In the microstructure of the composite Ti-20B-10Fe as shown in Figure 4.6 (c) and (d), the TiB has formed in clusters, but with a small amount of TiB whiskers growing independently into the β-Ti matrix. In general, the TiB phase in the composites Ti-10B10Fe and Ti-20B-10Fe is present as short and clustered whiskers. In the composite of composition Ti-30B-10Fe, the microstructure consisted of three phases, β-Ti, TiB, and TiFe, as shown in Figures 4.6 (e) and (f). The microstructure has a volume fraction of TiB greater than 0.70. As the B content increased, it seemed that the TiB formation occurred in the form of densely packed TiB whiskers. 84 4.4 Pseudo-binary Ti-B-Fe Phase Diagram with Constant B and Variation in Fe Content CALPHAD approach was also used to determine the composition and temperature window to produce the Ti-TiB MMCs with varying amount of Fe. Figure 4.7 illustrates the pseudo-binary phase diagram constructed using Thermo-Calc. The pseudo-binary phase diagram was calculated by keeping the B constant at 10 mol. % and varying Fe content. Three MMC compositions were selected as indicated by the circles in the diagram. The MMCs were EFAS processed at 900 °C with reaction sintering time as 4 hours. 4.4.1 Densification Behavior Figures 4.8 (a) and (b) illustrate the power input and the ram displacement during the EFAS processing. Table 4.9 presented the shrinkage percentage levels with changing Fe content. Figures 4.8 (a) and (b) show the power input and ram displacements during processing of the MMCs of varying Fe content with constant B, respectively. The power inputs of the three compositions were similar since they were subjected to similar heat cycle treatment. The ram displacement behavior for the three composition is also similar where increasing the Fe content helps in densification, as indicated in Table 4.9. Higher Fe in the composite would lead to the formation of the eutectic (Ti(L)+TiFe) as shown in TiFe phase diagram [84], which helps in the densification of the composite. The relative density in Table 4.10 revealed that all Ti-TiB MMCs with varying Fe fully densified to nearly 100 %. 85 4.4.2 X-ray Diffraction Analysis Figure 4.9 illustrates the room temperature X-ray diffraction patterns of the Ti-TiB MMCs within varying amount of Fe. The X-ray patterns confirm the presence of three phases (β-Ti, TiB, and TiFe) in the MMCs. It can be seen that the volume fraction of TiFe increased with increase in Fe as evident from the (110)TiFe peak in the patterns. The volume fractions of phase estimated using the direct comparison method are given in Table 4.11. 4.3.3 Microstructure Figures 4.10 (a) and (b) show the microstructures of the composite Ti-10B-20Fe at low and high magnifications. The microstructures consist of three phases: β-Ti, TiB, and TiFe phase. These phases are identified in Figure 4.10 (b). The TiB phase were observed as clusters with an average size of 12 μm and TiFe phases were observed in the form of islands with an eutectoid morphology. Figures 4.10 (c) and (d) show the microstructure at low and high magnifications, respectively, for the composite with composition Ti-10B-30Fe. Microstructure consisted of three phases, β-Ti, TiB, and TiFe. The TiFe phase is found to be in the form of clusters of particles embedded in β-Ti phase. Figures 4.10 (e) and (f) show the microstructures of Ti-20B-20Fe composite at low and high magnifications. TiB phase were seen as large grains around which the TiFe and β-Ti phases were present. 86 4.5 Mechanical Properties of Ti-B-Fe System with Variation in B Content 4.5.1 Hardness The Vickers hardness of TiB phase was determined as 1800 kg/mm2 [60]. The hardness of β-Ti in MMC could not be measured accurately because the sizes of indentations at 1 kg load were large enough that any indentation placed in β phase always intercepted other phases as well. Therefore, any hardness value depends on the location of the randomly placed indent on the as-polished surface of the MMC. For each composite composition, about twenty random hardness measurements were made in the as-polished condition. Because of the statistical placement of hardness indentations in the microstructure, each hardness value will be affected both by the hardness of β-Ti and TiB phases depending on the extent to which the indentation lied in the phases. The measured data are shown in Figure 4.11 for the three MMCs. The hardness values of the composites varied between 468-567, 612-794, and 10171144 kg/mm2, for the TiB volume fraction of 0.22, 0.27, and 0.79, respectively. As the volume fraction of TiB increased in the microstructure, the hardness level also increased. From the hardness distribution data, it is possible to a back calculation to determine the hardness of β-Ti for each MMC. For each composite, the distribution of the hardness followed a cumulative distribution function, which can be fitted using the Weibull type function: H− Hβ P = 1 – exp [−C {H TiB − Hβ n } ] (4.2) where H is the hardness of the composite and HTiB and Hβ are the hardness values of TiB and β-Ti phase, respectively. By taking HTiB = 1800 Kg/mm2, the hardness values of the beta phase were calculated as 390, 550, and 890 kg/mm2 for the composites with TiB 87 volume fractions of 0.22, 0.27, and 0.79, respectively, which is due to solid solution strengthening. These values give the best fit of Weibull function to the hardness determinations. From the hardness results, it is observed that the hardness of β-Ti increased with an increase in B content in the MMCs. This may be attributed to the strengthening from the formation of TiB whiskers in the β-Ti phase. This is consistent with microstructure of the composite Ti-30B-10Fe, which shows that there was increased formation of TiB whiskers with increasing B content. 4.5.2 Tensile Properties The composite with 0.22 volume fraction of TiB (Ti-10B-10Fe) was selected to evaluate the mechanical properties. Figure 4.12 presents the tensile data for the composite Ti-10B-10Fe. From the tensile data, the elastic modulus, ultimate tensile strength, and elongation were calculated, and these data are presented in Table 4.12. In two phase composites, the elastic modulus is strongly dependent on the volume fraction of the reinforcement phase. In case of Ti-TiB composites, a model developed by Chandran [91] on the basis of parallel and series arrangement of two phases, was found to accurately represent the variation of elastic modules with the reinforcement volume fraction in a number of MMCs. The elastic modulus of the composite as a function of the volume fraction of the two phases is given by: Ec = [Ep Em + E2m (1+C)2 − E2m ] (1+C) (Ep − Em )C+ Em (1+C)3 (4.3) where Ec is the elastic modulus of the composite and Ep and Em are the elastic modulus of particle or inclusion and the matrix material, respectively. The parameter C is given by: 88 1 1 C =[V ]3 − 1 p (4.4) where Vp is the volume fraction of the particle. For TiB Vf = 0.22, and with the elastic modulus of β-Ti (90 GPa) and the elastic modulus of TiB (370 GPa), the elastic modulus of any Ti-TiB composite can be calculated. The experimental elastic modulus values shown in Figure 4.13 are quite close to that predicted by the model. 4.5.2.1 Fractography The fracture morphology of the composite Ti-10B-10Fe was investigated using the scanning electron microscope. Figures 4.14 (a-d) illustrate the fracture surface morphology in the composite Ti-10B-10Fe. The micrographs show that the dominant mode of fracture is cleavage with no plastic deformation. In cleavage fracture, crack propagation occurs very fast with limited dislocation activity. The brittle fracture is indicated by the occurrence of the cleavage facts in the β-Ti matrix, with the crack propagating along the (001) cleavage planes [92]. At high magnification, the river lines contain some particles that are found in the β-Ti matrix. These particles could be nano-precipitates of TiFe phase. This can cause the brittleness of the βTi phase. 4.5.3 Flexure Strength The composite of composition Ti-30B-10Fe with TiB Vf = 0.79 was selected for flexure strength testing. Ten specimens were used for the flexure strength measurements by four-point bending test. The test results are presented in Figure 4.15. The flexural strength of the beam was calculated by using the following equation (4.5): 89 3PL S = 4bd2 (4.5) where P is the fracture load, L is the outer span (40 mm), b is the specimen width, and d is the specimen thickness. The flexural strength varied from 556 MPa to 727 MPa. The average strength is 645 MPa. Figure 4.15 (b) presents the results of flexure testing in the form of a Weibull plot. The parameters of the Weibull distribution are the Weibull modulus (m) and the characteristic strength (σθ). The line in the figure is a fit based on two parameter Weibull distribution: σ Pf = 1 − exp(−(σₒ))m (4.6) where Pf is the probability of failure, m is Weibull modulus, σ is the strength, and σθ is the characteristic strength corresponding to Pf = 0.632. The value of the characteristic strength and the Weibull modulus were obtained as: σθ = 674 MPa and m = 12. A previous study [93] showed that most of the structural ceramics showed Weibull moduli in the range of 5 to 15. For metals, these volume are in the range of 30 to 100. In the present composite, Weibull modulus is close to that of ceramics. The fracture surface of a flexure tested specimen was investigated using the scanning electron microscope (SEM). Figures 4.16 (a) and (b) illustrate the brittle fracture, occurring by cleavage in the flexural specimens. Crack propagation appears to occur without any occurrence of plastic deformation. To achieve better strength, the amount of Fe should be restricted to eliminate the formation of the intermetallic phase as TiFe. The presence of TiFe phase may also decrease the ductility due to precipitation of TiFe in beta phase and this may be the cause of brittle fracture. 90 4.5.4 Fracture Toughness The fracture toughness (KIC) of the composite of composition Ti-30B-10Fe was determined by using the single edge cracked beam method. Figure 4.17 illustrated the load displacement traces of five specimens. The average value of fracture toughness was 9.4 MPa√m. This value was relatively high compared to the value of 4.5 MPa√m for TiB ceramic, which was reported in previous work [94]. In that study [94], there was no Fe intended for the stabilization of the ductile phase of β-Ti. In general, the toughness of composite could be enhanced by having ductile β-Ti phase as ductile phase and by eliminating the intermetallic TiFe phase. However, in the present composite, there was no ductility in the β-Ti phase and there was significant amount of TiFe phase distributed as particles in the microstructure. Thus, the lack of high fracture toughness can be explained. 91 (a) Figure 4.1 Isothermal section of Ti-B-Fe at (a) 30 °C, (b) 700 °C and (c) 900 °C calculated from the thermodynamic data using Thermo-Calc. 92 (b) (c) Figure 4.1 Continued. 93 (a) (b) Figure 4.2 HT-XRD patterns for MMCs (a) Ti-10B-10Fe, (b) Ti-20B-10Fe, (c) Ti30B-10Fe, (d) Ti-10B-20Fe, (e) Ti-10B-30Fe, and (f) Ti-20B-20Fe; at room temperature, 700, and 900 °C. 94 (c) (d) Figure 4.2 Continued. 95 (e) (f) Figure 4.2 Continued. 96 Figure 4.3 Pseudo-binary phase diagram for Ti-B-Fe system calculated from the thermodynamic data (Alloy solutions Database v6.0) using Thermo-Calc. 97 3500 1 1400 (a) 3000 10 at.% B Power (W) 20 at.% B 30 at.% B 2000 1200 -1 1000 Ram Displacemrnt (mm) 2500 0 -2 1500 800 900 °C -3 1000 600 10 at.% B 20 at.% B -4 500 -5 0 0 5000 10000 15000 Time (Sec) 20000 400 30 at.% B Temperature (°C) (b) 200 -6 0 5000 10000 15000 0 20000 Time (Sec) Figure 4.4 EFAS process. (a) Power input versus time curves of EFAS MMCs processed at different compositions. (b) Ram displacement during EFAS processing and the temperature-time profile for the processing at 900 °C. 20 30 40 50 60 70 Figure 4.5 Room temperature X-ray diffraction of the synthesized MMCs. TiB (122) TiB (122) β-Ti (211) TiB (410) + TiB (312) TiB (312) β-Ti (211) TiB (020) β-Ti (200) TiB (020) β-Ti (200) TiB (112) TiFe (211) TiB (122) TiB (121) TiB (312) TiB (410) TiB (113) β-Ti (211) TiB (401) TiB (020) TiFe (200) β-Ti (200) TiB (112) TiB (301) TiB (301) TiB (301) TiB (112) TiB (211) TiB (211) TiB (211) TiB (210) TiB (102) TiFe (110) TiB (111) β-Ti (110) TiB (011) TiB (200) TiB (210) TiB (102) β-Ti (110) Ti20B10Fe TiB (210) TiB (102) TiB (111) TiB (111) β-Ti (110) Ti10B10Fe TiB (011) TiB (200) TiB (101) Ti30B10Fe TiB (011) TiB (200) Intensity (A.U) 98 80 99 Figure 4.6 Microstructures of the Ti-10B-10Fe (a-b), Ti-20B-10Fe (c-d), and Ti30B-10Fe (e-f) MMCs. 100 Figure 4.7 Pseudo-binary phase diagram for Ti-B-Fe system calculated from the thermodynamic data (Alloy solutions Database v6.0) using Thermo-Calc. 101 3500 1 1400 (a) (b) 30 at.% Fe 10 at.% Fe 20 at.% Fe 2000 1500 1000 500 0 0 10000 Time (sec) 20000 1200 0 1000 -1 800 -2 900 °C -3 600 10 at.% Fe 20 at.% Fe 30 at.% Fe -4 -5 0 5000 10000 15000 400 Temperature (°C) Power (W) 2500 Ram Displacement (mm) 3000 200 0 20000 Time (sec) Figure 4.8 EFAS process. (a) Power input versus time curves of EFAS MMCs processed at different compositions. (b) Ram displacement during EFAS processing and the temperature-time profile for the processing at 900 °C. 102 Figure 4.9 Room temperature X-ray diffraction of the synthesized MMCs. 103 Figure 4.10 Microstructures of the Ti-10B-20Fe (a-b), Ti-10B-30Fe (c-d), and Ti20B-20Fe (e-f) MMCs. 104 Figure 4.11 Cumulative probability distributions of hardness values. The lines represent fits to a Weibull-type function. 105 Figure 4.12 Tensile data of composite Ti-10B-10Fe. 106 Figure 4.13 The theoretical elastic modulus of Ti-TiB composite. 107 Figure 4.14 Fracture morphology. (a-d) Fracture surface of composite Ti-10B-10Fe. 108 Figure 4.14 Continued. 109 1 727 MPa (b) 0.5 0.75 Ln(Ln(1/(1-pf))) Probability of Fracture (a) 0.5 0.25 556 MPa 600 σ (MPa) -1.5 -2.5 0 500 -0.5 79 Vol.% TiB: m=12 = 674 MPa 700 6.3 6.4 6.5 ln(σ) Figure 4.15 Weibull distribution (a-b) composite Ti-30B-10Fe. 6.6 110 Figure 4.16 SEM images (a-b) fracture surface of Ti-30B-10Fe composite. 111 Figure 4.17 Fracture toughness of composite Ti-30B-10Fe. 112 Table 4.1 Total shrinkage in various MMCs in the Ti-B-Fe system ΔL Final height Initial height (mm) (mm) (mm) Ti-10B-10Fe 3.64 7.86 11.49 32 Ti-20B-10Fe 4.05 7.66 11.71 35 Ti-30B-10Fe 4.05 5.98 10.03 40 MMC Shrinkage % Table 4.2 Density and atomic/molecular weight of the elements/compounds Elements/Compounds Density (ρ) in g/cc Atomic/Molecular weight TiB 4.56 58.68 Ti 4.51 10.81 Fe 7.87 55.84 Table 4.3 MMC compositions and densities Density (g/𝐜𝐜) MMCs Relative density Experimental Theoretical Ti-10B-10Fe 4.75 4.77 0.996 Ti-20B-10Fe 4.78 4.80 0.996 Ti-30B-10Fe 4.83 4.84 0.998 113 Table 4.4 The effect of composition on the d-spacing of (110)β-Ti and (200)β-Ti Composite (110)β-Ti d (Å) (200)β-Ti d (Å) JCPDS 2.34 1.65 Ti-10B-10Fe 2.27 1.60 Ti-20B-10Fe 2.26 1.59 Ti-30B-10Fe 2.24 1.58 Table 4.5 Calculation of X-ray diffraction pattern for β-Ti phase Vₒ 2θ (°) |𝑭|𝟐 MMC (hkl) Ti-10B-10Fe 110 17.18 39.66 Ti-20B-10Fe 110 17.18 Ti-30B-10Fe 110 17.18 (Å𝟑 ) I β-Ti L P 1093 14.7 12 1268 11230 39.83 1093 14.6 12 1251 11123 40.15 1093 14.3 12 597 10932 (cps deg.) R Table 4.6 Calculation of X-ray diffraction pattern for TiB phase Vₒ I TiB 2θ (°) |𝐅|𝟐 L P 85.17 38.38 1426 15.82 8 69 2119 111 85.17 38.29 1426 15.90 8 91 2130 111 85.17 38.51 1426 15.70 8 209 2103 MMC (hkl) Ti-10B-10Fe 111 Ti-20B-10Fe Ti-30B-10Fe (Å𝟑 ) (cps deg.) R 114 Table 4.7 Calculation of X-ray diffraction pattern for TiFe phase I TiFe Vₒ MMC (hkl) Ti-30B-10Fe 211 (Å𝟑 ) 26.66 2θ (°) |𝐅|𝟐 L P R (cps deg.) 78.54 819 3.35 8 40 823 Table 4.8 Volume fraction of phases in the MMCs Actual volume fraction MMC β-Ti TiB TiFe Ti-10B-10Fe 0.78 0.22 - Ti-20B-10Fe 0.73 0.27 - Ti-30B-10Fe 0.13 0.79 0.08 Table 4.9 Shrinkage in various MMCs in the Ti-B-Fe system ΔL Final height Initial height (mm) (mm) (mm) Ti-10B-10Fe 3.64 7.86 11.50 32 Ti-10B-20Fe 3.54 7.21 10.75 33 Ti-10B-30Fe 3.67 6.80 10.45 35 MMC Shrinkage % 115 Table 4.10 Experimental and theoretical densities of Ti-TiB MMCs with varying Fe content Density (g/𝐜𝐜) MMCs Relative density Experimental Theoretical Ti-10B-10Fe 4.75 4.77 0.996 Ti-10B-20Fe 5.03 5.04 0.998 Ti-10B-30Fe 5.33 5.34 0.998 Ti-20B-20Fe 5.07 5.10 0.994 Table 4.11 Volume fraction of phases calculated using XRD data Actual volume fraction MMC β-Ti TiB TiFe Ti-10B-20Fe 0.66 0.26 0.07 Ti-10B-30Fe 0.49 0.26 0.25 Ti-20B-20Fe 0.27 0.49 0.23 116 Table 4.12 Tensile properties of composite Ti-10B-10Fe Ultimate MMC Vf of Tensile Ti-10B-10Fe TiB Strength Elongation (%) (MPa) Experimental Calculated Elastic Elastic Modulus Modulus (GPa) (GPa) Test 1 0.22 305.29 0 124.02 Test 2 0.22 578.80 0 119.81 Average 0.22 442.05 0 121.91 132.5 CHAPTER 5 CONCLUSIONS 1. The selected compositions of the MMCs based on Ti-TiB in the Ti-B-Fe system with 0.22, 0.27, and 0.79 volume fraction of boride phases were designed from CALPHAD approach and successfully synthesized to fully density by EFAS at a low temperature of 900 °C in one step; this sintering process achieved stable phases at room temperature. 2. The microstructure showed a uniform distribution of TiB and β-Ti phases in the MMCs. The β-stabilizer (Fe) helped to stabilize the β-Ti phase at room temperature. 3. The average values obtained from the Vickers hardness test were 530, 680, and 1080 kg/mm2 for volume fractions 0.22, 0.27, and 0.79 of boride phases, respectively. The hardness value showed variation, due to the random indentation on the surface of the MMC that contains two different phases (β-Ti and TiB). The Weibull-type described the variation of the Vickers hardness value by the exponential function. 4. The tensile test of Ti-10B-10Fe composite with a volume fraction of boride phase as 0.22 revealed an average elastic modulus of 122 GPa and ultimate tensile strength of 442 MPa. There was no ductility seen in this composite due to cleavage behavior of βTi matrix. TiFe was observed on the β-Ti matrix. 118 5. The flexural strength of Ti-30B-10Fe composite with a volume fraction of the boride phase as 0.79 was measured to be in the range of 556 and 727 MPa. The results reflected no ductility due to the high-volume fraction of the TiB phase. 6. The fracture surface of the composite revealed brittle failure caused by cleavage from the β matrix. The formation of Ti-Fe was also observed in the matrix. To get better strength, TiFe phase should be eliminated. 7. The fracture toughness showed an average value of 9.4 MPa√m for 0.79 volume fraction of boride phase. The nature of the fracture surface was observed to be brittle as evinced by cleavage along the β-Ti matrix. Toughness may be enhanced by the incorporation of more metallic β-Ti matrix along with the elimination of the intermetallic phase (TiFe). 8. Currently, the toughness of the composite is restricted by its brittle fracture. The scope of future research is based on enhancing the mechanical properties through addition of alloying elements and avoiding the formation of intermetallic phases by restricting the Fe content. CHAPTER 6 DESIGN, SYNTHESIS, AND PROPERTIES OF TITANIUM BORIDE CERMET MATERIAL BASED ON TI-B-FE-MN SYSTEM (Manuscript in preparation “Materials Science and Engineering: A”) 6.1 Introduction Rapid in-situ synthesis of bulk ceramic-metal composite (Cermet) containing a high volume fraction of titanium boride (TiB) as a reinforcement phase and β-Ti as a matrix phase was performed at a relatively low temperature using Electric-Field-ActivatedSintering (EFAS). The cermet consisted of TiB whiskers distributed uniformly throughout the volume of cermet. The cermet could be fully densified at a sintering temperature of 1025 °C, 4 hours holding time, and 10 MPa pressure. Optical microscopy and X-ray diffraction (XRD) analysis confirmed that TiB whiskers have been formed within the cermet structure. A pseudo binary diagram was constructed using CALPHAD method to determine liquid phase fields to enable identification of optimum processing temperature. Mechanical testing reveal that the fracture toughness was good (~9 MPa√m), but the flexural strength was not high. The microstructural factors giving good toughness and poor strength are discussed. Novel hard material applications required special mechanical properties such as 120 high strength and high ductility. Due to this need, a new family has been introduced (composite material) to accommodate for these mechanical properties. These composite materials are grouped into three main types: metal matrix composite (MMC), ceramic matrix composite (cermet), and bulk ceramic composite. The advantage of the composite material is that two phases are presented in the microstructure which are: a ceramic phase that provides hardness and stiffness and a metallic phase that provides ductility and toughness. The main difference between these three types of composite materials is based on the percentage of volume fraction of the hard phase. The MMC is considered to be having the lowest percentage of the volume fraction (< 30 % Vf), then followed by cermet (30-95 % Vf), and finally, the bulk ceramic composite (< 95 % Vf). In this research study, the focus is on designing and synthesis of a cermet composite because of the interesting feature of the microstructure. The microstructure of the cermet showed binding between the ceramic phase and the metallic phase, which is distributed uniformly discontinuously throughout the volume of the cermet. The industrial cermets known are tungsten carbide (WC), titanium carbonitride (TiCN), and titanium carbide (TiC). These cermets are used in drilling and cutting tools applications [95-97]. These cermets were manufactured by liquid phase sintering, which required high temperatures [98, 99] Metallic materials were added to these cermets to enhance mechanical properties and to reduce the processing melting point, the processing details of these cermets are beyond the scope of this study [100-105]. Electric-Field-Activated-Sintering (EFAS) can achieve high heating rate based on Joule’s heating principle, which can accelerate local heating and diffusion, reaction, and sintering of constituent powders giving the possibility of synthesizing cermet at low 121 temperatures. The research hypotheses are to determine whether using β-Ti as a metallic phase, to improve ductility, toughness, and investigating whether TiB can improve strength and hardness of the cermet. TiB is a new class of ceramic material phase that is presented as a reinforcement in the cermet, which enhances the mechanical properties such as hardness and stiffness. This material has been previously prepared by our research group by using hot pressing [106108]. The properties of TiB are found to be excellent in comparison with Si3N4 [110] where wear resistance is high, hardness is (~1600–1800 kg/mm2) [60], Young's modulus is (~ 427 GPa) [110], flexure strength is (~ 850 MPa) [60], and fracture toughness is (5.2 MPa√m). Previous studies performed by our research group showed that a complete sintering reaction and densification of TiB by hot pressing requires at least 15-20 MPa of pressure at a temperature of 1350 °C for a period of 2 hours. In this research work, the cermet material has been synthesized by EFAS at a temperature of ~1025 °C for 4 hours, which resulted in complete reaction sintering and fully densified cermet with uniform TiB distribution in the β-Ti matrix. It has been reported that TiB phase formed as whisker shape after the reaction sintering [5]. Other studies [111] show that needle shaped TiB2 can be obtained using spark plasma sintering (SPS) by using Ti and B powders. However, it has been reported that the composite was found to be not fully densified, due to lack of heat, which is evident by the presence of porosity in the composite microstructure that was sintered at 1500 °C. This study has further confirmed that at least a temperature of 1600 °C is required to obtain fully densified composite by SPS. According to the Ti-B binary phase diagrams, α or β-stabilizers are required to be added to lower the temperature of synthesis to obtain TiB. A Ti-B-Fe-Mn pseudo binary 122 phase diagram has been constructed by CALPHAD approach to investigate the possible phase formation and adequate process temperature from 950 °C to 1050 °C, which is close to transient liquid phase for the following elements Fe-Mn. The objective of this research study is to achieve fully densified cermet with the lowest temperature possible. Fe-Mn powders addition has been required to enhance the sintering reaction through formation of the transient liquid phase. The processing was done by EFAS using Ti, TiB2, Fe, and Mn powders. Cermet evaluation contains microstructure, density, X-ray analysis, and mechanical testing. 6.2 CALPHAD Approach CALPHAD approach is a computational method used for a system of materials to find phase equilibria for the system of interest produced under certain conditions such as temperature, pressure, and composition. This approach can predict to a very high precision the outcome of the phase fields for binary and complex multisystem such as ternary and pseudo-binary [112, 113]. The establishment of these phases is governed based on thermodynamic principals and thermodynamic system extrapolations. In this study, CALPHAD approach has been utilized to determine pseudo-binary system of Ti-B-Fe-Mn. CALPHAD approach has been used to determine the optimum conditions for the current cermet synthesis in the phase stability region of Ti-B-Fe-Mn. Temperature and pressure were modeled by thermo-Calc software. Figure 6.1 shows the pseudo-binary phase diagram of the Ti–B–Fe system. The pseudo-binary phase diagram was calculated by keeping the B constant at 10 mol. % and varying Fe content. Ti-B-Fe system has α-Ti forming at T < 600 °C due to (TiB+α- 123 Ti+TiFe) phase field. Therefore, Ti-B-Fe-Mn system was considered because of the potential of stabilizing β-phase with Mn while avoiding the formation of α-Ti phase. Figure 6.2 shows the pseudo-binary phase diagram of the Ti–B–Fe–Mn system. TiB-Fe-Mn pseudo-binary indicates that β-Ti phase is stable at T < 600 °C and down to 500 °C. The pseudo binary diagram has been constructed by making the B composition constant for the TiB phase and the variation stems from the fixed ratio of Fe:Mn 0.5 mol %. The cermet composition was designed to be close to the eutectic point promoting partitioning of Fe/Mn in the β-Ti phase due to the wide range of temperature occurrence. The sintering temperature has been chosen to be 92 % - 100 % of the melting temperature (Tm), i.e. 950 °C – 1050 °C to evaluate the pseudo-binary phase diagram. From Figures 6.1 and 6.2, it can be assessed that the addition of Mn has yielded advantageous points such as the eutectic reaction point temperature has been decreased slightly by 10° C, which resulted in lowering of the process temperature and for β-Ti (β-Ti +TiB+Ti(Fe,Mn)), phase field has become larger in size, which indicated higher stability by stabilizing β-Ti to a temperature of less than 600 °C and eliminates the α-Ti phase. To prove the accuracy of the hypothetical approach that has been acquired, four processing temperatures have been selected to investigate the system performance with the optimum conditions. 6.3 Experimental The starting powder mixture in this study consists of Ti powder (Atlantic Equipment Engineers, NJ, 99.7 % pure; average particle size: 45 μm), TiB2 powder (99.7 % pure; average particle size: 10 μm), iron powder (Alfa Aesar, 99.5 % pure; average 124 particle size: 10 μm), and manganese (99.9 % pure; average particle size: 44 μm). The powder was mixed in a steel jar under argon inert gas atmosphere and was blended for 24 hours with rotational speed of 350 rpm with 1:10 mass ratio of Ti ball to powder. The composition mixture of the synthesis cermet by mole percentage is as follows: 30 % B, 5 % Fe, 9 % Mn, and 56 % Ti. A 10 ton EFAS (Model 10-4, Thermal Technologies, Santa Rosa, CA) was used to synthesize cermets at processing temperatures between 950 °C to 1050 °C. The EFAS processing conditions are: pressure 10 MPa, heating rate 50 °C/min, holding time 4 hours, and cooling rate 22.5 °C/min to room temperature. Samples microstructures have been observed after a regress grinding followed by polishing and fine finishing using 0.5 μm colloidal silica. After fine polishing, the samples were etched using Kroll's reagent. Density measurements were evaluated by Archimedes’ method. Vickers hardness tester (LECO, M-400) was applied to the samples at 1 kg load to measure hardness of the cermets. Eight random measurements across the sample have been taken to ensure accurate hardness reading. X-ray analysis has been done using the model (Rigaku, Miniflex 600), at a scan rate of 0.5 °/min to obtain X-ray diffraction patterns. Mechanical properties analysis has been obtained from flexural strength and fracture toughness tests following ASTM C1161-13 and ASTM C1421-10 standards with displacement rates of 0.5 and 0.18 mm/min, respectively, using four-point binding configuration. The bar samples dimensions for both tests are 45 × 4 × 3 mm3, and were cut by Electric-Discharge-Machining (EDM). 125 6.4 Results and Discussion 6.4.1 Processing The sintering reaction in EFAS process is given by: Ti + TiB2 + Fe + Mn = 2 (Ti,Fe,Mn) B + β-Ti (6.1) The left side of the equation are the starting powders and the right side of the equation are the products of the reaction, which includes four phases, TiB, TiFe, TiMn, and β-Ti. Iron and manganese are known as β-stabilizer elements that will react with β-Ti phase in solid solution yielding TiFe and TiMn. The sintering goal is to maximize the hardness and strength by enhancing TiB whisker formation, and maximizing toughness through stabilizing β-Ti at room temperature using stabilizing elements such as Fe and Mn. Figure 6.3 (a) shows the EFAS sintering processing conditions for the power input v/s time. The power input curves are resembling each other due to the constant heating/cooling rate and holding time. In all the input power curves, once the targeted temperature is reached, the input power drops followed by small fluctuation, then steadystate power is observed. Figure 6.3 (b) shows ram displacement v/s time, which was recorded from EFAS sintering process. Figure 6.3 (a) to correlates the densification behavior with the power input v/s temperature. Rapid densification behavior has been observed during initial sintering reaction for temperature range 950 °C – 1050 °C due to quaternary eutectic phase at 1085 °C, which is observed from TiFe binary phase diagram [88]. The densification started at about ~ 500 °C, 500 s and completed at around ~ 1500 s. Densification increases with increasing temperature, which is evident from ram displacement behavior around 1025 °C with displacement of ~ 5 mm. 126 Figure 6.4 shows X-ray patterns for the sintered cermet. The sample that has been sintered at 950 °C has TiB2 phase whereas other cermets have no TiB2 phase present at a temperature range of 1000 °C – 1050 °C due to complete reaction sintering. The XRD pattern shows high intense peak (110) of β-Ti, which indicates β-Ti matrix is stabilized at the room temperature. The volume fraction percentage of the sintered cermet at 1050 °C is 65 % TiB, 3 % β-Ti, 10 % TiFe, and 22 % TiMn that has been calculated by using direct comparison method that was described excessively in our previous work [5]. 6.4.2 Microstructure Figure 6.5 (a-d) shows the microstructure of cermet sintered at 1025 °C and 1050 °C. Figure 6.5 (a,b) shows full densification at 1025 °C and above due to complete formation of TiB compared to Figure 6.5 (c,d). From observing the microstructures of sintered cermet at 1025 °C - 1050 °C, a close microstructure pattern is observed, which supports the evidence of having full densified cermet structure indicating that full reaction has been achieved. Densification has occurred due to the advantage of EFAS, which creates maximum localized high temperature at the particle contact region that allows the formation of quaternary liquid phase from Joule’s heating. The microstructural pattern showed high similarities between sintered cermet composition at 1025 °C and 1050 °C. In general, four phases can be observed from the microstructure pattern (TiB, β-Ti, TiFe, and TiMn). The whisker shape of TiB phase is packed together in a bundle distributed continuously throughout the volume with low aspect ratio ~1-2 and that was due to the reduction in mean-free-distance. The island shape of β-Ti phase is distributed randomly between the whisker bundles with average size of 8- 127 12 μm. The β-Ti phase island modes were presented as big and small islands that influenced the distribution of the whiskers bundle in a way that big islands are located between the whiskers bundle and the small islands are located within the whiskers bundle. The small dark regions are TiFe and TiMn that existed between the islands and the whisker bundles. Small grain size averaging 0.7-2.3 μm was evident in comparison to known cermets WC, TiCN, and TiC with average grain size of 1.0-2.5, 3.5, and 25-45 μm, respectively [114116]. 6.4.3 Mechanical Properties Figure 6.6 shows the comparison between density and hardness. Increasing the processing temperature will increase density and hardness for the cermet. Hardness (1055 kg/mm2) and density (4.9 g/cm3) reach maximum stable values at 1025 °C. Figure 6.7 (a) shows the flexural strength behavior for the cermet sintered at 1050 °C. The flexural strength range is 359 – 652 MPa for the volume fraction of 65 % TiB, 3 % β-Ti, 10 % TiFe, and 22 % TiMn. Weibull distribution function has been applied to fit the data based on the following equation, using the two parameters (Weibull moduli and characteristic strength), which determines the evaluation of cermet reliability. In this study, the Weibull modulus value was found to be 7, and is in the range of ceramic material (515) [93] due to high volume fraction percentage of 65 % TiB. σ Pf = 1 − exp(−(σₒ))m (6.2) where Pf is the probability of failure, m is the Weibull modulus, and σₒ is the characteristic strength value taken at probability failure of 0.632. The cermet characteristic strength and Weibull moduli are 613 MPa and 7, respectively. Figure 6.7 (b-e) shows fractured surface 128 of the cermet sintered at 1050 °C. The fracture surface behavior is brittle fracture, which is indicated by cleavage due to the high concentration formation of two intermetallic phases (TiFe and TiMn) that led to inferior strength. Figure 6.8 shows the fracture toughness (KIC), which is determined by using single edge cracked beam specimens four-point bending test. The average fracture toughness value for the cermet sintered at 1050 °C is 8.6 MPa√m with 65 % TiB volume fraction percentage. 6.5 Conclusions In summary, cermet has been designed based on CALPHAD approach and has accurately predicted the phases present in the microstructure and X-ray diffraction patterns. Cermet in the Ti-B-Fe-Mn system with composition 56 % Ti, 30 % B, 5 % Fe, and 9 % Mn has been successfully synthesized with full density at a low temperature ~1025 °C in a single step using EFAS. Fast reaction and densification of cermet has been triggered by the formation of a quaternary eutectic liquid phase resulting in the uniform distribution of TiB phase, which is evident from the observations of microstructure and average Vickers hardness values. Cermet flexural strength value 543 MPa with volume fraction percentage of 65 % TiB and brittle fracture behavior has been dominated by cleavage. The average value of fracture toughness is 8.6 MPa√m. In this study, the fracture toughness was limited due to the cleavage behavior of β-Ti phase. Current development is in progress to reduce the intermetallic phases, and to improve the mechanical properties of the cermet’s significantly. 129 Figure 6.1 Pseudo-binary phase diagram of the Ti–B–Fe system. 130 Figure 6.2 Pseudo-binary phase diagram of the Ti–B–Fe–Mn system. 131 1025 °C 1000 °C 950 °C 2000 Ram Displacement (mm) 1050 °C 2500 1025 °C 0 (b) 800 -2 600 -3 1500 1000 °C -4 1000 950 °C 1025 °C 200 -6 0 -7 0 0 5000 10000 15000 20000 Time (sec) 400 1050 °C -5 500 1000 -1 Temperature (°C) (a) 3000 Power (W) 1200 1 3500 0 5000 10000 15000 Time (sec) Figure 6.3 EFAS process. (a) EFAS sintering processing conditions for the power input v/s time. (b) Ram displacement v/s time. 132 Figure 6.4 X-ray patterns for the sintered cermet. 133 Figure 6.5 Microstructure. (a-d) The microstructure of cermet sintered at 1025 °C and 1050 °C. 134 Figure 6.5 Continued. 4.95 Density (g/cc) 4.9 1300 4.85 1100 4.8 4.75 900 4.7 700 4.65 4.6 500 930 960 990 102010501080 Temperature (°C) Vickers microhardness ((kg/𝒎𝒎𝟐 ) @ 1 kgf) 135 Figure 6.6 Comparison between density and hardness. 136 Probability of Fracture (a) 1 65 vol.% TiB: m=7 =613 MPa 652 MPa 0.75 0.5 0.25 359 MPa 0 200 300 400 500 600 700 σ (MPa) Figure 6.7 Cermet composite. (a) Flexural strength behavior for the cermet. (b-e) Fractured surface of the cermet sintered at 1050 °C. 137 Figure 6.7 Continued. 138 10 Average KIC 9 K (MPa.√m) 8 7 6 5 4 3 2 1 0 0 0.01 0.02 0.03 0.04 Displacement (mm) Figure 6.8 The fracture toughness (KIC) behavior for the cermet. CHAPTER 7 COMPUTATIONAL DESIGN, PHASE EQUILIBRIA, AND PROCESSING OF TITANIUM METAL MATRIX COMPOSITES IN TI-B-MO-FE-AL SYSTEM (Manuscript in preparation “Scripta Materilia”) 7.1 Introduction New metal matrix composite (MMC) material class containing titanium boride (TiB) and beta titanium (β-Ti) has been designed and processed, which has high toughness and high strength. This research study is based on using Electric-Field-Activated-Sintering (EFAS) to densify the powder by rapid reaction sintering. The powder of interest consists of Ti, TiB2, Mo, Fe, and Al. The MMC contains two phases that are present in the microstructure as TiB ceramic phase reinforcement and β-Ti metallic phase matrix. This phase has been formed in-situ upon the reaction between the powder particles. MMC has been evaluated based on the hardness, flexural strength, fracture toughness, and tensile strength. This MMC class yields excellent mechanical properties with a Vickers hardness of 375 kg/mm2, flexural strength of 1622 MPa, fracture toughness of 23 MPa√m, ultimate tensile strength of 1041 MPa, elastic modulus of 114 GPa, yield strength 910 MPa, and elongation of 0.27 %. The MMC showed remarkable results, which was reflected upon the fractured surface by having ductile fracture along the entire volume of the specimen that 140 allows the plastic deformation to occur in the matrix of the MMC, resulting in high fracture toughness values. This research study also includes a modeling design for the MMC by using calculation phase diagram (CALPHAD) approach. The MMCs contain two phases that govern the hardness and strength due to ceramic phase, toughness and ductility due to metallic phase. The ceramic phase is the reinforcement and the metal phase is the matrix. Both phases are coherent and thermodynamically stable at room temperature [58]. The volume fraction of the reinforcement ceramic phase is less than 0.3 Vf and it is uniformly and discontinuously distributed throughout the volume of the MMC. One of the biggest advantages of this novel MMC that contains ceramic phase (boride phase) is that in-situ reactions during the sintering are viable and fast, which results in whiskers containing high aspect ratio with constant density and no intermediate phases between matrix and whiskers. The boride phase is superior to the other popular reinforcements such as SiC, Al2O3, Si3N4, and B4C in terms of interface reaction where boride phase has no intermediate phase. This advantage makes the Ti-TiB MMC highly attractive to be developed and used in varies applications such as lightweight aerospace and automotive applications. In this study, new composition of MMC has been synthesized in which the ceramic phase (TiB) has been created in-situ by reaction sintering of powders during the EFAS process. This in-situ reaction gives two phases: TiB and β-Ti. Also, other alloying elements (Fe-Mo-Al) were added to provide strength that been described by solid solution strengthening principle. The greatest advantage of this Ti-B-Fe-Mo-Al system lies in its high flexibility, resulting in a variety of compositions that covers wide range in applications presented by low volume fraction (MMC) and high volume fraction close to 100 % nano- 141 ceramic. This MMC can be used in applications that require strength and high toughness and the nano-ceramic can be used in applications that require hardness and strength. Many benefits stem from the ceramic TiB phase such as high stiffness due to the high electron density of B–B bonds presented in the whiskers formed in the MMC. These whiskers enhance the mechanical properties such as strength and hardness in the MMCs [60]. The mechanical properties of the monolithic TiB phase according to previous studies are as follows: elastic modulus: 370–425 GPa [12, 110], fracture toughness: 4.5 MPa√m [94], and Vickers hardness: 1800 kg/mm2 [60]. Despite these high mechanical properties of TiB, further enhancement for the toughness is obtained by introducing metallic phase (β-Ti) as a matrix in MMC. In our previous research work [5, 106], it was proven that formation of two phases (β-Ti and TiB) can be acquired using TiB2 and Ti by reaction sintering, which is governed by the thermodynamically favorable simple reaction Ti + TiB2 → 2TiB. In the reaction sintering, the orthorhombic TiB whiskers are obtained from the TiB2 powder, where B diffuses in the Ti along B-B zigzag chain in [010] direction allowing the whisker growth with high aspect ratio. Other possible phases formed during the simple reaction sintering between Ti and B are TiB2 and Ti3B4. However, in this study, the reaction sintering completion has been achieved eliminating formation of TiB2 phase. Moreover, Ti3B4 phase formation can be only achieved at high elevated temperatures ~ 1690 °C [89]. The hypothesis of this research study is to achieve high toughness and ductility in the MMC provided from the metallic phase by introducing (β-Ti). Titanium has two crystal structures, namely, HCP packed α-Ti, and β-Ti, which has a BCC crystal structure. According to the Ti–B binary phase diagram, α-Ti is stable at room temperature and β-Ti 142 is stable at high temperatures ~ 883 °C. In this study, α-Ti phase has some challenges that made the phase undesirable due to limited slip systems in comparison with β-Ti. Interstitial elements such as O, N, and C in very small amounts can decrease the ductility and toughness significantly in the MMC [117]. However, β-Ti phase has more slip systems, which increase ductility and toughness, therefore, this phase is the point of interest for this research study. Since β-Ti phase is not stable at room temperature, adding β-stabilizing elements is required to achieve stability at room temperature. The objectives of this research study is to design and synthesize MMCs with discontinuous TiB whiskers as reinforcements with high aspect ratio and uniformly distributed throughout the volume of MMC. EFAS process produces the desired MMC in one step process rapidly. Producing MMC containing ductile metallic phase (β-Ti) stable at room temperature with low volume fraction of ceramic phase (TiB). This research study outlines the importance of processing through regress testing that contains characterization techniques and mechanical testing to investigate the novel class of MMC. 7.2 Design and Experimental Procedure 7.2.1 Design of MMC Composition and Processing The MMC that is the point of interest in this research study has no standard composition information to synthesize such as processing temperature and pressure, which are required to obtain a dense homogenous MMC structure. The importance of the processing temperature lies within two main points: first, to avoid excessive liquid phase formation that will lead to permanent die damage due to the oozing out of the liquid phase while under pressure; second, to achieve fully dense material. It is better to be as close as 143 possible to the liquid phase region by knowing the exact eutectic reaction point. The eutectic reaction point advantage was utilized in this research study by adding Fe to the MMC composition. The Ti–Fe binary phase diagram system shows that Fe has eutectic temperature of ~1085 °C at ~ 28 mol. % Fe [84]. In addition to that, Al element has been added to enhance the densification further due to the low melting point ~ 660 °C, and to eliminate the formation of the intermetallic TiFe phase. Also, Mo has been added to stabilize the β-Ti phase due to the coherency between the Mo-Ti resulting in no formation of intermetallic phase based on the Ti-Mo binary phase diagram [84]. CALPHAD approach has been used to aid in determining the optimum processing conditions (temperature and pressure) for the targeted MMC composition for this research study. This modeling approach can aid in determining liquid phase field, which helps in reaction sintering process for any system. Due to the previous reason, the pseudo-binary system was designed by CALPHAD approach. Thermo-Calc software has been used to model the phase stability regions in the Ti-B-Mo-Fe-Al system. The input data have been imported from the thermodynamic data base (Alloy solutions Database v6). Figure 7.1 (a-d) illustrates the pseudo-binary phase diagram of the Ti-B-Mo-Fe-Al system. The percentage mole composition is 10 % B, 3.7 % Mo, 4.3 % Fe, 2.9 % Al, and 79.1 % Ti. The MMC composition were designed to be close to the liquidus line (the red dot) to take the advantage of low temperature processing while stabilizing β-Ti phase. It is evident from the diagrams that the phase field region (β-Ti+TiB) is wider due to the stabilization of β-Ti phase. Also, the widening phase field region (β-Ti+TiB) indicates that there is no possible intermetallic phase formation. The processing temperature was chosen to be 1275 °C below the solidus line to allow solid state reaction enabling the maximum 144 densification possible. 7.2.2 Materials and Experimental Method In this research study, the MMC was synthesized from Ti, TiB2, Mo, Fe, and Al. The powder sources are as follows: Ti (Puris, West Virginia. 99.7 % pure, average particle size of 31 μm), TiB2 (Momentive, New York. 99.7 % pure, average particle size of 14 μm), Fe powder (Alfa Aesar, Massachusetts. 99.9 % pure, average particle size of 6 μm), Mo powder (Alfa Aesar, Massachusetts. 99.95 % pure, average particle size of 11 μm), and Al (Alfa Aesar, Massachusetts. 99.5 % pure, average particle size of 44 μm). The composition has been chosen such that Ti and TiB2 will yield low volume fraction of TiB (TiB < 20 % Vf percentage). Since the interest lies within β-Ti phase, Fe and Mo powders have been selected to stabilize the phase. Aluminum is an α-Ti stabilizer that has been added in a very small amount (2.9 mol. %) in comparison with Fe and Mo amounts. The addition of Al helps in lowering the melting temperature, which can help in densification due to formation of small amount of liquid phase during the process. The powders were ball milled for 24 hours in a container having Ar gas preventing any contaminations. After mixing, the mixture was loaded carefully into the cylindrical graphite die. The EFAS system (Thermal Technologies Model 10-4) was used to synthesis the MMC sample. The EFAS parameters were as follows: a heating rate of 50 °C/min and a holding time of 30 minutes to reach the target temperature of 1275 °C under a pressure of 10 MPa. After processing, the power was turned off, and the sample was cooled down to room temperature at the cooling rate of 22.5 °C/min, which is shown from the EFAS program. 145 The ingot that has been used for MMC analysis has dimensions of diameter 20 mm and thickness of 9 mm. The MMC density was obtain via Archimedes principle method. The ingot has been cut by a diamond saw to analyze the cross-section. Samples were prepared for metallographic analysis as follows: regress polishing followed by fine polishing. Extra fine polishing was performed by using colloidal silica (0.5 μm) to obtain mirror–like polished surface. The sample for optical microscopy was etched by Kroll’s reagent. The un-etched samples were used for hardness evaluation (Vickers hardness tester) with 1 kg load and holding time of 15 seconds and X-ray diffraction (XRD) analysis (Rigaku, Miniflex 600, Japan) with a scan rate of 0.5 °/min. The ingot that has been used for mechanical testing has the dimensions of diameter 60 mm and thickness 12 mm. The ingot was machined to produce the samples that are tested for flexural strength, fracture toughness, and tensile testing following ASTM standard. The samples dimensions for flexural strength are 45 × 4 × 3 mm3 (ASTM C116113). The samples dimensions for fracture toughness are 45 × 4 × 3 mm3 (ASTM C142110). Four-point bend configuration with displacement rates of 0.5 and 0.18 mm/min, respectively [86, 87], was used. The samples used for tensile testing have a diameter of 3.175 mm, with a total length of 50.80 mm and a gauge length of 12.70 mm (ASTM E8/E8M-15a) with strain rate of 2 × 10−4 /s. 7.2.3 Volume Fraction Calculation The MMC volume fraction calculations have been evaluated by using direct comparison method. The direct comparison method has been developed by Sahay et al. [5] to fit our research interests in Ti-TiB composites. This method yields relative volume 146 fractions of Ti and TiB phases by computing the highest integrated intensity for the nonoverlap peaks imported from the XRD pattern. The formula is given by: Vf = R RTi ITiB Ti ITiB + RTiB ITi (7.1) where I is the integrated intensity of the target peak. The parameter R, for each phase, is given by R= |Fhkl |2 PL Vₒ (7.2) where Fhkl is the structure factor, p is a multiplicity factor, L is the Lorentz polarization factor, and Vₒ the volume of the unit cell. For a detailed explanation of this method, our previous work shows extensive details [58]. In this research study, the relative volume fraction percentage of the MMC was calculated to be 18 % TiB and 82 % β-Ti. 7.3 Results and Discussion 7.3.1 EFAS Processing In this research study, it has been found that the synthesized MMC has a relative density of 99.7 %, which is confirmed by the micrograph reflected upon porosity absence. In addition to this, the ram displacement profile confirms that fully dense samples have been achieved. Figure 7.2 (a-b) illustrates ram displacement with power input versus time recorded from the EFAS process for the MMC composition with a target temperature at 1275 °C. From Figure 7.2 (a), it can be seen that the profile of the ram displacement shows a sharp decrease before reaching steady-state, which indicates that the MMC has reached full densification. Figure 7.2 (b) shows that power input increases until the target temperature is reached, then it will drop sharply, which is followed by a steady-state power. 147 7.3.2 X-ray Diffraction Data The XRD pattern of the synthesized MMC illustrated in Figure 7.3 has the processing temperature of 1275 °C. From the XRD pattern, it is observed that only β-Ti and TiB phases are present. Also, it can be understood from the XRD pattern that the sintering reaction at 1275 °C for 30 minutes is fully completed due to the absence of TiB2 phase. Moreover, β-Ti stability is achieved at room temperature due to the presence of the high intense characteristic peak (110) β-Ti. 7.3.3 Microstructure The MMC microstructure is illustrated in Figure 7.4 (a-b). The microstructure of the MMC shows two phases (β-Ti and TiB) as matrix and two modes (primary whiskers with aspect ratio of 3-5 μm and clusters with average size 20 μm), respectively. High aspect ratio has been achieved due to the high mean-free-path through the volume of the MMC. Also, it can be seen that low volume fraction of TiB and high volume fraction of β-Ti has been achieved. 7.4 Mechanical Properties 7.4.1 Hardness Evaluation The hardness evaluation for the MMC is done by Vickers hardness tester. In the previous study, it has been found that the Vickers hardness value of the TiB phase is ~1800 kg/mm2 [60]. The Vickers hardness measurements for the β-Ti phase is difficult to obtain due to the fact that the size of the indentation tip (55 μm) is larger than β-Ti phase (~ 30 μm). Also, the TiB whiskers present around the β-Ti phase affect the hardness value due 148 to random orientation of the whiskers. According to the previous reasons, cumulative probability distribution function has been used to fit the obtained hardness data, as shown in Figure 7.5. The average value of hardness is 310–427 kg/mm2. The variation in hardness stems from the fact that the two phases (soft β-Ti phase and hard TiB phase) are randomly distributed throughout the MMC. From the plot, it is possible to extrapolate the hardness value of β-Ti phase, which was found to be 300 kg/mm2 by curve fitting using Weibull type function and by taking hardness of the TiB phase as ~ 1800 kg/mm2. H− Hβ P = 1 – exp [−C {H TiB − Hβ n } ] (7.3) where H is the hardness of the MMC and HTiB and Hβ are the hardness values of TiB and β-Ti phases, respectively. 7.4.2 Flexure Strength The MMC flexural strength is shown in Figure 7.6. The average ultimate bending stress value obtained for the flexural strength value is 1622 MPa and the average yield bending stress value is 1170 MPa for the MMC containing 18 % volume fraction of TiB. The fractography analysis has been done using SEM images illustrated by Figure 7.7 (a-i). The fractured surfaces revealed dimpled fracture throughout the cross-section of MMC. The majority of the fracture is shown to be ductile fracture due to the plastic deformation of the metallic β-Ti matrix. 149 7.4.3 Fracture Toughness The fracture toughness (KIC) has been initiated by a single edge cracked beam in specimens bending for the MMC containing volume fraction of 18% TiB. Figure 7.8 illustrates the fracture toughness test. The average value of K IC obtained is 23 MPa√𝑚. High fracture toughness and flexural strength were evidenced as high values in comparison to some previous studies presented in Table 7.1. Figure 7.9 (a-h) illustrates the fracture behavior during the crack propagation. The fracture pattern shows an extensive plastic deformation through the β-Ti matrix. It is also shown that the crack is stable, which indicates further crack propagation requires high stress to continue due to the MMC resistance. 7.4.4 Tensile Properties The MCC with 18% volume percentage of TiB (Ti-10B-3.7Mo-4.3Fe-2.96Al) was mechanically evaluated by tensile test. Figure 7.10 presents the tensile data for Ti-10B3.7Mo-4.3Fe-2.96Al. The graph shows a significant improvement in the strength of the MMC represented by the new composition (Ti-10B-3.7Mo-4.3Fe-2.96Al) of MMC. There is ductility observed and that can be explained from fracture surface micrograph. From the tensile data, the elastic modulus, ultimate tensile strength, yield strength, and elongation as 114 GPa, 1041 MPa, 910 MPa, and 0.27%, respectively, was evidence as a high value in comparison to titanium alloys presented in Table 7.2. Figure 7.11 (a-d) illustrates the fracture behavior during the tensile test in the MMC (Ti-10B-3.7Mo-4.3-Fe-2.96Al). The fracture surface observed in the SEM shows a ductile 150 behavior throughout the β-Ti matrix, but the whiskers distribution prevented the β-matrix to be connected in a way that promotes in high ductility. The whiskers have debonded the matrix, which prevents the tensile force from transferring throughout the β-matrix, creating high local stress concentration that led to material failure. This phenomena explains the reason that there is some not very high ductility observed. 7.5 Conclusions 1- CALPHAD approach by Thermo-Calc. has been used to design and synthesize the novel MMC. The targeted MMC Ti-B-Mo-Fe-Al system was synthesized with full density by EFAS successfully in a single step with stable phases and microstructure. 2- The MMC microstructure revealed uniform distribution of TiB phase with stable β-Ti phase at room temperature due to the addition of β-stabilizer elements (Fe and Mo). 3- Weibull distribution function was the best fit for the obtained data from the hardness test, which yields average hardness value of 375 kg/mm2 and the hardness value of 300 kg/mm2 for β-Ti. 4- The average value of the flexural strength is 1622 MPa and the dominant fracture behavior is ductile fracture. 5- The average value of fracture toughness is 23 MPa√m. The high value of fracture came from ductile β-Ti phase. 6- The average ultimate tensile strength is 1041 MPa, elastic modulus of 114 GPa, yield strength 910 MPa, and elongation of 0.27 %. Micrographic images show the presence of ductile β-Ti phase and the presence of brittle TiB phase. 151 (a) Figure 7.1 CALPHAD approach. (a-d) The pseudo-binary phase diagram of the Ti-BMo-Fe-Al system. 152 (b) Figure 7.1 Continued. 153 (c) Figure 7.1 Continued. 154 (d) Figure 7.1 Continued. 155 1275 °C 0 1400 6000 1200 5000 (b) 1000 -1 800 -2 600 -3 18 vol.% TiB -4 400 200 -5 0 0 1000 2000 3000 4000 5000 Time (sec) Power (W) Ram Displacement (mm) (a) Temperature (°C) 1 18 vol.% TiB 4000 3000 2000 1000 0 0 1000 2000 3000 4000 5000 Time (sec) Figure 7.2 EFAS process. (a-d) Ram displacement with power input v/s time recorded from the EFAS process. 20 30 40 50 60 70 2θ Figure 7.3 Room temperature X-ray diffraction of the synthesized MMCs. TiB (122) TiB (312) TiB (020) β-Ti (200) TiB (112) TiB (301) TiB (211) β-Ti (211) β-Ti (110) Ti-10B-3.7Mo-4.3-Fe-2.9Al TiB (210) TiB (102) TiB (111) TiB (011) TiB (200) Intensity (A.U) 156 80 157 Figure 7.4 Microstructures (a-b) of the synthesis MMC. 158 1.0 0.9 Cumulative Probability 0.8 0.7 0.6 0.5 0.4 0.3 0.2 18% TiB (C=220, n=1.9) 0.1 0.0 0 200 400 600 800 1000 1200 1400 1600 1800 Hardness (Hv) Kg/mm2 Figure 7.5 Cumulative probability distributions of hardness values. The lines represent fits to a Weibull-type function. 159 1750 Test 1 Strength, MPa 1500 Test 2 Test 3 1250 Test 4 Test 5 1000 750 500 250 0 0 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 1.8 Displacement, mm Figure 7.6 Flexural strength of the MMC. 160 Figure 7.7 SEM analysis. (a-i) SEM images for the fracture surface of MMC. 161 Figure 7.7 Continued. 162 Figure 7.7 Continued. 163 Figure 7.7 Continued. 164 Figure 7.7 Continued. 165 25 Average KIC K (MPa.√m) 20 Test 1 Test 2 15 Test 3 Test 4 10 5 0 0 0.1 0.2 0.3 0.4 Displacement (mm) Figure 7.8 Fracture toughness of MMC. 0.5 166 Figure 7.9 Fracture surface. (a-h) Fracture behavior during the propagation of a crack through the MMC structure. 167 Figure 7.9 Continued. 168 Figure 7.9 Continued. 169 Figure 7.9 Continued. 170 Figure 7.10 Tensile data of MMC Ti-10B-3.7-Mo-4.3Fe-2.96Al. 171 Figure 7.11 Tensile fracture surface. (a-d) The behavior of MMC Ti-10B-3.7Mo4.3Fe-2.96Al. 172 Figure 7.11 Continued. 173 Table 7.1 Fracture toughness of Ti-10B-3.7Mo-4.3-Fe-2.96Al v/s titanium alloys MMC MMC Vf of Sintering Fracture Flexural Temperature Toughness Strength (°C) (KIC) (MPa) TiB Ref. Ti-TiB 0.06 1250 13.9 1443 [118] Ti–FeMo–B 0.10 1000 9.6 1560 [80] Ti-TiB 0.58 1400 9.3 608 [119] This study 0.18 1275 23 1622 Author Table 7.2 Tensile properties of Ti-10B-3.7Mo-4.3Fe-2.96Al v/s titanium alloys MMC Ultimate MMC Vf of Tensile TiB Strength Elongation (%) Yield Elastic Strength Modulus (MPa) (GPa) (MPa) Ref. Ti-6Al-4V 0.10 1000 0.25 - 137 [78] Ti 0.15 903 0.4 842 139 [120] Ti-5Al-2.5Fe 0.15 1092 0 * 151 [121] This study 0.18 1041 0.27 910 114 Author REFERENCES [1] K. K. Chawla, Composite Materials: Science and Engineering, Alabama: Springer, 2013. [2] R. W. Nichols, Advanced Materials by Design Advisory Panel, New York: NTIS, 1998. [3] J. R. Aikin, "The mechanical properties of in situ composites," The Journal of The Minerals, vol. 49, no. 8, pp. 35-39, 1997. [4] R. Orrù, R. Licheri, A. M. Locci, A. Cincotti and G. 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