Applied Science, Faculty of
Engineering, School of (Okanagan)
DSpace
UBCO
Pillari, Lava Kumar
Doctor of Philosophy - PhD
University of British Columbia
The automotive industry is globally transitioning towards lighter vehicles to meet stringent fuel standards and reduce CO₂ emissions. The B319 aluminum (Al) alloy finds widespread use in automotive powertrain components due to its low density, good strength, and excellent castability. However, the poor tribological properties of the B319 Al alloy hinder its broader application in powertrain components. Graphene has been considered a promising reinforcement for Al matrices due to its ability to enhance multiple properties simultaneously while maintaining the lightweight characteristics of the alloy. However, graphene agglomeration and its poor wettability pose significant processing challenges, particularly in the casting process, thereby limiting the advancement of Al-graphene composites and their industrial adoption. Therefore, this study aims to fabricate B319 Al alloy-graphene (0-0.7 wt.%) composites utilizing a novel distribution approach. This approach involves the fabrication of composite master alloy (F357-5 wt.% graphene) through a flake powder metallurgy approach involving ball milling and spark plasma sintering. Subsequently, the composite master alloy was introduced into the B319 Al alloy melt to fabricate B319-graphene composites via stir casting. The effect of graphene on the microstructure, hardness, tribological behavior, electrical conductivity, and thermal conductivity of the B319 Al alloy was investigated. The results indicate that the addition of graphene decreased the grain size of the B319 Al alloy, primarily attributed to the additional nucleation sites provided by the presence of graphene. Moreover, adding graphene refined the Si phase and changed its morphology from a needle-like to a more fibrous structure. Consequently, the hardness of B319 Al alloy was enhanced by 20% with the addition of 0.1 wt.% graphene, but it deteriorated with graphene additions beyond 0.3 wt.%. Regardless of the load, the B319-0.3 wt.% graphene composite exhibited the best anti-friction and wear response compared to the base alloy. This improvement was attributed to the formation of a stable mechanically mixed layer (MML) or graphene-containing tribolayer at the contact surfaces. Wear in the base alloy primarily occurred via MML delamination through subsurface crack propagation, while wear in the B319-0.3 wt.% graphene composite was driven by MML fragmentation, accompanied by adhesive, oxidative, and abrasive wear mechanisms.
FABRICATION OF GRAPHENE-BASED MASTER ALLOYS FOR THE DEVELOPMENT OF B319 ALUMINUM ALLOY-GRAPHENE COMPOSITES WITH ENHANCED TRIBOLOGICAL PROPERTIES by Lava Kumar Pillari B.Tech, Rajiv Gandhi University of Knowledge Technologies-Nuzvid, India, 2014 M.Tech (By Research), Vellore Institute of Technology-Chennai, India, 2018 A DISSERTATION SUBMITTED IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY in THE COLLEGE OF GRADUATE STUDIES (Mechanical Engineering) THE UNIVERSITY OF BRITISH COLUMBIA (Okanagan) June 2024 © Lava Kumar Pillari, 2024 ii The following individuals certify that they have read, and recommend to the College of Graduate Studies for acceptance, a dissertation entitled: Fabrication of Graphene-Based Master Alloys for the Development of B319 Aluminum Alloy-Graphene Composites with Enhanced Tribological Properties submitted by Lava Kumar Pillari in partial fulfillment of the requirements of the degree of Doctor of Philosophy . Dr. Lukas Bichler, School of Engineering, UBC Okanagan Supervisor Dr. Mohammad Arjmand, School of Engineering, UBC Okanagan Supervisory Committee Member Dr. Jian Liu, School of Engineering, UBC Okanagan Supervisory Committee Member Dr. Kasun Hewage, School of Engineering, UBC Okanagan University Examiner Dr. Gobinda Saha, Mechanical Engineering, University of New Brunswick External Examiner iii Abstract The automotive industry is globally transitioning towards lighter vehicles to meet stringent fuel standards and reduce CO2 emissions. The B319 aluminum (Al) alloy finds widespread use in automotive powertrain components due to its low density, good strength, and excellent castability. However, the poor tribological properties of the B319 Al alloy hinder its broader application in powertrain components. Graphene has been considered a promising reinforcement for Al matrices due to its ability to enhance multiple properties simultaneously while maintaining the lightweight characteristics of the alloy. However, graphene agglomeration and its poor wettability pose significant processing challenges, particularly in the casting process, thereby limiting the advancement of Al-graphene composites and their industrial adoption. Therefore, this study aims to fabricate B319 Al alloy-graphene (0-0.7 wt.%) composites utilizing a novel distribution approach. This approach involves the fabrication of composite master alloy (F357-5 wt.% graphene) through a flake powder metallurgy approach involving ball milling and spark plasma sintering. Subsequently, the composite master alloy was introduced into the B319 Al alloy melt to fabricate B319-graphene composites via stir casting. The effect of graphene on the microstructure, hardness, tribological behavior, electrical conductivity, and thermal conductivity of the B319 Al alloy was investigated. The results indicate that the addition of graphene decreased the grain size of the B319 Al alloy, primarily attributed to the additional nucleation sites provided by the presence of graphene. Moreover, adding graphene refined the Si phase and changed its morphology from a needle-like to a more fibrous structure. Consequently, the hardness of B319 Al alloy was enhanced by 20% with the addition of 0.1 wt.% graphene, but it deteriorated with graphene additions beyond 0.3 wt.%. Regardless of the load, the B319-0.3 wt.% graphene composite exhibited the best anti-friction and wear response compared to the base alloy. This improvement was attributed to the formation of a stable mechanically mixed layer (MML) or graphene-containing tribolayer at the contact surfaces. Wear in the base alloy primarily occurred via MML delamination through subsurface crack propagation, while wear in the B319-0.3 wt.% graphene composite was driven by MML fragmentation, accompanied by adhesive, oxidative, and abrasive wear mechanisms. iv Lay Summary The implementation of new fuel economy standards by governments internationally, alongside increasing consumer demand for more fuel-efficient vehicles, has increased pressure on automakers to improve fuel efficiency and reduce greenhouse gas emissions. Automakers have been exploring various approaches to meet fuel economy standards and improve fuel efficiency, with one strategy involving the substitution of heavier vehicle components with lighter alternatives. A significant emphasis has been placed on the vehicle’s engine block, as it is the heaviest component and offers the greatest potential for weight reduction. However, current engine materials often exhibit inadequate wear properties. Therefore, this dissertation focuses on improving the wear properties of current engine material through the addition of graphene. The newly developed material, known as the B319-graphene composite, demonstrates superior wear resistance. The findings of this dissertation could pave the way for the development of lightweight and high-performance wear-resistant components for next-generation automotive applications. v Preface This dissertation presents the original research work conducted by the author, with several parts of the study developed in collaboration with individuals from Canadian and Indian research facilities. These collaborators include Prof. Lukas Bichler and Mr. Kyle Lessoway from the School of Engineering at the University of British Columbia-Okanagan (Canada); Dr. Anthony Lombardi and Dr. Glenn Byczynski from Nemak, Windsor (Canada); Dr. Colin van der Kuur from Zentek Ltd., Guelph (Canada); Prof. B.S. Murty from Indian Institute of Technology, Hyderabad (India); and Dr. S.V.S. Narayana Murty from Liquid Propulsion Systems Centre, Thiruvananthapuram (India). Chapter 1, Chapter 3, Chapter 4, Chapter 5, and Chapter 6 of this dissertation include components from: • L.K. Pillari, K. Lessoway, C. van der Kuur, A. Lombardi, G. Byczynski, L. Bichler, Effect of graphene on the microstructure, thermal conductivity, and tribological behavior of cast B319 Al alloy, Wear 538–539 (2024) 205201. • L.K. Pillari, K. Lessoway, L. Bichler, Reciprocating dry sliding friction and wear behavior of B319 aluminum alloy-graphene composites, Tribology International 192 (2024) 109334. • L.K. Pillari, K. Lessoway, L. Bichler, Improving the wear behavior of as-cast B319 aluminum alloy through graphene addition, Proceedings of the 63rd Conference of Metallurgists (COM 2024), Accepted. I conducted the experimental work described in the aforementioned papers and wrote the original drafts of the manuscripts. K. Lessoway participated in the experimental work and contributed to the manuscript. C. van der Kuur provided materials and discussed the results. A. Lombardi and G. Byczynski contributed to the discussion of results and reviewed and edited the manuscript. L. Bichler secured funding, participated in experimental work, provided equipment and guidance, discussed results, and reviewed and edited all the manuscripts. Chapter 2 of this dissertation includes components from: vi • P.L. Kumar, A. Lombardi, G. Byczynski, S.V.S. Narayana Murty, B.S. Murty, L. Bichler, Recent advances in aluminium matrix composites reinforced with graphene-based nanomaterial: A critical review, Progress in Materials Science 128 (2022) 100948. • L.K. Pillari, K. Lessoway, L. Bichler, Carbon nanotube and graphene reinforced magnesium matrix composites: A state-of-the-art review, Journal of Magnesium and Alloys 11 (2023) 1825–1905. I conducted the literature review, data processing, and analysis, and wrote the original drafts of the manuscripts. K. Lessoway participated in data collection and contributed to the manuscript. A. Lombardi, G. Byczynski, S.V.S. Narayana Murty, and B.S. Murty reviewed and edited the manuscript. L. Bichler secured funding, provided equipment and guidance, discussed results, and reviewed and edited all the manuscripts. Some parts of the dissertation were also presented at the following conferences: • L.K. Pillari, K. Lessoway, C. van der Kuur, A. Lombardi, G. Byczynski, and L. Bichler, Effect of graphene on the microstructure and properties of as-cast B319 aluminum alloy, Canadian Materials Science Conference (CMSC), June 22-24, 2022, University of Toronto, Canada. • L.K Pillari, L. Bichler, Reciprocating dry sliding friction and wear response of B319 Al alloy-graphene composites fabricated by casting, International Conference on Processing and Characterization of Materials (ICPCM), December 8-10, 2023, National Institute of Technology, Rourkela, India. vii Table of Contents Abstract ......................................................................................................................................... iii Lay Summary ............................................................................................................................... iv Preface ............................................................................................................................................ v Table of Contents ........................................................................................................................ vii List of Tables ................................................................................................................................. xi List of Figures .............................................................................................................................. xii List of Abbreviations .................................................................................................................. xxi List of Symbols ......................................................................................................................... xxiii List of Elements ......................................................................................................................... xxv List of Units .............................................................................................................................. xxvi Acknowledgments ................................................................................................................... xxvii Chapter 1 : Introduction .............................................................................................................. 1 1.1 Background ...................................................................................................................... 1 1.2 Objectives of the Work ..................................................................................................... 4 1.3 Scope of the Dissertation.................................................................................................. 4 1.4 Novelty and Originality of the Research .......................................................................... 6 Chapter 2 : Literature Review ..................................................................................................... 7 2.1 Introduction ...................................................................................................................... 7 2.1.1 The Need for Lightweighting...................................................................................... 7 2.1.2 The Need for Al-Graphene based Nanocomposites (AGNCs) ................................. 10 2.1.3 Synthesis, Structure, and Properties of GNs (Pristine Graphene, GO, and rGO) ..... 14 2.2 Challenges in Developing AGNCs ................................................................................. 17 2.2.1 Agglomeration of GN ............................................................................................... 17 2.2.2 Poor Wettability and Al4C3 Formation ...................................................................... 19 viii 2.3 Processing of AGNCs..................................................................................................... 22 2.3.1 Powder Metallurgy (PM) .......................................................................................... 22 2.3.1.1 Incorporation of GN into the Al Matrix ............................................................ 22 2.3.1.1.1 Ball Milling (BM) ......................................................................................... 23 2.3.1.1.2 Solution-Assisted Mixing (SAM) ................................................................. 28 2.3.1.2 Consolidation of AGNC Powder ...................................................................... 30 2.3.1.2.1 Pressureless Sintering (PLS) ......................................................................... 30 2.3.1.2.2 Pressure-Assisted Sintering (PAS) ................................................................ 31 2.3.2 Casting ...................................................................................................................... 33 2.3.3 Friction Stir Processing (FSP) .................................................................................. 35 2.3.4 High-Pressure Torsion (HPT) ................................................................................... 37 2.3.5 Selective Laser Melting (SLM) ................................................................................ 39 2.3.6 Electrical Resistance Heat-Assisted Mechanical Impregnation (ERHAMI) ............ 41 2.4 Properties of AGNCs ...................................................................................................... 42 2.4.1 Mechanical Properties ............................................................................................... 43 2.4.1.1 Strengthening Mechanisms in AGNCs ............................................................. 47 2.4.2 Electrical and Thermal Properties ............................................................................. 49 2.4.3 Tribological Properties .............................................................................................. 51 2.5 Chapter Summary ........................................................................................................... 54 2.6 Problem Definition and Approach.................................................................................. 56 Chapter 3 : Materials and Experimental Details ..................................................................... 58 3.1 Materials ......................................................................................................................... 58 3.2 Fabrication of Composite Master Alloys ....................................................................... 58 3.3 Fabrication of Cast B319-Graphene Composites ........................................................... 59 3.4 Characterization of B319-Graphene Composites ........................................................... 60 ix 3.4.1 Spark Optical Emission Spectrometer (OES) ........................................................... 60 3.4.2 Friction and Wear Measurement ............................................................................... 61 3.4.3 Microstructural Analysis ........................................................................................... 63 3.4.3.1 X-ray Diffraction Analysis (XRD) ................................................................... 63 3.4.3.2 Raman Spectroscopy ......................................................................................... 64 3.4.3.3 Laser Particle Size Analysis.............................................................................. 64 3.4.3.4 Optical Microscopy ........................................................................................... 64 3.4.3.5 Scanning Electron Microscopy (SEM) ............................................................. 65 3.4.3.6 Electron Probe Microanalysis (EPMA) ............................................................ 66 3.4.3.7 Density Measurement ....................................................................................... 66 3.4.4 Thermal Analysis ...................................................................................................... 67 3.4.4.1 Differential Thermal Analysis/Thermogravimetric Analysis (DTA/TGA) ...... 67 3.4.4.2 Thermal Conductivity Measurement ................................................................ 68 3.4.5 Electrical Conductivity Measurement ....................................................................... 68 3.4.6 Hardness Measurement ............................................................................................. 68 Chapter 4 : Effect of Graphene on the Microstructure, Hardness, and Thermal Conductivity of Cast B319 Al Alloy .................................................................................................................. 69 4.1 Microstructural Analysis of Graphene ........................................................................... 69 4.2 Microstructural Analysis of Ball Milled F357-Graphene Powder Mixture ................... 72 4.3 Microstructural Analysis of SPSed Composite Master Alloy ........................................ 75 4.4 Microstructure and Phase Analysis of B319-Graphene Composites ............................. 80 4.5 Hardness, Electrical, and Thermal Conductivity of B319-Graphene Composites ......... 91 4.6 Chapter Summary ........................................................................................................... 93 Chapter 5 : Effect of Graphene on the Friction and Wear Behavior of B319-Graphene Composites ................................................................................................................................... 96 5.1 Frictional Behavior of B319-Graphene Composites ...................................................... 96 x 5.2 Wear Behavior of B319-Graphene Composites ........................................................... 103 5.2.1 Wear Track Analysis ............................................................................................... 103 5.2.2 Analysis of Wear Scar on Counterbody .................................................................. 106 5.2.3 Wear Properties ........................................................................................................110 5.2.4 Dissipated Frictional Energy ....................................................................................114 5.2.5 Wear Mechanisms ....................................................................................................115 5.2.5.1 Worn Surface Analysis ................................................................................... 115 5.2.5.2 Wear Debris Analysis ..................................................................................... 129 5.2.5.3 Subsurface Analysis of Wear Track ............................................................... 133 5.2.5.3.1 Subsurface Microstructure of Wear Track .................................................. 133 5.2.5.3.2 Subsurface Hardness of Wear Track ........................................................... 148 5.3 Chapter Summary ......................................................................................................... 151 Chapter 6 : Conclusions and Future Work............................................................................. 156 6.1 Conclusions .................................................................................................................. 156 6.1.1 Composite Master Alloy ......................................................................................... 156 6.1.2 Microstructure, Hardness, Electrical, and Thermal Conductivity of B319-Graphene Composites .......................................................................................................................... 157 6.1.3 Friction Behavior of B319-Graphene Composites ................................................. 158 6.1.4 Wear Behavior of B319-Graphene Composites ...................................................... 158 6.2 Challenges Faced During the Research ........................................................................ 159 6.3 Limitations of the Research.......................................................................................... 160 6.4 Recommendations for Future Work ............................................................................. 160 References .................................................................................................................................. 162 xi List of Tables Table 2.1. Trend in average weight of vehicles in the U.S. (Source: U.S. Environmental Protection Agency, 2023 EPA Automotive Trends Report). ............................................................................ 9 Table 2.2. Properties of graphene-based nanomaterials (Pristine graphene, GO, and rGO). ....... 15 Table 3.1. Chemical composition of F357 Al alloy, in wt.%. ...................................................... 60 Table 3.2. Chemical composition of B319-graphene composites, in wt.%. ................................. 60 Table 4.1. The average particle size of F357-graphene powder. .................................................. 75 Table 4.2. The average composition of secondary phases in B319 Al alloy, obtained by EDS analysis. ......................................................................................................................................... 81 Table 4.3. EPMA-WDS chemical analysis of phases in as-cast B319 Al alloy. .......................... 83 Table 4.4. Microstructural characteristics, relative density, and hardness of B319-graphene composites..................................................................................................................................... 94 Table 5.1. Average surface roughness of polished samples as measured by a profilometer. ....... 99 Table 5.2. Effect of graphene on the average wear track depth and maximum depth of B319-graphene composites tested under 5 N and 10 N loads. ............................................................. 106 Table 5.3. Effect of normal load on the Hertzian contact radius and maximum contact pressure.......................................................................................................................................................113 Table 5.4. Average running-in and steady-state CoF of B319-graphene (0-0.7 wt.%) composites...................................................................................................................................................... 152 Table 5.5. Average overall (running-in and steady-state regime) CoF of B319-graphene (0-0.7 wt.%) composites. ....................................................................................................................... 152 Table 5.6. Average wear volume (mm3) of B319-graphene composites evaluated by weight loss and laser microscope method. ..................................................................................................... 154 Table 5.7. Average specific wear rate (mm3/N.m×10-4) of B319-graphene composites evaluated by weight loss and laser microscope method. .................................................................................. 154 xii List of Figures Figure 1.1. Flow chart illustrating the research methodology and key aspects explored in this thesis. .............................................................................................................................................. 5 Figure 2.1. Contribution of various sectors to greenhouse gas (GHG) emissions in Canada. (Source: National Inventory Report 1990–2018: Greenhouse gas sources and sinks in Canada). . 8 Figure 2.2. Trend in weight percentage of steels, cast iron, and aluminum in light-duty vehicles [7]. ................................................................................................................................................. 10 Figure 2.3. Photograph of V-8 engine block with cast-in iron liners [8]. .................................... 12 Figure 2.4. A schematic showing the crystal structure of graphene-based materials (a) pristine graphene, (b) graphene oxide (GO), and (c) reduced graphene oxide (rGO). The inset shows the C-C bond length (0.142 nm) and type (covalent) in pristine graphene. ....................................... 16 Figure 2.5. High-resolution TEM image of vacancy defect in single-layer graphene, (a) pristine graphene, (b) graphene with vacancy defect, and (c) with atomic configuration [121]. .............. 16 Figure 2.6. TEM image of (a) Al+0.5 wt.% GNF, showing a uniform distribution of GNF [124]; (b) Al+1.0 wt.% GNF, showing GNF agglomeration in Al matrix [124]; (c) Al+1 wt.% GNP, showing GNP agglomeration in Al matrix [125]; and (d) scanning electron microscope (SEM) image of fracture surface of Al+1.5 wt.% rGO, showing agglomeration of rGO in Al matrix [47]........................................................................................................................................................ 18 Figure 2.7. The wetting behavior of Al on (a) graphite substrate, (b) Cu-coated graphite substrate [128], and (c) Ni-coated graphite substrate [129] captured during the sessile drop test. ............. 19 Figure 2.8. TEM image showing needle shape Al4C3 (a) in Al5083+GNP composite, (d) Al5083+GO composite; (c) TEM image of the interface between GNP and Al5083, (e) high-resolution image of the selected area in (d) and (b) diffraction pattern of the Al [130]. .............. 20 Figure 2.9. Contribution (%) of various processing methods used by researchers to develop AGNCs. ......................................................................................................................................... 22 Figure 2.10. A schematic showing the agglomeration of graphene caused by van der Waals and interaction forces between the graphene sheets. ........................................................................... 23 Figure 2.11. A schematic illustrating the (a) movement of milling media, (b) fracture, and (c) cold-welding mechanism during the milling of Al-graphene powder mixture in a planetary ball mill. 24 xiii Figure 2.12. Effect of milling time on the dispersion of GNS in Al matrix (a) 4 hours, (b) 6 hours. (c) 8 hours. and (d) 10 hours [48]. ................................................................................................ 25 Figure 2.13. HRTEM image of Al6063-0.3 vol.% GNS milled for (a) 1h, showing the thin layer of amorphous Al2O3 between GNS and Al, and (b) 3h, showing the direct contact between Al and GNS [140]. .................................................................................................................................... 26 Figure 2.14. Milling parameters used by various researchers to disperse GNs in Al matrix (a) milling speed, (b) ball-to-powder-ratio, and (c) milling time. The data for this plot were extracted from the literature and plotted by the author. ................................................................................ 27 Figure 2.15. (a) SEM image showing uniform adsorption of GO on Al flake, and (b) EDS of the selected area in (a) [46]. ................................................................................................................ 29 Figure 2.16. Electron backscatter diffraction (EBSD) grain boundary maps of (a) Al-Mg alloy, (b) FSPed Al-Mg alloy, and (c) FSPed Al-Mg alloy+3 vol.% graphene [177]. ............................ 36 Figure 2.17. STEM image of Al-5wt.% graphene composite HPT processed at (a, b) 298 K for 1 turn, showing the presence of graphene at grain boundary; (c, d) 473 K for 5 turns, showing the presence of graphene at both the grain boundary and within grain [188]. ................................... 38 Figure 2.18. A schematic of selective laser melting (SLM) setup used for the synthesis of AGNCs [197]. ............................................................................................................................................. 40 Figure 2.19. A schematic showing the experimental setup of the mechanical impregnation technique used for producing Al-GNP surface composites [202]. ................................................ 41 Figure 2.20. Relative density of GN-reinforced Al matrix composites processed by pressureless sintering (conventional, microwave) and pressure-assisted sintering (hot pressing, hot isostatic pressing, spark plasma sintering). The data for this plot were extracted from the literature and plotted by the author...................................................................................................................... 43 Figure 2.21. Percentage of improvement or deterioration in mechanical properties of AGNCs (processed by different techniques) as a function of GN content: (a) hardness, (b) elastic modulus, (c) yield strength, (d) ultimate tensile strength, and (e) ultimate compressive strength. The data for this plot were extracted from the literature and plotted by the author. ......................................... 45 xiv Figure 2.22. Optimum GN content reported in individual studies of AGNCs fabricated by PM, casting and FSP. The data for this plot were extracted from the literature and plotted by the author........................................................................................................................................................ 46 Figure 2.23. (a) Grain size and (b) Percentage of increase or decrease in grain size of AGNCs as a function of GN content, (c) Trend in reported strengthening mechanisms (LT: Load transfer strengthening; OS: Orowan strengthening; DS: Dislocation strengthening; GR: Grain refinement strengthening) by researchers in AGNCs, (d) Average grain size of AGNCs (processed by different techniques) at an optimum content of GN reported in the individual studies. The data for this plot were extracted from the literature and plotted by the author. ....................................................... 49 Figure 2.24. Percentage of increase or decrease in (a) CoF and (b) wear rate of AGNCs as a function of GN content. The data for this plot were extracted from the literature and plotted by the author............................................................................................................................................. 53 Figure 2.25. A schematic depicting the GNs effects during sliding wear (a) wrapping effect, (b) bridging effect, (c) interlayer sliding, and (d) formation of rollers [35]. ...................................... 54 Figure 3.1. A schematic showing a hybrid distribution approach used in the present study to fabricate B319-graphene composites. ........................................................................................... 59 Figure 3.2. Photograph of the pin-on-disk tribometer used in this study, showing its main parts........................................................................................................................................................ 61 Figure 3.3. Schematic representation of a wear track, showing the position of the cross-section examined in this work and sliding direction. ................................................................................ 65 Figure 4.1. SEM images of graphene at (a) low, (b,c) high magnification and EDS mapping of Fig. (c), and the EDS data of region (A) marked in Fig. (c). ........................................................ 70 Figure 4.2. (a) TGA curves for proximate analysis, (b) Particle size distribution curve, (c) Raman spectra, and (d) XRD pattern of raw graphene powder. ............................................................... 71 Figure 4.3. SEM images of F357 Al alloy powder (a) As received, (b) 1-hour ball milled, and (c) EDS area mapping of Figure (a). .................................................................................................. 72 Figure 4.4. SEM images and EDS area elemental mapping of ball milled powder (a) F357-2.5 wt.% graphene, (b) F357-5 wt.% graphene, and (c) F357-10 wt.% graphene. ............................. 74 Figure 4.5. Particle size distribution curves for F357 Al alloy and F357-graphene powder. ....... 75 xv Figure 4.6. Optical microstructure of SPSed composite master alloy: (a) F357-2.5 wt.% graphene, (b) F357-5 wt.% graphene, and (c) F357-10 wt.% graphene; (d) Relative density and hardness of composite master alloys. ............................................................................................................... 76 Figure 4.7. Backscattered electron (BSE) images and corresponding EDS area maps of SPSed composite master alloys (a) F357-2.5 wt.% graphene, (b) F357-5 wt.% graphene, and (c) F357-10 wt.% graphene............................................................................................................................... 77 Figure 4.8. Photographs of (a) unmelted F357-10 wt.% graphene composite master alloy and (b) clusters of ball-milled F357-10 wt.% graphene powder. .............................................................. 79 Figure 4.9. (a) XRD pattern of as-received F357 alloy powder and SPSed composite master alloy, and (b) DTA curves of as-received F357 alloy powder and F357-5 wt.% graphene milled powder........................................................................................................................................................ 79 Figure 4.10. Optical microstructure of B319 Al alloy in (a) ingot and (b) as-cast condition; (c) XRD patterns of B319 Al alloy in ingot and as-cast condition..................................................... 81 Figure 4.11. XRD pattern of as-cast B319 Al alloy-graphene (0-0.7 wt.%) composites. ............ 82 Figure 4.12. SEM image and EDS area elemental mapping of B319 Al alloy. ........................... 82 Figure 4.13. SEM-BSE image of as-cast B319 Al alloy and corresponding EPMA-WDS elemental mapping......................................................................................................................................... 83 Figure 4.14. (a) SEM image and (b) Optical microstructures of B319 Al alloy, showing the precipitation of Al2Cu on Al15(Fe, Mn)3Si2 and Si particles. ........................................................ 84 Figure 4.15. The possible solidification reactions and sequence in B319 Al alloy. ..................... 85 Figure 4.16. Optical microstructure of cast samples captured at 100x (a) B319, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, (f) B319-0.7 wt.% graphene.......................................................................................................... 86 Figure 4.17. Macrostructure of cast samples etched with Keller's reagent: (a) B319, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene and (f) B319-0.7 wt.% graphene. .................................................................................. 87 Figure 4.18. Relative density of B319 Al alloy-graphene composites. ........................................ 88 xvi Figure 4.19. Optical microstructure of cast samples: (a) B319, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene and (f) B319-0.7 wt.% graphene......................................................................................................................... 89 Figure 4.20. Effect of graphene on the morphological characteristics of Si: (a) Length, (b) Width, and (c) Aspect ratio. ...................................................................................................................... 90 Figure 4.21. (a) SEM image and (b) EDS elemental mapping of B319-0.1 wt.% graphene composite, showing the presence of graphene in the growth front and on the surface of Si. ....... 91 Figure 4.22. SEM image and corresponding EDS elemental maps of B319-0.3 wt.% graphene composite. ..................................................................................................................................... 91 Figure 4.23. Effect of graphene content on the (a) hardness, (b) % IACS, measured by eddy current electrical conductivity meter, (c) electrical conductivity, measured by four-point probe method, and (c) thermal conductivity of B319 Al alloy. ............................................................................. 92 Figure 5.1. Evolution of CoF as a function of sliding time for the B319 Al alloy composites reinforced with varying amounts of graphene tested at 5 N: (a) 0 wt.%, (b) 0.1 wt.%, (c) 0.2 wt.%, (d) 0.3 wt.%, (e) 0.5 wt.%, and (f) 0.7 wt.%. “I” and “II” zones in the figures represent the running-in and steady-state regimes, respectively. ..................................................................................... 97 Figure 5.2. Evolution of CoF as a function of sliding time for the B319 Al alloy composites reinforced with varying amounts of graphene tested at 10 N: (a) 0 wt.%, (b) 0.1 wt.%, (c) 0.2 wt.%, (d) 0.3 wt.%, (e) 0.5 wt.%, and (f) 0.7 wt.%. “I” and “II” zones in the figures represent the running-in and steady-state regimes, respectively. ..................................................................................... 98 Figure 5.3. Average CoF (including running-in and steady-state regimes) of B319-graphene (0-0.7 wt.%) composites tested at 5N and 10N loads. ..................................................................... 101 Figure 5.4. Secondary electron SEM images of the wear track: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 5 N load. ............................................ 104 Figure 5.5. Secondary electron SEM images of the wear track: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 10 N load. .......................................... 104 xvii Figure 5.6. Surface profile maps of wear tracks captured by laser confocal microscope: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 5 N load. The color scale indicates wear track depth. ................................................................................................. 105 Figure 5.7. Surface profile maps of wear tracks captured by laser confocal microscope: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 10 N load. The color scale indicates wear track depth. ................................................................................................. 105 Figure 5.8. SEM images of wear scar on the counterbody (100Cr6 balls) after wear testing of (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene composites tested under 5 N load. ...................................................................................................................................... 107 Figure 5.9. (a) SEM micrograph of wear scar on steel ball slide against B319-0.3 wt.% graphene composite under 5 N, marked region (A) in Figure 5.8(d) and (b) the corresponding EDS area map...................................................................................................................................................... 108 Figure 5.10. SEM images of wear scar on the counterbody after wear testing of (a) B319-0 wt.% graphene, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (d1) high magnification image of a region (A) marked in Fig. (d), (d2) carbon map of Fig. (d1), (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 10 N load. ................. 108 Figure 5.11. Average wear scar area on the counterbody (100Cr6 balls) after wear testing of B319-graphene composites. .................................................................................................................. 109 Figure 5.12. Effect of graphene on the mass loss of B319 Al alloy-graphene composites tested under 5 N and 10 N loads. ........................................................................................................... 111 Figure 5.13. Effect of graphene on the wear volume loss of B319 Al alloy tested under (a) 5 N and (b) 10 N loads. The wear volume loss was calculated using the weight loss method and a laser microscope. ..................................................................................................................................112 Figure 5.14. Effect of graphene on the specific wear rate of B319 Al alloy tested under (a) 5 N and (b) 10 N loads. The specific wear rate was calculated using the wear volume loss determined by the weight loss method and 3D laser microscope. ..................................................................113 xviii Figure 5.15. Average wear volume and dissipated frictional energy of B319 Al alloy-graphene composites tested at (a) 5 N and (b) 10 N loads. .........................................................................114 Figure 5.16. Worn surface morphologies of B319 Al alloy tested at 5 N load: (a) SEM image, (b) laser microscope image, (c) SEM micrograph of the region marked in Fig. (a) and (b), (d) SEM micrograph of the region marked in Fig. (c) and corresponding EDS area map (e). ...................118 Figure 5.17. SEM micrograph of the worn surface of B319 Al alloy tested under a 5 N load, showing the different wear features. ............................................................................................119 Figure 5.18. Worn surface morphologies of B319-0.3 wt.% graphene composite tested at 5 N: (a) SEM image, (b) laser microscope image, (c) SEM micrograph of the region marked in Fig. (a) and (b), (d) SEM micrograph of the region marked in Fig. (c) and corresponding EDS area map (e)...................................................................................................................................................... 120 Figure 5.19. (a) SEM micrograph of B319-0.3 wt.% graphene composite tested under a 5 N load, showing the different wear features; (b) SEM image of graphene marked in Figure (a) and corresponding EDS area map of carbon (c). ............................................................................... 121 Figure 5.20. SEM images of worn surfaces of the B319-graphene composites tested under 10 N: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, (f) B319-0.7 wt.% graphene and (g) EDS area elemental mapping of B319 Al alloy (a). .................................................................................................... 123 Figure 5.21. (a) SEM image of the worn surface of the B319-0.1 wt.% graphene composite tested under a 10 N load and its corresponding EDS area maps (b). .................................................... 124 Figure 5.22. (a) SEM image of the worn surface of the B319-0.2 wt.% graphene composite under a 10 N load and its corresponding EDS area maps (b). .............................................................. 125 Figure 5.23. SEM image of worn surfaces of (a) B319 Al alloy at 100x magnification and (b) B319-0.3 wt.% graphene composite at 500x magnification. ...................................................... 126 Figure 5.24. SEM images of worn surfaces and their corresponding EDS area maps for B319-0.5 wt.% graphene composite. .................................................................................................... 127 Figure 5.25. SEM images of worn surfaces and their corresponding EDS area maps for B319-0.7 wt.% graphene composite. .................................................................................................... 127 xix Figure 5.26. (a) SEM image of a worn surface in B319-0.3 wt.% graphene composite, showing the presence of graphene; High magnification images of graphene marked in Figure (a): (b) region A, (c) carbon map of Figure (b), (d) region B, and (e) optical image of B319-0.3 wt.% graphene composite’s worn surface captured by Raman spectrometer, (f) Raman spectra of graphene (region D) marked in Figure (e), (g,h) SEM images of worn surfaces, showing tribo-morphological changes of graphene, (i) EDS elemental spectrum of a region (red square box, C) marked in Figure (d) and its quantification data. .................................................................................................... 128 Figure 5.27. SEM images of wear debris from B319-graphene composites with varying amounts of graphene tested under a 5 N load: (a) 0 wt.% graphene, (b) 0.1 wt.% graphene, (c) 0.2 wt.% graphene, (d) 0.3 wt.% graphene, (e) 0.5 wt.% graphene, and (f) 0.7 wt.% graphene. ............... 131 Figure 5.28. SEM images of wear debris from B319-graphene composites with varying amounts of graphene tested under a 10 N load: (a) 0 wt.% graphene, (b) 0.1 wt.% graphene, (c) 0.2 wt.% graphene, (d) 0.3 wt.% graphene, (e) 0.5 wt.% graphene, and (f) 0.7 wt.% graphene. ............... 132 Figure 5.29. Cross-sectional SEM-BSE micrographs of the worn surface of B319 Al alloy tested under 5 N load: (a) at low magnification, (b) at high magnification, and (c) EDS area elemental maps of Figure (b), showing three different zones. .................................................................... 135 Figure 5.30. Cross-sectional SEM-BSE micrographs of the worn surface of B319 Al alloy tested under a 10 N load: (a) at low magnification, (b) at high magnification, and (c) EDS area elemental maps of Figure (b), showing three different zones. .................................................................... 136 Figure 5.31. (a) Cross-sectional SEM-BSE micrograph of the worn surface of B319-0.3 wt.% graphene composite tested under a 5 N load at low magnification and corresponding EDS area maps (b), showing three different zones. .................................................................................... 137 Figure 5.32. (a) Cross-sectional SEM-BSE micrograph of the worn surface of B319-0.3 wt.% graphene composite tested under a 5 N load at high magnification and corresponding EDS area elemental maps (b), showing three different zones and stable MML. ........................................ 138 Figure 5.33. Cross-sectional SEM-BSE micrographs of the worn surface of B319-0.3 wt.% graphene composite tested under a 10 N load, (a) at low magnification, (b) at high magnification, and (c) EDS area elemental mapping of Figure (b), showing stable MML and its sublayers. ... 139 xx Figure 5.34. Cross-sectional SEM-BSE micrographs of the worn surface of B319 Al alloy tested under a 5 N load: (a) at low magnification, (b) at high magnification, and (c) EDS area elemental maps of Figure (b), showing the presence of cracks and delamination of MML. ...................... 142 Figure 5.35. (a) Cross-sectional EPMA-BSE micrograph of the worn surface of B319 Al alloy tested under a 10 N load, and (b) WDS area elemental mapping of the region (red square box) marked in Figure (a), showing the presence of cracks and delamination of MML. ................... 143 Figure 5.36. (a) EPMA-BSE cross-sectional micrograph of the worn surface of B319-0.3 wt.% graphene composite tested under a 10 N load, and (b) WDS area elemental mapping of the region (red square box) marked in Figure (a). ....................................................................................... 144 Figure 5.37. Cross-sectional SEM-BSE micrograph of the worn surface of B319-0.5 wt.% graphene composite tested under a 10 N load: (a) at low magnification, (b) high magnification image of the region marked in Figure (a) and corresponding EDS area elemental maps (c), showing the presence of cracks and graphene agglomeration. .................................................................. 146 Figure 5.38. Cross-sectional SEM-BSE micrograph of the worn surface of B319-0.7 wt.% graphene composite tested under a 10 N load: (a) at low magnification, (b) high magnification image of the region marked in Figure (a) and corresponding EDS area elemental maps (c), showing the presence of cracks and delamination of MML. ..................................................................... 147 Figure 5.39. Cross-sectional microhardness of subsurface layers as a function of distance from worn surface for B319 Al alloy and B319-0.3 wt.% graphene composite tested under loads: (a) 5 N and (b) 10 N. ........................................................................................................................... 148 Figure 5.40. Wear process in the B319 Al alloy: (a) initial microstructure with coarse as-cast grain structure, (b) formation of a thick MML and a deformed layer in B319 Al alloy, (c) subsurface crack initiation, (d) MML delamination, generating coarse wear debris; Wear process in the B319-0.3 wt.% graphene composite: (e) initial microstructure with refined as-cast grain structure, (f) formation of compact and graphene-containing MML or tribolayer, (g) crack initiation at the interface of sublayers within MML, and (h) MML fragmentation and generation of equiaxed wear debris; Wear process in the B319-0.5, 0.7 wt.% graphene composites: (i) initial microstructure, showing the presence of pores and graphene agglomeration, (j) formation of MML, (k) crack initiation at pores and regions of graphene agglomeration, (l) MML delamination, generating sheet-like coarse wear particles. ................................................................................................. 150 xxi List of Abbreviations AGNCs Aluminum-Graphene based Nanocomposites APT Atom Probe Tomography BSE Backscattered Electron BM Ball Milling CAFE Corporate Average Fuel Economy CNTs Carbon Nanotubes CoF Coefficient of Friction CTE Coefficient of Thermal Expansion DTA Differential Thermal Analysis EBSD Electron Backscatter Diffraction EDS Energy Dispersive Spectroscopy EPA Environmental Protection Agency EPMA Electron Probe Microanalyzer ERHAMI Electrical Resistance Heat-Assisted Mechanical Impregnation FSP Friction Stir Processing GHG Greenhouse Gas GN Graphene-based Nanomaterial GNFs Graphene Nanofibers GNPs Graphene Nanoplatelets GNSs Graphene Nanosheets GO Graphene Oxide HP Hot Pressing HPT High-Pressure Torsion IACS International Annealed Copper Standard MML Mechanically Mixed Layer MMM Mechanically Mixed Material MPG Miles Per Gallon OES Optical Emission Spectrometer PAS Pressure Assisted Sintering xxii PIC Pressure Infiltration Casting PLS Pressureless Sintering PM Powder Metallurgy rGO Reduced Graphene Oxide SAM Solution Assisted Mixing SDAS Secondary Dendrite Arm Spacing SEM Scanning Electron Microscopy SLM Selective Laser Melting SPS Spark Plasma Sintering STEM Scanning Transmission Electron Microscopy TEM Transmission Electron Microscopy TGA Thermogravimetric Analysis UTS Ultimate Tensile Strength WDS Wavelength Dispersive Spectroscopy XRD X-ray Diffraction YS Yield Strength xxiii List of Symbols 𝑎 Contact Radius 𝛼 Thermal Diffusivity 𝐶𝑝 Heat Capacity d Interlayer Spacing 𝐸∗ Reduced Modulus 𝐸1 Elastic Modulus of the B319 Al alloy 𝐸2 Elastic Modulus of the 100Cr6 Ball 𝐹𝑓 Friction Force 𝐹𝑛 Normal Load 𝐻 Hardness 𝑘 Specific Wear Rate 𝐾 Wear Coefficient 𝜅 Thermal Conductivity 𝜆 Wavelength 𝑚𝑎 Samples Weight in Air 𝑚𝑤 Sample Weight in Water µ Coefficient of Friction 𝑝𝑜 Maximum Hertzian Contact Pressure 𝑅 Ball Radius ⍴ Density ⍴𝐻2𝑂 Density of Water ⍴𝑠𝑎𝑚𝑝𝑙𝑒 Density of Sample ⍴𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 Theoretical Density 𝑆 Sliding Distance 𝑡 Sliding Time 𝜃 Bragg’s Angle 𝑉 Wear Volume ʋ𝑠 Sliding Speed 𝜗1 Poisson's Ratio of the B319 Al Alloy xxiv 𝜗2 Poisson's Ratio of the 100Cr6 Ball 𝛥𝑊 Weight Loss xxv List of Elements Al Aluminum C Carbon Cr Chromium Cu Copper Fe Iron Mg Magnesium Mn Manganese Ni Nickel O Oxygen Pt Platinum Si Silicon Sn Tin Sr Strontium Ti Titanium Zn Zinc xxvi List of Units cm Centimeter °C Degree Celsius g Gram GPa Giga Pascal HV Vicker’s Hardness J Joule K Kelvin Kg Kilogram kV Kilo Volts m Meter mg Milligram min Minute mm Millimeter ms Millisecond nA Nanoampere nm Nanometer µm Micrometer MPa Mega Pascal N Newton rpm Revolutions Per Minute s Second S Siemens W Watt xxvii Acknowledgments As the culmination of my doctoral studies, which commenced five years ago, approaches, I would like to express my profound gratitude to all who have supported me throughout this journey. First and foremost, I express my heartfelt gratitude to my supervisor, Prof. Lukas Bichler, for his ongoing support, invaluable guidance, and constant encouragement, which have motivated me to continually challenge myself during this journey. I have been fortunate to have a supervisor like him, whose mentorship not only provided me with invaluable guidance but also fostered an environment of academic freedom, allowing me to explore my research ideas confidently. Additionally, the opportunities he provided for collaboration on various other projects not only enhanced my skills and knowledge but also enriched my overall academic experience. Moreover, I am grateful for the valuable lesson he imparted on staying grounded and humble, a quality that has greatly influenced my approach to both my professional and personal life. Thank you, Professor, for making my graduate studies an enjoyable and memorable experience that I will cherish for life. I sincerely thank Mr. Kyle Lessoway for his help with the metal casting experiments. Working on this project with him has been thoroughly enjoyable. I cannot forget the many trials of Al castings we did to distribute graphene uniformly. Thank you, Kyle, for the wonderful times and for always being my go-to person whenever I needed help. I am grateful for the financial support of my research provided by the Natural Science and Engineering Research Council of Canada (NSERC). I am also indebted to our industry collaborators for their technical and financial backing. I extend my thanks to Dr. Anthony Lombardi and Dr. Glenn Byczynski from Nemak USA/Canada for their insightful discussions and knowledge sharing on automotive Al metal castings. I also express my gratitude to Dr. Colin van der Kuur from Zentek Ltd. for providing the graphene material. I would like to thank my doctoral committee members, Dr. Mohammad Arjmand and Dr. Jian Liu, for their support and helpful insights. I am deeply thankful to Dr. Sudip Shrestha from The Fipke Laboratory for Trace Element Research (FiLTER) for his expert assistance in SEM and EPMA. He generously allotted SEM slots whenever I needed them, even accommodating last-minute requests. I also extend my gratitude to Dr. Mark xxviii Button for his help with ICP-OES data acquisition and friendly support. Furthermore, I appreciate the enjoyable conversations I’ve had with both of you during the SEM sessions. I am grateful for the help and cooperation from the technicians at the UBC Okanagan, School of Engineering. A big thanks to Dr. Matthew Brown for his help with XRD and Raman Spectroscopy. Additionally, I appreciate all the training and support provided by Mr. Raymond Seida, Mr. Durvin Bossy, Mr. Praveen Rajan, and Mr. Alec Smith. They generously shared their technical knowledge and provided assistance with machining my samples and graphite molds. Most importantly, I thank all the past and present members of our research group at UBC Okanagan, especially Dr. Anil Prasad, from whom I learned SPS, and Dr. Somi Doja, from whom I learned DSC. I truly appreciate both of your support during the initial days of my PhD. Additionally, I extend my gratitude to Dr. Kris Mackowiak for his support. I would also like to thank Mr. Mitchell Barrett and Mr. Abhinav Karanam for all their help whenever I needed it. Special thanks to Mr. Adil Shaik, Dr. Satish Idury, and Dr. Joshua Stroh for their valuable discussions and suggestions. I am deeply grateful to Prof. B.S. Murty from the Indian Institute of Technology Hyderabad for his insightful suggestions and constant support. Additionally, I extend my thanks to Dr. S.V.S. Narayana Murty from the Liquid Propulsion Systems Centre, Thiruvananthapuram, India, for fruitful discussions and timely motivation. Their words of encouragement were a beacon of hope during challenging times. I am grateful to both for their meticulous correction of the review article. Their constant source of inspiration has been invaluable to me. I extend my heartfelt thanks to my mom and sisters for their constant support and patience throughout my academic journey. I am forever indebted to my mom for her sacrifices. Thank you, Maa. Finally, I thank my wife, Lavanya, for her endless love, selfless support, and care. Without the support and encouragement of these amazing people, this degree would not have been possible. I always miss my father (Janardhana Reddy) and my best friend (Jyothiswar Reddy), who unfortunately are no longer with us. I deeply miss the support and encouragement you both offered, and this dissertation is a tribute to their profound impact on my life. 1 Chapter 1: Introduction 1.1 Background The ever-increasing demand for improved fuel efficiency and reduced CO2 emissions in the automotive sector continues to drive the increased use of aluminum (Al) and its alloys. The trend of increasing production of electric and hybrid cars has brought lightweighting into new focus due to the heavy weight of battery packs and the associated increase in vehicle curb weight. As a result, the development of new materials for powertrain components is critical since these components account for ∼24–28% of the total vehicle weight [1]. Al alloy B319 (Al-Si-Cu) has been generally used to make powertrain components, such as cylinder heads, engine blocks, and transmission cases, due to its good thermal and mechanical properties at ambient and elevated temperatures (up to ∼200ºC) [2–5]. However, a significant drawback of the B319 Al alloy lies in its low seizure load and insufficient wear resistance, restricting its widespread application in powertrain components [6]. Various technologies, primarily surface coatings [7], cylinder liners [8], and surface modification [9], have been employed to enhance the wear resistance of Al-Si alloy-based components. The use of iron-based cylinder liners or coatings introduces a challenge related to the generation of residual tensile stresses at the cylinder bore region. This is attributed to the large mismatch in the coefficient of thermal expansion (CTE) between the Al alloy (21.5-23.5 ×10-6/ºC [10]) and the coating or liner (gray cast iron: 10.5-13 ×10-6/ºC [11]), eventually resulting in spalling [8,12,13]. The surface modification by electrochemical etching [9] or laser surface texturing [14] has also been used commonly to enhance the friction and wear resistance of Al-Si alloys. In electrochemical etching, the surface of the Al-Si alloy is selectively etched to remove the soft Al phase, ensuring that the hard Si phase comes into contact with the counterbody (e.g., piston ring), thereby reducing wear on the Al-Si alloy. Similarly, in laser surface texturing, an array of microdimples on the alloy surface is created through a material ablation process using a pulsating laser beam. The microdimples act as lubricant reservoirs, ensuring more consistent and effective lubrication of the contact surfaces, thereby improving tribological properties [15]. However, potential drawbacks include the limited depth of modification, inconsistency in surface finish, and higher costs, thereby limiting the application of surface modification techniques for mass production of cylinder blocks. 2 Therefore, improving the tribological properties of the B319 Al alloy is preferred so that the necessity for coatings, liners, and surface modification of the cylinder bore is not diminished. Over the past three decades, researchers have tried various microstructural engineering approaches, including grain refinement [16], heat treatment [17,18], Si modification [16,19], and the addition of alloying elements [20,21] or ceramic particles [22–25] to enhance the tribological properties of Al-Si alloys. Among these approaches, adding ceramic particles, including SiC, B4C, Al2O3, and TiC, was the most effective method to improve the wear resistance of as-cast Al-Si alloys. Although this approach has achieved some success, some authors have reported a critical load above which the composites showed an increased wear rate over the base alloy [26–28]. The mechanism of wear rate transition at higher loads remains not fully understood, but it was suggested that the fracture of ceramic particles at the contact surface was accompanied by material delamination due to subsurface cracks. Furthermore, the abrasive nature of fractured ceramic particles can cause wear of both the body and the counterbody, thus deteriorating the overall wear resistance of the tribo-system [28]. Finally, to achieve a meaningful improvement in wear resistance, a high content of ceramic particles was required, which contributed to increasing the component's weight and cost. Recently, graphene has emerged as an alternative to ceramic reinforcement due to its desirable combination of mechanical, thermal, electrical, and solid lubrication properties [29–32]. Small concentrations of graphene were seen to improve the friction and wear resistance of Al alloys due to graphene’s high surface area (2630 m2/g [33]), high in-plane strength (130 GPa [34]), and weak inter-layer bonding. As a result, numerous studies have been carried out to improve Al alloy's mechanical, thermal, and tribological properties [31]. Zhang et al. [35] fabricated Al-Si alloy A355 reinforced with 1 vol.% graphene composite by ball milling (BM) followed by hot pressing (HP), and the composite showed a reduction of ∼39% and ∼85% wear rate and coefficient of friction (CoF) over the base alloy, respectively. Similarly, Singh et al. [36] reported incorporating 0.7 wt.% reduced graphene oxide (rGO) into ADC12 alloy decreased the wear volume by ∼50% and CoF by ∼90%. In another study, Tiwari et al. [37] reported a 25% increase in the thermal conductivity of Al by adding 0.1 wt.% graphene. Although there are several studies on the wear behavior of Al-graphene based nanocomposites (AGNCs) fabricated by powder metallurgy (PM) [31,38–41], limited research is available on the 3 tribological performance of cast AGNCs [36,42–44]. Specifically, there is a paucity of systematic studies investigating graphene's effect on the microstructure and tribological properties of cast B319 Al alloy. Additionally, existing research on the wear behavior of AGNCs has primarily focused on the effect of graphene addition on wear properties, primarily identifying wear mechanisms based on the topography of worn surfaces. However, there has been a notable absence of subsurface analysis of worn surfaces, which is essential for a comprehensive understanding of the impact of graphene on wear mechanisms in these composites. The scarcity of literature on cast AGNCs can be attributed to the challenges associated with uniformly dispersing graphene in the Al melt during the casting process. These challenges arise from several factors, including the agglomeration of graphene and its segregation at the melt’s surface driven by strong van der Waals forces and density mismatch, respectively [31]. Consequently, significant attention has been directed towards distribution approaches in the past decade. Among various dispersion methods, the flake-PM approach in the PM route [45–48] and the utilization of composite master alloy in the casting route [49] have shown promising results in dispersing graphene. For instance, Mirjavadi et al. [49] dispersed graphene in an Al-10Zn-3.5Mg-2.5Cu alloy through a composite master alloy approach and the addition of 0.7 wt.% graphene improved the ultimate tensile strength (UTS) of Al-10Zn-3.5Mg-2.5Cu alloy by 35%. Considering these gaps in the existing literature, this study focuses on the fabrication of B319 Al alloy composites reinforced with varying amounts of graphene (0-0.7 wt.%) using a novel distribution approach. Initially, a composite master alloy of F357 Al alloy-5 wt.% graphene was developed through a flake-PM process involving BM and spark plasma sintering (SPS). Subsequently, the composite master alloy was introduced into the B319 Al alloy melt. Upon melting of the master alloy within the liquid alloy, graphene was released, thereby overcoming common challenges associated with graphene dispersion. The obtained composites were subjected to comprehensive characterization to evaluate the influence of graphene on the microstructure, hardness, electrical conductivity, and thermal conductivity of the as-cast B319 Al alloy. Furthermore, the friction and wear behavior of the B319-graphene composites were studied under reciprocating dry sliding conditions. The worn surface and its subsurface microstructure were characterized to understand the role of graphene in enhancing the wear resistance of the B319 Al alloy and to identify the wear mechanisms in these novel composites. 4 1.2 Objectives of the Work With the above background, the objectives of the present work were: 1. To fabricate castable B319-graphene composites using a flake-PM approach involving BM and SPS. 2. To investigate the effect of graphene on the microstructure, hardness, electrical conductivity, and thermal conductivity of the B319-graphene composites. 3. To study the friction and wear behavior of the B319-graphene composites under reciprocating dry sliding conditions. 4. To elucidate the role of graphene in enhancing the tribological properties of the B319 Al alloy. 5. To identify the dominant wear mechanisms in the B319-graphene composites through detailed subsurface microstructural analysis of wear tracks. 1.3 Scope of the Dissertation This dissertation focuses on fabricating B319-graphene (0-0.7 wt.%) composites using a flake-PM approach involving BM and SPS followed by casting. It investigates the influence of graphene on the microstructure, hardness, electrical conductivity, and thermal conductivity of the as-cast composite materials. Furthermore, it discusses the reciprocating dry sliding friction and wear behavior of B319-graphene composites, complemented by detailed microstructural analysis. Figure 1.1 presents a flow chart outlining the research conducted in this thesis, providing an overview of the methodology and key aspects explored. The dissertation has been organized as follows: Chapter 1 provides an overview of the background and objectives of the current work, while also outlining the scope and structure of this thesis. Chapter 2 offers a comprehensive literature review, disseminating state-of-the-art knowledge on the processing challenges, fabrication techniques, and properties of AGNCs. Chapter 3 presents the materials and experimental details of processing, and characterization techniques employed in this study. 5 Chapter 4 deals with the results and discussion of the synthesis and microstructure of composite master alloys, as well as the effect of graphene on the microstructure, hardness, electrical conductivity, and thermal conductivity of the B319 Al alloy. Chapter 5 focuses on the results and discussion of the frictional and wear behavior of B319-graphene composites. Furthermore, it examines the wear mechanisms in these composites through detailed microstructural characterization of wear tracks and their subsurface microstructure. Chapter 6 summarizes the major findings of the dissertation and outlines opportunities for future research. Figure 1.1. Flow chart illustrating the research methodology and key aspects explored in this thesis. 6 1.4 Novelty and Originality of the Research The novelty and originality of the research presented in this dissertation lies in the following four key aspects: 1. Composite system: To the best of the author’s knowledge, this is the first attempt to explore the development of cast B319 Al alloy-graphene composites. 2. Distribution approach: The B319-graphene composites were fabricated using a novel distribution approach. This approach involved fabricating a composite master alloy (F357-5 wt.% graphene) through a flake-PM approach involving BM and SPS. Subsequently, the composite master alloy was introduced into the B319 Al alloy melt to create the B319-graphene composites via stir casting. 3. Property evaluation: For the first time, the effect of graphene addition on the friction and wear behavior of the B319 Al alloy was investigated. 4. Wear mechanisms: An in-depth analysis of the subsurface microstructure of wear tracks was carried out to understand the role of graphene in enhancing the tribological performance of the B319 Al alloy and to identify the dominant wear mechanisms in the novel B319-graphene composites. 7 Chapter 2: Literature Review This chapter provides a state-of-the-art literature review on Al matrix composites reinforced with graphene-based nanomaterials (GNs), focusing on their fabrication and properties. The chapter begins by highlighting the need for “lightweighting” and AGNCs in Section 2.1, followed by a brief introduction to the structure and properties of GNs. Section 2.2 addresses the current challenges in the processing of AGNCs, including issues such as low wettability, interfacial reactions, and agglomeration of GNs. Different processing techniques employed for the fabrication of AGNCs are discussed in Section 2.3. Section 2.4 critically analyzes various properties of AGNCs, encompassing mechanical, electrical, thermal, and tribological properties. Additionally, this section elucidates the active strengthening mechanisms observed in AGNCs, such as load transfer, Orowan strengthening, dislocation strengthening, and grain refinement. The current understanding of the AGNCs is summarized in Section 2.5. Finally, Section 2.6 outlines the problem identification and methodological approach adopted in this study. 2.1 Introduction 2.1.1 The Need for Lightweighting The enactment of new fuel economy standards by governments internationally, coupled with the growing consumer demand for more fuel-efficient vehicles, has increased pressure on the automotive and aerospace sectors to improve fuel efficiency and mitigate CO2 emissions. For instance, the contribution of various sectors to greenhouse gas (GHG) emissions in Canada is depicted in Figure 2.1. Notably, transportation ranks among the largest sectors, contributing 25% of total GHG emissions in Canada. Similarly, in the United States (U.S.), the transportation sector contributed around 28% of total U.S. GHG emissions in 2018, according to data from the Environmental Protection Agency (EPA) published in April 2020. In the U.S., CAFE (Corporate Average Fuel Economy) standards have been established to improve vehicle's fuel economy (passenger cars and light trucks) and control CO2 emissions. The CAFE standard has increased from 18 MPG (miles per gallon) in 1978 to the current 49.2 MPG for passenger cars, with a projected increase to 59.4 MPG by 2026. Every country has its own fuel economy standards, many of which are more stringent than those of the U.S. However, all countries share the dual goal of reducing CO2 emissions and improving fuel efficiency. 8 Figure 2.1. Contribution of various sectors to greenhouse gas (GHG) emissions in Canada. (Source: National Inventory Report 1990–2018: Greenhouse gas sources and sinks in Canada). Manufacturers have been exploring different approaches to meet fuel economy standards and governmental regulations, with one strategy involving the reduction of vehicle weight. This reduction can be achieved by either downsizing the vehicle or using lightweight materials. However, downsizing may potentially compromise passenger safety and comfort. Therefore, a focus on substituting heavier materials with lighter ones while maintaining vehicle performance is often pursued. Research indicates that a 10% decrease in vehicle weight can lead to an 8 to 10% improvement in fuel economy, and a reduction of 100 kg in vehicle weight may yield a 12.5 g/km decrease in CO2 emissions [50]. Moreover, reducing vehicle weight can enhance performance without compromising occupant safety. Table 2.1 presents the trend in the average weight of U.S. vehicles. It clearly shows that the reduction in vehicle weight began in the late 1970s to meet the CAFE-1978 standards. Automakers 9 achieved this reduction by employing lighter materials or downsizing vehicles. Consequently, the average weight of U.S. vehicles decreased from 1841 kg in 1975 to 1483 kg in 1985. However, the introduction of numerous new features in the early 1990s resulted in increased vehicle weight. These features included safety enhancements (such as airbags and anti-block systems), in-car entertainment systems, and emission control devices. As of 2022, the average vehicle weight of U.S. vehicles stands at 1952 kg, and the trends are suggesting a weight of 2013 kg for 2023. It is expected that this trend will continue in the coming years as manufacturers introduce new models with enhanced performance and increased luxury. In view of the increasing demand for CO2 control, fuel efficiency, and lightweighting, there is an immediate need to develop high-performance, lightweight materials for next-generation automotive applications. Table 2.1. Trend in average weight of vehicles in the U.S. (Source: U.S. Environmental Protection Agency, 2023 EPA Automotive Trends Report). Year Average vehicle weight lbs kgs 1975 4060 1841 1985 3271 1483 1990 3426 1554 2000 3821 1733 2010 4001 1814 2012 3979 1804 2014 4060 1841 2016 4035 1830 2017 4093 1856 2018 4137 1876 2019 4155 1884 2020 4166 1890 2021 4289 1945 2022 4303 1952 10 2.1.2 The Need for Al-Graphene based Nanocomposites (AGNCs) Before the 1970s, low-carbon steel and cast iron were the most used automotive industry materials. The cylinder head and engine block are among the two critical components of a passenger vehicle. Both components have historically been made of cast iron due to its good high-temperature strength, excellent wear resistance, and castability. However, cast iron is very dense (7.87 g/cm3), making the engine block one of the heaviest components, accounting for 3-4% of the vehicle's total weight [3,51]. As a result, these components present unique opportunities for immediate weight reduction with the inclusion of alternative lightweight materials. Figure 2.2. Trend in weight percentage of steels, cast iron, and aluminum in light-duty vehicles [7]. Data extracted from (Q. Dai, J.C. Kelly, A. Elgowainy, Vehicle Materials: Material composition of U.S. light-duty vehicles, 2016). Al alloys were considered an immediate substitute for cast irons and steels due to their lightweight (2.7 g/cm3), which is about 65% lower than that of cast iron. The replacement of cast iron with Al alloys in engine blocks offered good potential for weight reduction of up to 45% [52]. Manufacturing of engine blocks with Al alloys started in the late 1970s, and the cast iron cylinder 11 heads have been entirely substituted by Al alloys during the past decade [53]. Significant research is currently underway to replace steel in other automotive components with Al alloys in next-generation electric vehicles. The trend in Al weight percentage along with steel and cast iron in U.S. light-duty vehicles is shown in Figure 2.2. Since 1975, the weight percentage of Al usage in light-duty vehicles has increased almost five-fold (from 2% in 1975 to 9.77% in 2014) [54], and it is expected to reach 16% by 2028, according to a study conducted by Ducker Worldwide. Currently, cast Al alloys, mainly A356 (Al-Si-Mg) and B319 (Al-Si-Cu) are predominantly used in the cylinder head and engine block applications. The principal alloying element in these alloys is silicon (Si), which contributes to the alloy’s excellent hot tear resistance, high castability, and fluidity. These alloys are age hardenable, with the presence of magnesium (Mg), and Copper (Cu) leading to the formation of Mg2Si and Al2Cu precipitates, with a concomitant enhancement of fatigue resistance [18,55,56]. However, the Cu-bearing B319 Al alloy exhibits better creep and tensile properties at elevated temperatures compared to the A356 Al alloy, as Al2Cu precipitates are more stable than Mg2Si precipitates [57]. Moreover, the exceptional thermal conductivity of the B319 Al alloy enables efficient heat dissipation from the combustion chamber compared to cast iron, allowing engines to operate at higher temperatures and deliver increased power per cylinder size [58]. Over the past decade, the operating temperature of cylinder heads has risen from 170℃ to a maximum of 200℃ with the substitution of Al alloys [57]. However, the operating temperature of cylinder heads must increase further to meet the latest CAFE standards. Therefore, improving the thermal conductivity of the B319 Al alloys is possibly one of the best ways to increase a combustion engine's operating temperature, such that the combustion heat is quickly dissipated into the coolant. Considering the piston’s reciprocating motion within the cylinder bore and the continuous contact of the piston ring with the cylinder bore wall, the essential characteristics sought in engine block materials include wear resistance and a low CoF between the cylinder bore and piston ring. Friction among engine components accounts for a substantial portion of total vehicle power consumption, with over 50% of this power loss attributed to frictional losses between cylinder bores and the piston assembly [3]. Studies have indicated that a 10% reduction in friction loss can yield a 3% improvement in fuel economy [59]. Consequently, minimizing frictional losses between the 12 cylinder bore surface and the piston ring emerges as a critical parameter in the development of alternative materials for cylinder blocks [60,61]. Al-Si alloys are softer than cast iron, and as a result, they exhibit poorer friction and wear properties. Consequently, the Al-Si alloy cylinder bore walls are typically equipped with cast-in or pressed-in iron liners. Figure 2.3 shows a V8 engine with cast-in iron liners on the cylinder bore surface. However, the mismatch in CTE between the Al alloy and liner, reduced thermal conductivity, increased cost, and engine weight due to the inclusion of iron liners are significant drawbacks. These drawbacks prompt the exploration of alternative approaches, such as employing surface coatings with electroplating or thermal spraying [7,60,62,63], and implementing surface modification through electrochemical etching or texturing [9,14,15] to enhance the wear properties of cylinder bore surfaces. Despite the advantages provided by these approaches, potential limitations include the delamination of coatings over time [64], the possibility of galvanic corrosive attack [65], and the costly and complex nature of the process [3]. Javidani and Larouche [3] have provided a comprehensive review discussing alternative approaches for manufacturing liner-less Al-Si alloy cylinder blocks. Figure 2.3. Photograph of V-8 engine block with cast-in iron liners [8]. Reprinted from J. Mater. Process. Technol., 189, E. Carrera, A. Rodríguez, J. Talamantes, S. Valtierra, R. Colás, 206-210, Copyright (2007), with permission from Elsevier. 13 In light of the stringent weight reduction requirements in the automotive sector mentioned above, there persists a pressing demand for immediate high-performance Al-Si-based materials that offer multi-functional properties such as improved mechanical properties, good wear resistance, and high thermal and electrical conductivity, all without increasing the component’s weight. However, most conventional Al-Si alloys fail to simultaneously meet these requirements, prompting researchers to explore alternative approaches to enhance the tribological properties of the Al-Si alloy, thereby eliminating the need for coatings, liners, and surface modification of the cylinder bore. One effective alternative approach is modifying the Al-Si alloy cylinder block with suitable reinforcements. The resulting composites exhibit a combined property of the matrix (Al and its alloys) and reinforcement (ceramic, carbon-based materials). Previous studies have shown that adding SiC [66], Al2O3 [67], and B4C [23] to Al-Si alloy enhances the wear resistance, strength, and hardness of the Al alloy. However, composites reinforced with micron-sized particles are prone to failure through easy crack initiation at the weak matrix-reinforcement particle interface and require a high fraction of reinforcement to achieve the desired level of properties [68]. For instance, Sajjadi et al. [69] developed 356 Al alloy composites reinforced with micron (20 µm) and nano (50 nm) sized Al2O3 and observed that a 1 wt.% Al2O3 nanocomposite exhibited better compressive strength than a 10 wt.% Al2O3 microcomposite. Al reinforced with nanosized reinforcement such as carbon nanotubes (CNTs), and graphene is a possible system showing a promise to address the limitations observed with microcomposites. Recently, graphene's leading derivatives, such as graphene oxide (GO) and rGO, are being intensely investigated as a potential reinforcement for Al and its alloys. In comparison to CNTs, GN is a better reinforcement for Al due to its higher surface area per unit volume [70]. Moreover, GNs have a lower tendency to twist than CNTs, which makes GNs disperse relatively easily in the metal matrix. Reducing the reinforcement size to nano-level (<100 nm) effectively impedes dislocation motion and improves mechanical properties. Even at lower reinforcement content, nano-sized reinforcements contribute to significant improvements in properties, primarily attributable to their higher specific surface area [71]. However, a sufficient volume fraction of reinforcement and ensuring its homogeneous distribution in the matrix is crucial to obtaining improved properties, even with the nano reinforcement. Shin and Bae [72] demonstrated that the addition of only 0.7 vol.% of graphene increased the yield 14 strength (YS) of AA2024 alloy by 50%. Zhang et al. [35] improved the tribological properties of Al by reducing the CoF by 39% and wear rate by 85%, respectively, with a 1 vol.% addition of graphene. Wang et al. [46] prepared graphene-reinforced Al composites and observed a 62% improvement in tensile strength over unreinforced Al with a 0.3 wt.% graphene addition. The literature consistently highlights that incorporating GN in Al alloys can improve the mechanical and tribological properties of Al without adversely affecting weight. Therefore, there is an immediate need to pursue further research on AGNCs for applications where strength, wear resistance, and lightweight characteristics are key requirements. 2.1.3 Synthesis, Structure, and Properties of GNs (Pristine Graphene, GO, and rGO) Graphene is a single layer of carbon atoms arranged in a two-dimensional (2D) crystalline hexagonal lattice [73]. Graphene (as reinforcement in Al alloys) has attracted a lot of attention in the last few years due to its high Young’s modulus, fracture strength, large surface area, low density, and excellent thermal and electrical conductivity. Various properties of GN are presented in Table 2.2. The theoretical Young’s modulus of pristine graphene is as high as 1 TPa, significantly higher than Young’s modulus of current Al reinforcing agents: SiC (450 GPa) [74], Al2O3 (386 GPa) [75], TiC (310-458 GPa) and TiN (440-616 GPa) [76]. The properties of GNs largely depend on their structure, shape, and size, which are primarily influenced by the synthesis technique. There are significant structural differences between pristine graphene, GO, and rGO that determine their physical, mechanical, electrical, and thermal properties. The schematic structure of GNs is illustrated in Figure 2.4. The ideal structure of pristine graphene consisting of single-atom-thin fully sp2 bonded carbon atoms arranged in a hexagonal lattice structure with no defects is shown in Figure 2.4(a). The carbon atoms are tightly bonded covalently with a bond length of 0.142 nm [77]. Pristine graphene can be synthesized by various methods, including thermal exfoliation [78], chemical exfoliation [79], mechanical exfoliation [80], chemical vapor deposition [81], epitaxial growth of graphene on various substrates [82,83], and production of graphene from graphite oxide [84,85]. GO is a highly oxidized form of graphene, that contains various oxygen functional groups in the carbon lattice, as shown in Figure 2.4(b). GO can be produced using various methods, such as the Brodie method [86], the Staudenmaier method [87], and the Hummers method [88]. 15 Table 2.2. Properties of graphene-based nanomaterials (Pristine graphene, GO, and rGO). Property* Pristine Graphene Graphene Oxide (GO) Reduced Graphene Oxide (rGO) C:O ratio No Oxygen [89] 2-4 [89] 8-246 [89] Density (g/cm3) 1.8 to 2.2 [68] 1.8 [90] 2.05 [91], 2.26 [92] Electrical conductivity (S/m) 100 × 106 [93] 0.5 [92], 3×10-4 [94] 200-35100 [95], 1500 [94], 1192 [96] Electron mobility (cm2/Vs) 200,000 [97] Insulator or semiconductor, depending on the degree of oxidation 372 [98], 1596 [99] Fracture strength 125 GPa [34] 120 MPa [90] N/A Optical transmittance (%) 96 [100], 97.7 [101] 90 [102] 80 [102,103] Sheet resistance (Ω/sq) 125 [104] 4×1010 [84] 4×106 [84] Specific surface area (m2/g) 2600 [105], 2630 [33] 368-382 [106], 463 [107], 925 [108] 467 [91], 705 [33], 320 [109] Thermal conductivity (W/mK) 4840±0.44- 5300±0.48 [110] 776 [111] 2180-2275 [111] Thermal stability 450-650 (in air) [68] Mass loss occurs above 60-80℃ [112] No significant mass loss up to 800℃ (in N2) [113] Thermal expansion coefficient (10-6K-1) -1 [68,114], -1.3 [115] -67 [116], -11.7 to -130.14 [117] 14.9-33.6 [118] Young’s modulus (TPa) 1 [34] 0.042±0.002 [90], 0.2076±0.0234 [119] 0.25±0.15 [120] N/A: Not available *The method of synthesis and the content of oxide functional groups, defects, and impurities influence the properties of graphene and its derivatives. 16 Figure 2.4. A schematic showing the crystal structure of graphene-based materials (a) pristine graphene, (b) graphene oxide (GO), and (c) reduced graphene oxide (rGO). The inset shows the C-C bond length (0.142 nm) and type (covalent) in pristine graphene. Figure 2.5. High-resolution TEM image of vacancy defect in single-layer graphene, (a) pristine graphene, (b) graphene with vacancy defect, and (c) with atomic configuration [121]. Adapted with permission from J.C. Meyer, C. Kisielowski, R. Erni, M.D. Rossell, M.F. Crommie, A. Zettl, Direct imaging of lattice atoms and topological defects in graphene membranes, Nano Lett. 8 (2008) 3582–3586. Copyright (2008), American Chemical Society. The principle in all GO synthesis methods involves treating graphite with suitable oxidizing agents such as nitric acid, potassium permanganate, potassium chlorate, and sulfuric acid. Subsequent exfoliation of graphite oxide results in the formation of GO. The oxidation of graphite increases the interlayer spacing from 0.335 nm to 0.625 nm, and the hydrophilic nature of graphite oxide further increases the spacing to 1.15 nm due to water adsorption [122,123]. Therefore, the energy required to exfoliate the graphene layers is low when the precursor material is graphite oxide rather than graphite. Additionally, water can exfoliate graphene layers from GO due to their hydrophilic 17 nature, making this route the most cost-effective. However, the presence of oxygen functional groups deteriorates the mechanical and conductive properties of graphene. These oxygen functional groups in GO can be eliminated by a reduction process (chemical or thermal) to generate a pristine graphene-like structure with enhanced properties, known as rGO (Figure 2.4(c)). However, point defects are created during the reduction process, and it is impossible to remove all the functional groups. The vacancy defect in single-layer graphene captured with high-resolution TEM (transmission electron microscopy) is shown in Figure 2.5(a-c) [121]. The presence of oxygen functional groups in GO and defects and functional groups in rGO is the main reason for their lower set of properties than pristine graphene. The control of oxygen and functional groups in the carbon lattice can be achieved by tuning the oxidation conditions of graphite or the reduction conditions of GO. 2.2 Challenges in Developing AGNCs 2.2.1 Agglomeration of GN One of the major challenges in the development of AGNCs is achieving a uniform distribution of GN in the Al matrix. The strong van der Waals forces between the graphene layers, high surface area (2630 m2/g), and the density difference between Al (2.7 g/cm3) and GN (1.8–2.2 g/cm3) leads to an agglomeration of GN in the Al matrix. Although the van der Waals forces are weak intermolecular forces, they become significant at the nanoscale due to the large surface area of the GN. The van der Waals attractions between the GN layers lead to the self-assembly of GN layers, resulting in large cluster formations in the liquid Al matrix. The agglomeration becomes more severe, especially with the large fraction of GN. Li et al. [124] dispersed various wt.% (0.5, 1, 1.5, and 2) GNF (graphene nanofibers) in the Al matrix by mechanical mixing and cryomilling. This mixing approach was able to disperse 0.5 wt.% GNF uniformly in the Al matrix, as presented in Figure 2.6(a). However, this mixing approach was not effective for reinforcement fractions beyond 0.5 wt.% GNF, when clustering of GNF along the grain boundaries in the Al matrix was observed, as seen in Figure 2.6(b). These regions were seen to readily nucleate cracks, thereby degrading the composite's mechanical properties. In another study, Bhadauria et al. [125] used ultrasonication to disperse graphene nanoplatelets (GNPs) of 0.25, 0.5, and 1 wt.% in the Al matrix. They also observed an agglomeration in Al + 1 18 wt.% GNP composite (Figure 2.6(c)), which caused the tensile strength to decrease drastically from 171 (0.5 wt.% GNP) to 71 MPa (1 wt.% GNP). Also, in a study performed by Li et al. [47], Al with 0.5, 1.5, and 3 vol.% rGO was synthesized through ultrasonication, HP, followed by hot extrusion. Increasing the rGO content from 0.5 vol% deteriorated the tensile ductility from 14.8% to 5.9 (1.5 vol.%) and 2.1% (3 vol.%). This was attributed to the agglomerations of rGO in the 1.5 vol.% composite, as seen in the fracture surface of the Al+1.5 vol.% composite (Figure 2.6(d)). Thus, agglomeration is a serious problem and must be avoided to take advantage of graphene's superior properties in the Al matrix. Figure 2.6. TEM image of (a) Al+0.5 wt.% GNF, showing a uniform distribution of GNF [124]; (b) Al+1.0 wt.% GNF, showing GNF agglomeration in Al matrix [124]; (c) Al+1 wt.% GNP, showing GNP agglomeration in Al matrix [125]; and (d) scanning electron microscope (SEM) image of fracture surface of Al+1.5 wt.% rGO, showing agglomeration of rGO in Al matrix [47]. (a, b) Adapted from Mater. Sci. Eng. A, 626, J.L. Li, Y.C. Xiong, X.D. Wang, S.J. Yan, C. Yang, W.W. He, J.Z. Chen, S.Q. Wang, X.Y. Zhang, S.L. Dai, Microstructure and tensile properties of bulk nanostructured aluminum/graphene composites prepared via cryomilling, 400-405, Copyright (2015), with permission from Elsevier. (c) Adapted from J. Alloys Compd., 748, A. Bhadauria, L.K. Singh, T. Laha, Effect of physio-chemically functionalized graphene nanoplatelet reinforcement on tensile properties of aluminum nanocomposite synthesized cia spark plasma 19 sintering, 783-793, Copyright (2018), with permission from Elsevier. (d) Adapted from Carbon, 95, Z. Li, G. Fan, Q. Guo, Z. Li, Y. Su, D. Zhang, Synergistic strengthening effect of graphene-carbon nanotube hybrid structure in aluminum matrix composites, 419-427, Copyright (2015), with permission from Elsevier. 2.2.2 Poor Wettability and Al4C3 Formation The key factor in deciding the mechanical behavior of AGNCs is the ability to transfer the applied load between the Al and GN. In AGNCs, under a given load, a certain fraction of the load is carried by the GN and the rest by the Al. The main function of the Al matrix is to facilitate the transfer and distribution of the applied load to the individual GN sheets. The extent of load transfer from Al to GNs depends on their wettability and interfacial reactions [126,127]. Wettability describes the ability of a liquid (Al) to spread over a solid surface (GN) and can be measured by the wetting angle (Φ, contact angle). A wetting angle of Φ < 90° indicates that the liquid wets the solid surface, whereas Φ > 90° signifies that the liquid does not wet the solid surface. Research indicates that liquid Al exhibits a high wetting angle on graphene, resulting in poor wettability. Oh et al. [128] examined the wetting angle between molten Al and graphite, obtaining a measured angle of 140°, as shown in Figure 2.7(a). However, coating graphite with copper (Cu) (Figure 2.7(b)) or nickel (Ni) (Figure 2.7(c)) reduced the wetting angle [129]. Consequently, coating GN with a suitable material enhances the wettability of Al with GN. It is important to note that the presence of oxygen functional groups in GO and rGO can have a significant effect on wettability. Figure 2.7. The wetting behavior of Al on (a) graphite substrate, (b) Cu-coated graphite substrate [128], and (c) Ni-coated graphite substrate [129] captured during the sessile drop test. 20 (a,b) Adapted from J. Alloys Compd., 542, Se-Il Oh, J-Y. Lim, Y-C. Kim, J. Yoon, G-H. Kim, J. Lee, Y-M. Sung, J-H. Han, Fabrication of carbon nanofiber reinforced aluminium alloy nanocomposites by a liquid process, 111-117, Copyright (2012), with permission from Elsevier; (c) Adapted from Mater. Sci. Eng. A, 244, S.W. lp, R. Sridhar, J.M. Toguri, T.F. Stephenson, A.E.M. Warner, Wettability of nickel coated graphite by aluminum, 31-38, Copyright (1998), with permission from Elsevier. Figure 2.8. TEM image showing needle shape Al4C3 (a) in Al5083+GNP composite, (d) Al5083+GO composite; (c) TEM image of the interface between GNP and Al5083, (e) high-resolution image of the selected area in (d) and (b) diffraction pattern of the Al [130]. Adapted from Composites Part A, 109, P. Shao, W. Yang, Q. Zhang, Q. Meng, X. Tan, Z. Xiu, J. Qiao, Z. Yu, G. Wu, Microstructure and tensile properties of 5083 Al matrix composites reinforced with graphene oxide and graphene nanoplates prepared by pressure infiltration method, 151-162, Copyright (2018), with permission from Elsevier. Interfacial reactions between Al and GN at high temperatures can result in the formation of Al4C3 during composite synthesis. The presence of this phase significantly impacts the mechanical, wear, and thermal properties of AGNCs. There remains an ongoing debate in the scientific community as to whether the formation of Al4C3 is beneficial or detrimental to the strength of the AGNCs. 21 Interestingly, while some researchers have reported a positive contribution of Al4C3 [131,132], others have reported a negative impact [133–135] on the mechanical properties of AGNCs. Shao et al. [130] synthesized AA5083 Al matrix composites reinforced with GNP/GO through BM, pressure infiltration casting (PIC), and subsequent extrusion at 480℃. Molten Al was infiltrated into the AA5083-GNP/GO preform at 550℃ under 10 MPa. The researchers observed needle-shaped Al4C3 in both composites, a finding further corroborated by HRTEM, as shown in Figure 2.8. However, the Al4C3 content in the AA5083-GNP composite (Figure 2.8(a) and (b)) was lower than that in the AA5083-GO composite (Figure 2.8(d)). It was proposed that defect regions in GNP/GO act as nucleation sites for Al4C3 formation. The higher defect concentration in GO led to the formation of a higher content of Al4C3 in the AA5083-GO composite. Conversely, the lower content of Al4C3 in AA5083-GNP composites potentially facilitated a stronger interface bond between AA5083 and GNP, as shown in Figure 2.8(c). The GNP composite exhibited a notably higher tensile strength of 331 MPa compared to the GO-based composite (tensile strength: 282 MPa). The high content of Al4C3 might be the reason for the poor mechanical properties in the AA5083-GO composite. Sun et al. [136] observed an increase in hardness with increasing rGO content from 0.15 to 0.30 and 0.50 wt.%, but poor tensile properties were observed in a composite of 0.50 wt.% rGO. They further explained that this decrease in tensile properties was due to the high content of brittle Al4C3 in the 0.50 wt.% rGO composite compared to the 0.30 wt.% rGO composite. Thus, while the presence of Al4C3 can improve hardness and tensile properties, higher levels of Al4C3 can lead to degradation in the tensile properties of AGNCs. High processing temperatures and defects in GN are favorable conditions for Al4C3 formation in AGNCs. Therefore, careful attention is required to maintain the structural integrity of GNs during AGNC processing. A thin Al4C3 layer (5–20 nm) is considered optimal for improving interfacial bonding, whereas thicker Al4C3 adversely affect mechanical properties [137]. Consequently, the structural integrity of GN is a crucial characteristic that influences interfacial reactions between Al and the GN (carbon) and must be maintained to achieve the desired amount and size of Al4C3, thereby resulting in improved alloy properties. To tailor the AGNCs' properties for potential applications, the challenges mentioned above must be overcome during processing. Researchers have explored various approaches and processing methods aimed at overcoming these challenges, which are reviewed in the following section. 22 2.3 Processing of AGNCs Over the past decade, researchers used various processing techniques, including powder metallurgy (PM), casting, friction stir processing (FSP), mechanical impregnation, selective laser melting, etc., to develop AGNCs. The next most used techniques followed by PM are casting and FSP (Figure 2.9). Other than these techniques, researchers attempted various approaches to synthesize AGNCs using 3D printing, high-pressure torsion, mechanical impregnation, etc. The selection of processing technique depends on various factors, such as GN content, GN distribution, intended application, and final shape/size of AGNC. This section provides an overview of these processing techniques and research efforts to overcome the challenges outlined above and insights into individual studies. Figure 2.9. Contribution (%) of various processing methods used by researchers to develop AGNCs. 2.3.1 Powder Metallurgy (PM) The PM-based techniques for developing AGNCs involve three major steps: (i) incorporation of graphene into the Al matrix powder, (ii) consolidation of AGNC powder, and (iii) secondary processing (rolling, extrusion, etc.) of AGNCs. 2.3.1.1 Incorporation of GN into the Al Matrix The dispersion of GN in the Al matrix is a crucial step as it dramatically affects the mechanical, thermal, and electrical properties of the composite. GN should be dispersed uniformly and without 23 agglomeration in the Al matrix in order to take advantage of the nanostructure and high surface area of GN. Typically, the GNs available in the market are highly agglomerated due to the van der Walls and interaction forces (Figure 2.10). The interaction energy between interlayer graphene layers is 2 eV/nm2. To overcome this interaction energy and deagglomerate GN, a force of at least 300 nN/µm2 is required. The most common methods for dispersing GNs are BM and solution-assisted mixing (SAM). Researchers attempted different approaches using BM and SAM to overcome the agglomeration of GN in Al matrices, and these are discussed in the following sections. Figure 2.10. A schematic showing the agglomeration of graphene caused by van der Waals and interaction forces between the graphene sheets. 2.3.1.1.1 Ball Milling (BM) BM, also commonly called “mechanical alloying” or “mechanical milling”, is one of the most common methods for dispersing GNs in metal matrices. The movement of milling media (Steel, WC, Zirconia balls, etc.) during the milling of Al-GN powder is illustrated in Figure 2.11(a). In a high-energy planetary ball mill, the milling vial rotates opposite the supporting disc, generating centrifugal force. Energy transfer to the powder particles occurs in these mills by the impact of high-velocity milling media. During the milling of Al-GN powders, some particles of Al and GN become trapped between moving milling media. The impact forces of the milling media deform the Al powder particles due to their ductile nature, leading to the fracturing and delamination of GN sheets at the early stage of milling. The fractured /delaminated GN sheets adhere to the Al particles. The deformed Al particles become brittle with further milling due to work hardening, 24 resulting in fracture either by fragmentation or a fatigue failure mechanism (Figure 2.11(b)) [138]. With continued milling, composite powder undergoes repeated fracturing and cold-welding [139]. The fractured GN sheets are repeatedly enclosed and embedded into the cold-welded Al particles, leading to a layered structure of AGNC powder (Figure 2.11(c)). After milling for a certain duration, the rate of cold welding and fracture attains a steady-state equilibrium. Subsequently, this leads to the achievement of GN dispersion in the Al matrix. Figure 2.11. A schematic illustrating the (a) movement of milling media, (b) fracture, and (c) cold-welding mechanism during the milling of Al-graphene powder mixture in a planetary ball mill. Jiang et al. [48] conducted a study investigating the milling time effect on graphene nanosheets (GNSs) dispersion in the Al matrix. The powder milled for 4 h showed clusters of GNS on the Al surface (Figure 2.12(a)). However, uniform dispersion was achieved after 6 h of milling (Figure 2.12(b)). Further extension of milling time to 8 h (Figure 2.12(c)) and 10 h (Figure 2.12(d)) resulted in the reclustering or heterogenous distribution of GNS on the Al surface. This trend is evident in the distance between clusters of GNSs; the mean distance increased from 436±156 nm (6 h) to 831±296 nm (8 h) and 930±490 nm (10 h). The uniform distribution in 6 h milled samples can be 25 attributed to the steady-state equilibrium between welding and fracturing. The dominance of the cold-welding mechanism leads to an increase in particle size and the formation of clusters of Al and GN, whereas the dominance of the fracture mechanism causes a reduction in particle size. Figure 2.12. Effect of milling time on the dispersion of GNS in Al matrix (a) 4 hours, (b) 6 hours. (c) 8 hours. and (d) 10 hours [48]. Adapted from Composites Part A, 111, Y. Jiang, Z. Tan, R. Xu, G. Fan, D-B. Xiong, Q. Guo, Y. Su, Z. Li, D. Zhang, Tailoring the structure and mechanical properties of graphene nanosheet/aluminum composites by flake powder metallurgy via shift-speed ball milling, 73-82, Copyright (2018), with permission from Elsevier. Yu et al. [140] investigated the effect of milling time on GNS dispersion, defect creation in GNS, Al4C3 formation, mechanical and electrical properties of AA6063-0.3 wt.% GNS composite. GNS dispersion improved with milling time, achieving a uniform dispersion after 3 h. In a composite milled for 1 h, a thin Al2O3 film was observed between Al and GNS (Figure 2.13(a)), while direct bonding of Al and GNSs was observed after 3 h milled composite (Figure 2.13(b)). The Al2O3 26 layer on the surface of gas-atomized Al particles was disrupted during milling due to the milling balls' impact, completely diminished after 3 h of milling. The formation of Al4C3 was not observed after 1 h or 2 h of milling. Also, the size of Al4C3 increased from 10 to 20 nm to 50 nm as milling time increased from 3 h to 4 h, respectively. The increased Al4C3 size was mainly attributed to the rise in defect formation with milling. Additionally, the ID/IG of powders increased from 1.1 to 1.9 as milling time increased from 0 to 4 h, respectively. The diameter of GNSs and grain size of consolidated compacts decreased with increasing milling time up to 3 h. Consequently, composites milled for 3 h exhibited the highest tensile strength and electrical conductivity. However, the presence of a large amount of Al4C3 in the 4 h milled composite degraded the properties of AA6063. Figure 2.13. HRTEM image of Al6063-0.3 vol.% GNS milled for (a) 1h, showing the thin layer of amorphous Al2O3 between GNS and Al, and (b) 3h, showing the direct contact between Al and GNS [140]. Adapted from Carbon, 141, Z. Yu, W. Yang, C. Zhou, N. Zhang, Z. Chao, H. Liu, Y. Cao, Y. Sun, P. Shao, G. Wu, Effect of ball milling time on graphene nanosheets reinforced Al6063 composite fabricated by pressure infiltration method, 25-39, Copyright (2019), with permission from Elsevier. These studies have demonstrated that the defects in GN act as seed points for the nucleation of Al4C3. Therefore, it is essential to optimize the milling parameters to achieve the desired properties of AGNCs. The structural integrity of GN and its dispersion in Al can be controlled by tailoring 27 milling parameters/conditions, such as the ball-to-powder weight ratio (BPR), milling atmosphere, milling media, process control agent, milling speed, and time. These milling parameters/conditions play a crucial role in determining the final size and morphology of powder particles as well as the quality of dispersion. Figure 2.14. Milling parameters used by various researchers to disperse GNs in Al matrix (a) milling speed, (b) ball-to-powder-ratio, and (c) milling time. The data for this plot were extracted from the literature and plotted by the author. Literature indicates that there exists an optimum content of GN above which agglomeration occurs. Alipour et al. [141] observed agglomeration of GNPs along the grain boundaries in the AA7068 Al matrix with an addition of more than 0.5 wt.% GNP. Similarly, it was observed in the author’s work that the addition of more than 0.5 wt.% graphene resulted in agglomeration of graphene in 28 the Al matrix [29,30]. EI-Ghazaly et al. [38] also showed agglomeration of graphene in AA2124 with an addition beyond 3 wt.%. The milling parameters used by various researchers to disperse GNs in Al matrices are shown in Figure 2.14. The plots clearly indicate that the majority of researchers have employed milling speeds, BPR, and milling times in the range of 200 to 400 rpm (Figure 2.14(a)), 5:1 to 20:1 BPR (Figure 2.14(b)), and 1 to 6 h (Figure 2.14(c)), respectively. 2.3.1.1.2 Solution-Assisted Mixing (SAM) The SAM technique involves two main steps: (i) dispersion of GNs in a suitable liquid medium (such as acetone, ethanol, isopropanol, water, etc.) to break the van der Waals and interaction forces in GNs, thereby impeding agglomeration and promoting the exfoliation of graphene layers; and (ii) mixing the dispersed GNs with Al powders using ultrasonication or stirring. Ultrasonication involves the use of sound energy with frequencies exceeding 20 kHz to agitate particles and induce cavitation in the solvent, thus achieving uniform mixing. The sound energy in the solvent during sonication can be generated either with an ultrasonic bath or probe. Stirring is typically performed in mechanical agitation systems. The level of contamination in SAM is relatively low compared to BM due to the absence of a milling jar or media. However, some contamination from the liquid medium may still occur. SAM involves low shear or impact forces compared to BM; thus, it has a negligible effect on changing the particle morphology or size. While the formation of defects in the form of open edges or vacancies is possible, the level of defect creation during SAM is generally lower than during BM. Additionally, the exfoliation of graphene layers can be achieved in a suitable solvent at higher frequencies/speeds [32]. The binding energy between the graphene sheets in graphite is about 2 eV/nm2, and a force of 300 nN/µm2 is required to exfoliate graphene [141]. Sonication and mechanical agitation can easily overcome these binding energies, mainly due to the high local shear stresses developed during processing. For example, Latief et al. [142] dispersed GNP in acetone for 1 h at 50 kHz to exfoliate/deagglomerate GNP, followed by 5 h sonication with Al powder to achieve a uniform dispersion of GNP in the Al matrix. Similarly, Wang et al. [46] exfoliated GO into single-layered sheets by ultrasonication in deionized water, followed by the uniform adsorption of GO onto the Al surface via mechanical stirring. Liu et al. [143] studied the effect of solvent on the dispersion of GO and found that stirring in acetone provided better dispersion than ethanol-water and ethanol. Thus, effective dispersion and good structural integrity 29 of GNs with few/single layers can be achieved through SAM by optimizing the parameters such as mixing time, speed/frequency, and using a suitable liquid medium. Fattahi et al. [144] sonicated the mixture of Al + GNSs for 30 min prior to the BM and achieved uniform dispersion. Sonication of GN prior to milling deagglomerated the GN. In another study, Saboori et al. [145] dispersed GNPs in ethanol for 45 min before mechanical agitation. It has been suggested that the flake powders are geometrically more compatible with nano reinforcements than spherical powders [45]. Thus, flaky powder particles promote good interfacial bonding with GNs, enhancing effective load transfer from Al to the GNs and further improving the mechanical properties of AGNCs. Consequently, the flake-PM approach involves three steps: (i) synthesis of flake powders by BM, (ii) incorporation of reinforcement into flake powders, and (iii) consolidation of composite powders. Many researchers, including Wang et al. [46], Li et al. [47], Jiang et al. [48], and Pu et al. [146], have successfully employed this approach to disperse GNs in Al matrices. The SEM image showing the uniform absorption of GO on the Al flake surface prepared by the flake-PM approach is provided in Figure 2.15(a), with energy dispersive spectroscopy (EDS) further confirming the presence of GO on the Al flake (Figure 2.15(b)). Figure 2.15. (a) SEM image showing uniform adsorption of GO on Al flake, and (b) EDS of the selected area in (a) [46]. Adapted from Scr. Mater., 66, J. Wang, Z. Li, G. Fan, H. Pan, D. Zhang, Reinforcement with graphene nanosheets in aluminum matrix composites, 594–597, Copyright (2012), with permission from Elsevier. 30 2.3.1.2 Consolidation of AGNC Powder Following the incorporation of GNs into the Al powder matrix, consolidation is the crucial step in PM processing of AGNCs. PM processes can be broadly classified into two categories: (i) pressureless sintering and (ii) pressure-assisted sintering. 2.3.1.2.1 Pressureless Sintering (PLS) In PLS processes, the composite powder is consolidated by compaction and subsequent sintering. In general, the compaction process involves applying external pressure to the powder at room temperature, commonly known as cold compaction. This pressure applied during compaction deforms the powder particles, facilitating the formation of mechanical bonds/cohesion between them. However, the resulting green compact typically possesses high porosity, leading to the poor compact’s mechanical and thermal properties. Thus, additional thermal treatment, often called sintering, is required to enhance these properties. Sintering is typically carried out below the melting point (0.6-0.8Tm) of the matrix (Al) in a suitable furnace for an appropriate time. This enables diffusion-driven metallurgical bonding between powder particles, resulting in a dense microstructure that imparts structural integrity and mechanical strength to the compact. The sintering process can be further categorized into conventional and microwave sintering processes based on the heating mechanism. In conventional sintering, the green compact is heated through conduction, convection, or radiation mechanisms. Consequently, compacts experience a temperature gradient from their surface inwards and thus may require longer processing to reach the desired homogenous temperature. In contrast, microwave sintering involves the conversion of electromagnetic energy into thermal energy within the material, leading to rapid and uniform heating. As a result, microwave sintering typically yields a more uniform and fine microstructure compared to conventional sintering methods [147]. Latief et al. [148] used a PLS process to synthesize Al-1, 3, and 5 wt.% GNPs composites and studied the effect of sintering temperature (400, 500, 600℃) on the mechanical properties of AGNCs. The hardness and compressive strength of composites increased with increasing sintering temperature, reaching the highest value at 600℃. Similarly, studies by Saboori et al. [149], Baig et al. [150], and Gurbuz et al. [151] demonstrated that increasing the sintering temperature improved sinterability and resulted in enhanced properties. This enhancement is attributed to the increased diffusion of atoms, leading to a denser microstructure [152]. Bustamante et al. [153] and 31 Gurbuz et al. [151] observed an increase in hardness with sintering time up to 2 h and 3 h, respectively, beyond which further processing resulted in hardness degradation. The degradation of hardness at longer sintering times was associated with grain coarsening caused by enhanced atomic diffusion [154]. Additionally, Saboori et al. [149] and Baig et al. [155] demonstrated that the density or hardness of AGNCs increased with increasing compaction pressure up to 500 MPa. However, any further increase in pressure resulted in property degradation. This degradation was attributed to the damage of GNs structure at high pressures or entrapment of air in the composites, which could not be removed during subsequent sintering. 2.3.1.2.2 Pressure-Assisted Sintering (PAS) PAS processes use the simultaneous application of both pressure and temperature to consolidate the Al-GN composite powder. As a result, both compaction and sintering occur simultaneously. A variety of PAS techniques have been explored for consolidating nanocomposites. Among them, HP and SPS have been used extensively to consolidate AGNCs. In Both HP and SPS methods, a uniaxial press coupled with a heating device is typically utilized. The application of pressure together with temperature helps increase the density, leading to a reduction in the distance over which diffusion must occur. Consequently, fully densified compacts can be produced by PAS processes at lower temperatures and shorter sintering times compared to PLS processes. The main difference between HP and SPS lies in the heating mode. In HP, sample heating occurs through heat transfer via conduction from the heating elements to the sample, akin to conventional sintering. Conversely, in SPS, sample heating is achieved through Joule heating, generated by the application of a pulsed direct current directly onto the powdered sample. Direct heating in SPS helps achieve high heating rates and uniform heating, which further leads to the rapid consolidation of powder samples at low compaction pressure. In the case of SPS, heat dissipation occurs at the contact points of the powder particles, which helps maintain the nanostructure in sintered compacts [156,157]. According to the literature [31], the consolidation of Al-GN composite powder by SPS was usually carried out in the pressure and sintering time range of 30-50 MPa and 1-20 min, respectively, which are lower than in conventional and HP techniques. Furthermore, the commonly used heating rate in SPS is in the range of 50-200℃/min, compared to 5-15℃/min in other methods. 32 Ghasali et al. [158] synthesized Al-GNP composites through conventional, microwave sintering, and SPS techniques and investigated the effect of processing techniques on the mechanical properties of the composites. SPS-processed composites exhibited higher density and bending strength when compared to conventional and microwave-sintered composites. This was mainly attributed to the direct heating of powdered samples and the fine as-sintered microstructure achieved during SPS. However, the SPSed composites showed lower hardness values than conventional and microwave-sintered composites, which was attributed to the presence of a hard Al4C3 phase in both conventional and microwave-sintered composites. Sintering at higher temperatures during conventional (700℃) and microwave sintering (600℃) compared to SPS (420℃) was the main reason for Al4C3 formation. Also, Zhou et al. [159] observed the formation of Al4C3 in Al-0.25 vol.% rGO at an SPS temperature of 610℃, but not in the sample sintered at 560℃. The lower activation energy at high temperatures makes Al react with rGO (especially at defected regions), which leads to the formation of Al4C3. Li et al. [160] observed that Al4C3 did not form, although the HP was carried out for a longer time (4 h) at a high temperature (580°C). The use of defect-free and stable GNSs might be the possible reason for the absence of Al4C3 formation. It clearly shows that Al4C3 may form below the melting point of Al at the defect sites of GNs, but the carbide reaction between defect-free GNs and Al does not occur when the sintering temperature is above the melting point of Al. Zhao et al. [161] synthesized Al-1 wt.% GNF composites by SPS at 560℃ using 30 MPa consolidation pressure and studied the influence of SPS time on Al4C3 formation and mechanical properties of the composite. The increase in sintering time from 4 min to 8 min led to the formation of Al4C3, which further resulted in a decrease in elongation (↓32.3%), UTS (↓34.4%), and YS (↓36%) of the composite. The effect of SPS pressure (40 and 50 MPa) on the microstructure, density, and hardness of AGNCs was investigated in the author’s work [30]. It was observed that increasing compaction pressure from 40 to 50 MPa improved the density and hardness of composites. Increasing the pressure during SPS caused the powder agglomerates to break, increasing the contact area between powder particles, thus enabling enhanced diffusion. The enhanced diffusion promotes particle rearrangement, and interparticle bonding, leading to higher density and hardness. Consequently, optimization of the SPS parameters is required to achieve the desired properties. 33 2.3.2 Casting Casting is another popular processing route for synthesizing metal matrix composites due to its ability to produce large quantities of bulk material in a single run. However, challenges related to the difficulty in achieving uniform dispersion of GNs, poor wettability between Al and GNs, and the formation of Al4C3 limit its use in the processing of AGNCs. Usually, the processing of AGNCs through casting involves melting Al in a protective atmosphere, mixing GNs in molten Al under ultrasonication or mechanical stirring, followed by pouring, and solidification. Although various casting methods were explored for making metal matrix composites, stir casting has been the most common method of producing AGNCs. Generally, GNs tend to suspend on the top surface of the molten Al due to the difference in density between GNs [1.8–2.2 g/cm3] and Al [2.7 g/cm3]. This problem can be overcome to some extent using stir casting, where the GNs are mixed in molten Al using a mechanical stirrer (often made of graphite), which helps retain the GNs in suspension [162]. Boostani’s group has carried out significant research on developing hybrid composites made of Al+SiC+graphene through casting. They investigated the effect of graphene addition on the distribution of SiC in A356 [163] and A357 [164–166] matrix and the mechanical properties of the composite. They observed a uniform dispersion of SiC with 1 vol.% graphene addition, leading to 84% and 45% improvement in ductility and YS, respectively [163]. An enhancement of 258% and 350% in tensile ductility and YS was observed in the A357 alloy matrix with the graphene addition [164]. The addition of graphene improved the distribution of SiC by encapsulating SiC with onion-shaped shells of graphene. Additionally, mechanical stirring and ultrasonic vibration at 605℃ helped break solid dendrites and further induced a uniform distribution of SiC within the matrix. Furthermore, rod-shaped Al4C3 phase formation was observed in both samples. Interestingly the content of Al4C3 was lower in the AGNC, and the phase formed primarily at the defected regions of graphene. Mirjavadi et al. [186] developed Al-10Zn-3.5Mg-2.5Cu matrix composites reinforced with graphene (0-1 wt.%) using a composite master alloy by stir casting technique. First, an Al-graphene powder mixture was prepared using ultrasonication and BM. Subsequently, the milled powder was cold compacted and sintered at 560°C for 4 h to produce the composite master alloy. In the next step, the composite master alloy was introduced into the molten Al-10Zn-3.5Mg-2.5Cu alloy, 34 followed by stirring for 15 min, pouring, and extrusion. The extruded composites underwent T6 heat treatment. The grain morphology transformed into dendritic, and the grain size reduced with the graphene addition. However, graphene agglomeration was observed beyond 0.7 wt%, resulting in decreased mechanical properties of the composites. This fabrication method employing a composite master alloy appears effective in dispersing graphene in molten Al alloy, with the Al phase acting as a carrier for graphene. When the sintered composite master alloy was introduced into the Al melt, the Al phase carried the graphene within the melt and began to slowly melt. During this process, graphene was gradually released into the melt, and stirring facilitated the uniform dispersion of graphene. However, the use of pure Al powder may potentially dilute the Al-10Zn-3.5Mg-2.5Cu matrix, which can affect the mechanical properties of the resultant composites. Researchers also explored synthesizing AGNCs by PIC. In this technique, molten Al is pressurized into AGNC preforms, filling the spaces between the Al and GNs. Pressurization of the molten Al during PIC can be performed by an external pressure using inert gas or mechanical force [167,168]. However, studies on AGNCs using the PIC are limited, possibly due to the severe chemical reaction between molten Al and GNs. The first report on Al-20Si+GNF composites by PIC was published by Yang et al. [169] in 2017. First, GNFs were mixed with Al-20Si powder by BM at 300 rpm for 3 h, followed by consolidation of the powder mixture to prepare a preform. Then the Al-GNF preform was preheated to 500℃, while the Al-20Si was melted at 770℃. Subsequently, the molten Al-20Si was infiltrated into the preform under 5 MPa pressure for 10 min. The resultant composites were annealed at 400℃ for 2 h. The addition of 1.5 wt% GNF improved the tensile and bending strength of the Al-20Si alloy by 130% and 230 %, respectively. Despite the high processing temperatures, no Al4C3 formation occurred. This was primarily due to the presence of Si, which decreased the solubility of C in Al and improved the wettability of Al by GNF, thereby inhibiting the Al4C3 formation [170]. Moreover, the use of high-quality graphene (with fewer defects) can further inhibit Al4C3 formation [171]. In another study, Shao et al. [130] observed the formation of needle-shaped Al4C3 in Al-GNP and Al-GO composites prepared by PIC. However, the content and size of Al4C3 were siginificantly higher in the GO composites compared to GNP composites. The presence of more defects in the GO provides enhanced nucleating sites for the formation of Al4C3. All these studies clearly demonstrate that the structural integrity of GNs is crucial for achieving better properties of as-cast 35 materials. Production of AGNCs with minimal or no Al4C3 formation can be achieved by casting using high-quality GNs. 2.3.3 Friction Stir Processing (FSP) The FSP process is a solid-state or severe plastic deformation technique primarily derived from friction stir welding to refine surface microstructure. In this process, a non-consumable rotating tool is plunged into the workpiece surface, performing three simultaneous operations: deformation, stirring, and heating of the workpiece surface. Friction and plastic deformation induced by the tool heat and soften the material surface, leading to dynamic recrystallization and a fine-grain microstructure. Additionally, the rotating tool stirs the softened material from the adjoining side and deposits it on the trailing side, resulting in material intermixing in the processed region. Consequently, the surface properties of the workpiece are improved due to either alloying or grain refinement. Intermixing material under frictional shear forces provides an opportunity to make AGNCs using FSP by dispersing GNs on the Al surface. To prepare AGNCs by FSP, several grooves or holes are machined on the Al surface, which are then filled with GNs. These holes or groves act as reservoirs to prevent GNs from being swept away from the Al surface by the rotating tool [172–175]. Controlling the temperature during FSP is crucial as it affects the microstructure and mechanical properties of the workpiece. The temperature of the workpiece material can be controlled by altering process parameters such as force, rotational speed, and traverse speed. The first report on the processing of AGNCs with FSP was by Jeon et al. [175], published in 2014. They directly coated the GO-water colloid onto the surface of an Al alloy (5052-H32) instead of filling pre-machined holes/grooves. Then, FSP was performed with a rotational speed of 700 rpm and a traverse speed of 70 mm/min. The GO was reduced to rGO during the FSP, which was further confirmed by Raman spectroscopy. The increase in ID/IG after FSP indicates the formation of rGO, resulting in improved thermal conductivity and ductility compared to unreinforced Al. Khodabakhshi et al. [176,177] prepared an Al-Mg alloy (AA5052) matrix reinforced with GNPs using a multi-pass FSP process and investigated the effect of GNP addition on the microstructure and mechanical properties of AA5052. An improvement of 53% in hardness and more than a 3-fold improvement in YS was observed with an addition of 3 vol.% GNP compared to the base material. This was mainly due to microstructural refinement during FSP and uniform dispersion of GNPs. The grain size of Al-Mg alloys decreased from 10.7 µm (Figure 2.16(a)) to 9.7 µm 36 (Figure 2.16(b)) after FSP, and a further decrease in grain size to 2.1 µm (Figure 2.16(c)) was observed with GNP addition. The intense grain refinement after GNP addition was attributed to the combination of Zener pinning and particulate-stimulated nucleation during recrystallization. The interfaces in AA5052/GNP composite acted as nucleating sites during recrystallization. GNPs at the grain boundaries inhibited grain growth after recrystallization, thereby enhancing mechanical properties. Figure 2.16. Electron backscatter diffraction (EBSD) grain boundary maps of (a) Al-Mg alloy, (b) FSPed Al-Mg alloy, and (c) FSPed Al-Mg alloy+3 vol.% graphene [177]. Adapted from Surf. Coat. Technol., 335, F. Khodabakhshi, M. Nosko, A.P. Gerlich, Effects of graphene nano-platelets (GNPs) on the microstructural characteristics and textural development of an Al-Mg alloy during friction stir processing, 288-305, Copyright (2018), with permission from Elsevier. 37 In a study by Sharma et al. [178], AGNCs were prepared by PM-assisted FSP approach and the authors studied the effect of GNP content (2.5, 5 wt.%) and rotational speed (550, 1120 rpm) on the mechanical properties of Al6061, structural integrity, and distribution of GNPs. Prior to FSP, Al-2.5, and 5 wt% GNP composites were prepared by BM, compaction, and sintering. The resulting composites were utilized as consumable rods during FSP to develop surface composites. Uniform distribution of GNPs was achieved at 2.5GNP/550 rpm composite due to the uniform rod consumption during FSP. Conversely, non-uniform consumption of the rod was observed in the 5GNP/550 rpm composite, attributed to the high hardness of the rod. Additionally, surface composites exhibited higher ID/IG values compared to sintered composites, indicating an increase in disorder during FSP. 2.3.4 High-Pressure Torsion (HPT) HPT is a severe plastic deformation technique commonly used to refine the microstructure to enhance the strength of metals and alloys. In this technique, a disk-shaped sample is clamped between two anvils and subjected to torsional strain by rotating either the lower or upper anvil. Simultaneously, it experiences several GPa of compressive pressure at any temperature. The shear deformation of the disk sample occurs by surface frictional forces under a quasi-hydrostatic pressure. Typically, disk samples prepared by the HPT technique generally have a thickness of 0.2-0.5 mm and a diameter of 10-20 mm [179,180]. Researchers have used the HPT technique not only to refine microstructures [181] but also to consolidate powders of pure metals [182], alloys [183], and composites reinforced with Al2O3 [184], SiC [185], CNT [186] & GN [187–190] directly. One of the distinctive features of HPT is its ability to apply torsion under a high compressive pressure at any temperature, enabling the consolidation of powders into high-density samples while maintaining the nanostructure. Additionally, the consolidation of AGNCs at lower temperatures (25 to 200℃) using HPT has been observed to mitigate undesirable interfacial reactions, such as Al4C3 formation, commonly encountered during sintering. The earliest report on the processing of AGNCs using the HPT technique was conducted by Czeppe et al. [187,189]. Initially, they prepared Al-0.05 wt.% graphene powder by BM at 150 rpm for 7.5 h. Subsequently, the powder was consolidated by HPT at a pressure of 4 GPa with 5/10 rotations. Comparatively, the ball-milled and HPT-treated samples exhibited higher hardness values than the directly consolidated sample. However, the addition of graphene did not yield a statistically 38 significant improvement in hardness, likely due to the small amount (0.05 wt.%) of graphene added. It was observed that disorder (ID/IG) in graphene and grain refinement increased with the increasing anvil rotations from 5 to 10. Figure 2.17. STEM image of Al-5wt.% graphene composite HPT processed at (a, b) 298 K for 1 turn, showing the presence of graphene at grain boundary; (c, d) 473 K for 5 turns, showing the presence of graphene at both the grain boundary and within grain [188]. In another study, Huang et al. [188] prepared Al+5 wt.% graphene composites by HPT and investigated the effect of temperature and the number of HPT turns on graphene dispersion, the Al/graphene interface, grain size, and the hardness of the composites. Increasing the number of turns from 0 to 20 decreased the number of layers in graphene and improved their dispersion in Al. Scanning transmission electron microscopy (STEM) images of the composite processed at 298 39 K for 1 turn are presented in Figure 2.17(a,b). It was observed that the Al (1 1 1)/graphene (0001) interface exhibited numerous orientations, indicating graphene occupying the grain boundaries of Al in the 1-rotation sample. Moreover, the measured interlayer distance of graphene and Al was ∼0.35 and ∼0.23 nm, respectively. However, the sample processed at 473 K for 5 turns showed graphene not only at the grain boundary (Figure 2.17(c)) but also within the grain (Figure 2.17(d)). Plastic flow/diffusion and graphene exfoliation occurred in composites processed at 473 K for 5 turns, resulting in effective graphene dispersion and good interfacial bonding between Al and graphene. The composites processed at 298 K for 20 turns exhibited maximum hardness and tensile strength. Furthermore, an increase in the processing temperature to 473 K led to a deterioration in mechanical properties. The increase in the degree of grain coarsening with increasing processing temperature was the main reason for the lower hardness and tensile strength observed at 473 K. 2.3.5 Selective Laser Melting (SLM) SLM is one of the most versatile additive manufacturing techniques, capable of producing high-density components directly from powders. In SLM, computer-aided design (CAD) data containing a series of steps from component initiation to removal is uploaded to the SLM equipment, detailing the geometric characteristics of the component. Subsequently, a thin layer of powder, typically 20 to 100 µm in thickness, is evenly spread onto the substrate via a powder delivery system. A high-energy laser beam is then utilized to scan selected areas of the powder surface in accordance with the CAD data. Upon completion of scanning a powder layer, the substrate platform lowers by a single powder layer thickness along the z-axis, followed by the deposition of a fresh layer of powder onto the previously scanned layer. This layer-by-layer process continues until the entire component is built up. The heat generated during laser scanning is sufficient to melt the powder particles, leading to the formation of a melt pool. As soon as the laser beam leaves the melt pool area, solidification occurs by epitaxial growth of the partially melted grains from the previous layer or substrate. The cooling rate, temperature gradient, and direction of heat flow during SLM decide the grain size, morphology, and crystallographic orientation in the final component. Consequently, optimizing SLM parameters such as layer thickness, scanning speed, and laser power is important to achieving the desired microstructure. One advantage of the SLM process is the ability to fabricate near-net-shape components [191,192]. Literature indicates that the SLM technique can be utilized not only 40 for making components from metal or alloy powders [193,194] but also from composite powders [195,196]. The first report on the preparation of AGNCs by the SLM technique was by Hu et al. [197]. They used a customized SLM system to prepare Al-0.5, 1, and 2.5 wt.% graphene composites. A schematic of their SLM setup is shown in Figure 2.18. Prior to the SLM process, Al-graphene powder was prepared by BM at 200 rpm with a 5:1 BPR for 2 h. A layer of milled powder was then deposited onto an AA6061 substrate, followed by laser scanning at 90 W power, 3 mm/s scanning speed, and a pulse duration of 2.4 ms. Argon gas was continuously purged onto the laser scanning area to avoid oxide and nitride formation during the SLM process. The micro and nano hardness of Al improved by 75% and 70%, respectively, with 2.5 wt.% graphene addition. Similarly, Wang et al. [198] also employed BM and SLM techniques to prepare AlSi10Mg-0.5 wt.% GNP composites. However, there was no significant improvement in the tensile strength of Al with the addition of graphene, mainly due to the high porosity in the composite. It is possible that the SLM process parameters for composite synthesis were not optimized in this investigation. Bai et al. [199], Zhao et al. [200] and Wu et al. [201] investigated the tribological properties of SLM-processed AGNCs and observed improved wear resistance of Al with the addition of graphene. Despite the benefits of SLM, it may introduce tensile residual stresses in the surface or subsurface regions, thereby diminishing the fatigue performance of the material. Figure 2.18. A schematic of selective laser melting (SLM) setup used for the synthesis of AGNCs [197]. Reprinted from J. Alloys Compd., 746, Z. Hu, F. Chen, J. Xu, Q. Nian, D. Lin, C. Chen, X. Zhu, Y. Chen, M. Zhang, 3D printing graphene-aluminum nanocomposites, 269-276, Copyright (2018), with permission from Elsevier. 41 2.3.6 Electrical Resistance Heat-Assisted Mechanical Impregnation (ERHAMI) ERHAMI is a novel solid-state processing technique primarily utilized for developing surface composites by impregnating nano reinforcements onto the workpiece surface. The experimental setup comprises two water-cooled Cu electrodes, as depicted in Figure 2.19, wherein the workpiece is held, and the current is passed. Prior to ERHAMI, nano reinforcement is applied to the surface of the workpiece. Increased heat generation at the interface occurs through Joule heating due to the high resistance at the interface between the workpiece surface and the Cu electrode. Consequently, softening of the workpiece surface occurs, facilitating the subsequent impregnation of reinforcement onto the surface. The Cu electrode plays multiple roles, including holding the workpiece, conducting current, and applying pressure, with the aid of which pressure can be pneumatically controlled. Thus, adjustments in the current and heating time enable control over the surface temperature [202,203]. Figure 2.19. A schematic showing the experimental setup of the mechanical impregnation technique used for producing Al-GNP surface composites [202]. Reprinted from Mater. Des., 116, B. Sahoo, J. Joseph, A. Sharma, J. Paul, Surface modification of aluminum by graphene impregnation, 51-64, Copyright (2017), with permission from Elsevier. ERHAMI technique was first proposed by Sahoo et al. [203], who demonstrated this technique by impregnating graphite particles onto the surface of AA6063. First, a paste containing graphite 42 flakes dispersed in 10% polyvinyl alcohol (as a binder) was drop-cast onto the AA6063 surface with a thickness of 0.19 mm, followed by drying in a vacuum furnace. Subsequently, the substrate was positioned between Cu electrodes, and a current ranging from 1000–3500 A, along with a force of 1.5 N, was applied to the substrate for a fraction of seconds (0.2–0.6 s). The heat generated at the interface between the graphite coating and the Al substrate led to the formation of surface composites through the impregnation of graphite flakes onto the alloy surface. The increase in current and heating time contributed to an improved hardness of the surface composites, reaching a maximum hardness at 2000 A and 0.6 sec of heating. In another study by Sahoo et al. [204], the effect of carbonaceous material (CNT, graphene, and graphite particles) and graphite particle size on the mechanical properties of AA1100 surface composites was investigated. Al-CNT/GNP surface composites exhibited higher hardness compared to graphite-reinforced composites. The high specific surface area and aspect ratio associated with CNTs, and GNPs increase dislocation density and degree of grain refinement and improve Al hardness by inducing compressive stress. Additionally, CNTs and multi-layer graphene reinforced composites demonstrated superior wear resistance compared to graphite-reinforced composites due to their enhanced mechanical properties [205]. 2.4 Properties of AGNCs The primary goal of adding GN to the Al matrix is threefold: (i) increasing mechanical properties, (ii) improving tribological properties, and (iii) enhancing the electrical and thermal conductivity of the material. These three effects stem from the self-lubricating properties of GNs, as well as their higher strength, stiffness, and conductivity compared to the Al matrix. Available literature demonstrates that a significant improvement in hardness, wear resistance, YS, and UTS compared to monolithic Al is achievable. For example, automotive applications demand materials with high strength, wear resistance, and lightweight characteristics. Furthermore, in addition to increased strength, wear resistance, and lightweight, good conductivity is essential for certain components, such as conductors and battery enclosures/cases in electric vehicles. The mechanical, tribological, thermal, and electrical properties of AGNCs collected from the literature are summarized in the author’s review article [31]. This section discusses the effect of GN addition on the mechanical, and functional properties of Al. Additionally, it discusses the strengthening mechanisms contributing to the improvement of mechanical properties in AGNCs. 43 2.4.1 Mechanical Properties The mechanical properties of Al can be improved by introducing an optimal content of GN uniformly into the Al matrix. However, attaining these benefits requires fabricating a composite with low porosity (gas or shrinkage), since porosity is well known to degrade the mechanical, thermal, and electrical properties of an alloy [206–208]. Moreover, the presence of pores diminishes the thermal and electrical conductivity of the composites due to enhanced phonon scattering. Nevertheless, achieving a high relative density (i.e., a low porosity) in AGNCs poses challenges, especially with PM-based techniques. The literature indicates that the relative density of AGNCs typically decreases with increasing concentration of GN in the alloy [31]. This decline can be attributed to GN agglomeration and the significant difference in the melting points of the Al and the GN. Figure 2.20. Relative density of GN-reinforced Al matrix composites processed by pressureless sintering (conventional, microwave) and pressure-assisted sintering (hot pressing, hot isostatic pressing, spark plasma sintering). The data for this plot were extracted from the literature and plotted by the author. The tendency of bond formation between Al and GN is low due to the large difference in their melting points - Al (660℃) and GN (4625℃) [209]. Consequently, the presence of high melting 44 point GN at the particle/grain boundaries of Al restricts particle boundary diffusion, resulting in reduced sinterability [29]. Furthermore, the tendency for GNs to agglomerate at the grain/particle boundaries increases with higher GN content, leading to the formation of pores or voids. The relative density data of AGNCs produced by PM-based techniques is presented in Figure 2.20. Most research groups have employed a low weight percentage of GN (≤3 wt.%) with reported relative density values generally exceeding 92%. A majority of the composites (with ≤ 3 wt.% GN) consolidated by PAS techniques exhibited a relative density higher than 98%. The maximum amount of GNs used in PLS and PAS-fabricated composites was 5% and 15%, respectively. As discussed in Section 2.3.1.2.2, direct heating of powders and a simultaneous application of pressure and temperature in PAS techniques can result in higher densities even with a large content of GNs. A vast majority of research has been conducted to investigate the effect of GN addition on the mechanical properties of AGNCs. However, the extent of property improvement resulting from GN addition depends on various factors, such as GN dispersion, interfacial interactions, and load transfer capability. These factors, in turn, are strongly influenced by the processing techniques used for fabricating the composites. The mechanical properties most widely studied in AGNCs include hardness, YS, UTS, fracture strain, and ultimate compressive strength. The percentage improvement or deterioration in various mechanical properties of AGNCs compared to unreinforced Al, as a function of GN content, is depicted in Figure 2.21(a-e). In summary, the maximum percentage improvements achieved for various properties of AGNCs are as follows: 242% for hardness [210], 77% for elastic modulus [169], 182% for YS [70], 215% for UTS [211], 135% for fracture strain [210], and 586% for ultimate compressive strength [160]. However, it can be clearly seen from Figure 2.21 that there is a large scatter of percentage improvement in properties among AGNCs processed by similar processing techniques, even with similar GN content. This scatter in data can likely be attributed to differences in the size, surface area, and number of layers in the GNs utilized in these studies. Shin et al. [70,72] demonstrated that composite strength increased with increasing surface area of few-layer graphene. GNs with large surface areas and single layers provide a larger interfacial region between Al/GN where a more efficient load transfer can occur. Researchers have used GNs with varying numbers of layers, leading to scattering in the mechanical properties of AGNCs [31]. 45 Figure 2.21. Percentage of improvement or deterioration in mechanical properties of AGNCs (processed by different techniques) as a function of GN content: (a) hardness, (b) elastic modulus, (c) yield strength, (d) ultimate tensile strength, and (e) ultimate compressive strength. The data for this plot were extracted from the literature and plotted by the author. 46 Dutkiewicz et al. [212] observed a 50% improvement in the hardness of Cu-GNP composite as the GNP thickness decreased from 10-20 nm to 2-4 nm. As discussed in previous sections, Van der Waals forces between GN layers constitute weak intermolecular forces. With an increasing number of layers, the weak phase volume fraction in the composite rises, resulting in a lower load-bearing capacity and leading to a premature failure of the composite at a stress lower than theoretically predicted. Additionally, large-sized GN also act as defects in the composite due to poor interfacial bonding. Therefore, composites reinforced with a single layer of GNs with a large surface area are expected to offer superior strengthening compared to those reinforced with a few layers of GNs. The simultaneous improvement in hardness by 242% and the fracture strain by 86% compared to unreinforced Al was achieved with in-situ synthesis of GN by a one-step casting process that resulted in excellent interfacial bonding and uniform distribution of GN [210]. Figure 2.22. Optimum GN content reported in individual studies of AGNCs fabricated by PM, casting and FSP. The data for this plot were extracted from the literature and plotted by the author. A trend deduced from the examination of the mechanical properties suggests that there is an optimal GN content for improving the alloy’s properties. Beyond this optimum content/concentration, the percentage improvement in properties tends to decrease with increasing GN content, as shown in Figure 2.21(a-e). This phenomenon is mainly ascribed to the 47 agglomeration or clustering of GNs at higher concentrations, resulting in weakened regions that facilitate the formation of microcracks and pores. The optimal GN content, representing the maximum improvement in the mechanical properties of AGNCs reported in individual studies, is compiled and presented in Figure 2.22. It is evident that the optimal GN content is scattered between 0 and 6 wt.%. However, the majority of studies indicate a value of ≤ 1 wt.%, suggesting that GNs should not be added beyond this concentration. 2.4.1.1 Strengthening Mechanisms in AGNCs In general, there are two types of mechanisms associated with strengthening in AGNCs: (i) direct strengthening, caused by the effective load transfer from the Al matrix to the GN (load transfer strengthening), and (ii) indirect strengthening, resulting from the influence of GNs on the deformation mode and Al matrix microstructure (dislocation-based and grain refinement strengthening). The probable reasons for this anomaly can be attributed to the non-uniform dispersion of GNs in the Al matrix, the formation of interfacial phases, the presence of defects in the end product, and a large number of carbon layers in GN, which in turn affects the strengthening efficiency in AGNCs. This section provides an overview of the various active strengthening mechanisms relevant to AGNCs. In the load transfer mechanism, the strengthening effect of GN is greatly dependent on the efficiency of load transfer from the Al matrix to the GNs, which occurs through interfacial shear stresses developed during loading. Hence, strong interfacial bonding between Al and GNs is necessary to achieve theoretically predicted mechanical properties. Strengthening via the Orowan looping mechanism is based on the restriction of dislocation motion by hard particles or precipitates. When a dislocation interacts with particles sufficiently hard to impede dislocation movement, the dislocation bends around the particle, forming semi-circular dislocation loops known as “Orowan loops”. The loops hinder subsequent dislocation motion, resulting in improved material strength. The GNs are considered non-shearable dispersoids or precipitates, and thus “dispersion strengthening is active”. However, it is unclear whether the Orowan looping mechanism can operate at all or only under certain circumstances in AGNCs. Typically, the Orowan looping mechanism is more prominent in composites reinforced with low aspect ratio nanoparticles [213]. Consequently, the strengthening effect due to the Orowan looping mechanism is reduced in composites reinforced with high aspect ratio GNs. 48 In dislocation strengthening, the large mismatch in CTE and elastic modulus between Al (CTE: 23.5×10-6K-1; Elastic modulus: 62 GPa) and GN (CTE: -1×10-6K-1; Elastic modulus: 1 TPa) can lead to the formation of thermal residual stresses during processing, resulting in the generation of dislocations at the Al/GN interface. The increased dislocation density in the vicinity of GNs can enhance composite strength by restricting the movement of other dislocations. Additionally, the incompatible mechanical deformation characteristics between Al and GNs contribute to the generation of geometrically necessary dislocations at the Al/GN interface, thereby further enhancing the strengthening effect in AGNCs. In grain size refinement strengthening, the addition of GNs to the Al matrix can refine the microstructure either by acting as a nucleating site or by restricting grain growth. During AGNC processing, GNs offer more nucleation sites at the start of solidification. The presence of GNs at particle/grain boundaries further restricts grain growth by grain wrapping or pinning during the growth stage. The grain size of AGNCs reported in the individual studies is related to the GN content in Figure 2.23(a). Most of the studies demonstrated a decrease in Al grain size with GN addition and reported a grain size in the range of 0.07 µm to 250 µm for AGNCs. The maximum refinement in Al grains with the addition of GN reported in the literature was -80% [132], -87% [214], and -91% [215] with the PM, casting, and FSP routes, respectively (Figure 2.23(b)). The increased fraction of grain boundaries caused by the addition of GNs hinders the dislocation motion during the plastic deformation of the composite and leads to increased strength [154]. The trend in reported strengthening mechanisms by researchers in AGNCs is shown in Figure 2.23(c). It indicates that GNs contribute strength to the Al matrix predominantly by load transfer followed by the Orowan strengthening mechanism in AGNCs processed by PM and casting route. However, as mentioned earlier, the Orowan mechanism is more pronounced in composites reinforced with ceramic nanoparticles than with 2D and large aspect ratio GNs. The strengthening in FSP processed AGNCs is mainly through the grain refinement mechanism. It is well known that the grain size influences the strengthening mechanism; thus, the contribution of these mechanisms will be more or less pronounced at different grain sizes [1]. It is evident that the grain size of the matrix significantly influences the strengthening mechanisms that ultimately determine the reinforcement role in strengthening composites. Figure 2.23(c,d) indicates that load transfer and 49 Orowan strengthening dominate at coarser grain sizes, and grain refinement strengthening dominates at ultrafine grain sizes in AGNCs. Figure 2.23. (a) Grain size and (b) Percentage of increase or decrease in grain size of AGNCs as a function of GN content, (c) Trend in reported strengthening mechanisms (LT: Load transfer strengthening; OS: Orowan strengthening; DS: Dislocation strengthening; GR: Grain refinement strengthening) by researchers in AGNCs, (d) Average grain size of AGNCs (processed by different techniques) at an optimum content of GN reported in the individual studies. The data for this plot were extracted from the literature and plotted by the author. 2.4.2 Electrical and Thermal Properties In addition to its exceptional mechanical properties, graphene also has remarkable thermal and electrical properties, including a low CTE and high thermal and electrical conductivity, as seen in 50 Table 2.2. As a result, graphene addition to the Al matrix is expected to enhance these properties, along with improving mechanical properties. However, there are fewer reports investigating the impact of GN addition on thermal and electrical properties compared to those focusing on the mechanical properties of AGNCs. The first study to investigate the effect of GN addition on the thermal and electrical properties of Al was performed by Hahm et al. [216]. They fabricated Al-1 wt.% graphene by pulverization, BM, and sintering followed by extrusion and studied the effect of pulverization time (10 s, 1000 s) on the thermal and electrical properties of the composite. The composite reinforced with GN pulverized for 10 s and 1000 s exhibited enhanced thermal conductivity by 5% and 13%, respectively. Moreover, the electrical contact resistance decreased by 76.6% and 87.3% with a pulverization time of 10 s and 1000 s, respectively. The increase in pulverization time elevated the degree of exfoliation, resulting in a decrease in the number of layers in GN from several layers to approximately three layers. The improved dispersion and reduced number of layers in GN within the Al matrix with a pulverization time of 1000 s were the primary factors contributing to the improved thermal properties of the composite. Wang et al. [217] reported a maximum improvement in thermal and electrical conductivity at a 0.5 wt.% GN addition level, followed by a decline in these properties at higher addition levels. Therefore, as with the mechanical properties, there exists an optimal GN content to achieve maximum thermal and electrical conductivity in AGNCs. Several researchers have observed that the addition of GN beyond optimum content in AGNCs results in inferior thermal and electrical conductivity compared to unreinforced Al. This is primarily attributed to the tendency of GN to form clusters or agglomerates with increasing content. In research works by Guo et al. [218], Zhou et al. [219], and the author’s own research [30], no improvement in the electrical conductivity of AGNCs was observed with the addition of GN. Literature suggests that the improvement in thermal and electrical conductivity of AGNCs with GN addition is significantly smaller than theoretical predictions. In fact, some researchers have reported a deterioration in the electrical conductivity of Al with GN addition. The final thermal and electrical conductivity of AGNCs is influenced by various factors, including the GN type (pristine graphene, GO, and rGO), distribution of GN, directional alignment of GN within the matrix, presence of pores, presence of precipitates, interfacial and microstructural characteristics. 51 These factors, in turn, are influenced by the processing route used to synthesize GN and AGNCs. Saboori et al. [145] investigated the effect of processing technique on the thermal conductivity of AGNCs and observed that as-sintered composites exhibited superior thermal conductivity over as-extruded composites. This could be ascribed to the thermal resistance caused by increased point defects and grain boundaries during extrusion [220,221]. Pristine graphene typically possesses better conductivity than GO and rGO. The presence of oxygen functional groups and defects in GO and rGO reduce conductivity, as shown in Table 2.2. 2.4.3 Tribological Properties Tribological properties are important for components involved in repeated sliding and/or rolling movements with a contacting counterbody. In the automotive application domain, structural and mechanical components such as pistons, piston rings, brake rotors, bearings, connecting rods, etc., are subjected to various repeated movements and contact forces. Therefore, tribological behavior becomes a critical aspect, especially in the design of these components with Al and its alloys, as they generally exhibit poor wear resistance [222,223]. There are several ways to control the tribological behavior of components during operation, including the use of grease or liquid-type lubricants and/or a solid lubricant coating. However, the use of grease or liquid-type lubricants is not recommended due to environmental concerns. Solid lubricant coatings aim to improve the tribological behavior of component surfaces under extreme operating conditions. Unlike conventional lubricants or greases, solid lubricant coatings are capable of offering the desired level of durability and performance. Nevertheless, they also come with several drawbacks, such as poor adhesion, susceptibility to oxidation, replenishment challenges, low thermal conductivity, and limited service life and durability [224,225]. Some researchers attempted the addition of hard, micron-sized reinforcements, including TiB2 [226], SiC [227], AlN [228], TiC [229], and Al2O3 [230] into the Al matrix to improve tribological behavior. Although these reinforcements improved wear resistance, they posed challenges such as increased density, reduced ductility, and poor thermal conductivity [68,231]. Additionally, a large mismatch in CTE between Al and most ceramic particles led to the generation of thermal stresses at the interface, resulting in the detachment of hard particles from the surrounding Al matrix [38]. Subsequently, these detached particles acted as abrasives, further accelerating surface damage [224]. In recent years, graphite particles have emerged as solid lubricants suitable for the Al matrix 52 due to their excellent self-lubricating behavior [232–234]. However, the addition of soft graphite particles notably decreased the alloy’s mechanical properties (hardness, strength, etc.), thus limiting the alloy’s application to humid environments [38,218]. According to the Archard equation [235], the wear resistance of a composite is directly proportional to its hardness (Equation (2.1)). 𝑉 = 𝐾𝐹𝑛𝑆𝐻 (2.1) Where 𝑉 is the wear volume, 𝐾 is the wear coefficient, 𝐹𝑛 is the normal load, 𝑆 is the sliding distance, 𝐻 is the hardness. Maurya et al. [236] investigated the influence of carbonaceous material (graphite, graphene, CNT) on the wear performance of AGNCs. AGNCs exhibited superior wear resistance compared to graphite and CNT composites. A similar observation was made by Sharma et al. [237] in FSP-processed hybrid composites, where the Al-SiC-GNP hybrid composite showed better wear resistance than the Al-SiC-Graphite composite. The better wear performance of graphene-containing composites was attributed to their high hardness. As seen in Section 2.4.1, the addition of GN can improve hardness by 242% compared to unreinforced Al. Therefore, adding GN into the Al matrix is expected to enhance both the tribological and mechanical properties of the composite. Generally, the tribological performance of AGNCs is assessed based on CoF and wear rate or wear volume. The first investigation into the tribological properties of AGNCs was conducted by Ghazaly et al. [238]. They produced AA2124-1, 3, 5 wt.% GNS composites by PLS followed by extrusion. In their study, 3 wt.% GNS composite exhibited the highest wear resistance. Moreover, it is evident from Figure 2.24 that most studies demonstrated a reduction in CoF and wear rate with GN addition. This reduction can be primarily attributed to the improved load-bearing capacity of AGNCs with GN addition. Further, Zhang et al. [35] reported two effects, such as strengthening and lubrication of GN during dry sliding wear, as contributing to the enhanced tribological performance of AGNCs. Many researchers have reported the formation of continuous GN-rich tribofilms on the sliding surface in AGNCs [38,239]. In the strengthening effect, the formation of 53 stable GN-rich tribofilms reduces the wear rate of AGNCs by wrapping wear debris and bridging subsurface cracks, as shown in Figure 2.25(a) and (b). The GNSs wrapping wear debris can strengthen the tribolayer by reducing the available low-energy pathways for crack propagation, thereby improving wear resistance [172]. The self-lubricating ability of GNs decreases the CoF of AGNCs either by interlayer sliding or roller formation, as schematically illustrated in Figure 2.25(c) and (d). The weak interlayer bonds and the strong in-plane covalent bonds in GN enable interlayer sliding under the shear forces induced during sliding wear. Additionally, the formation of rollers from the packets of graphene planes on the counter surface can reduce the CoF of AGNCs. Figure 2.24. Percentage of increase or decrease in (a) CoF and (b) wear rate of AGNCs as a function of GN content. The data for this plot were extracted from the literature and plotted by the author. Nyanor et al. [240] reported a decrease in wear rate with increasing rGO content, and the 0.4 wt.% rGO composite exhibiting the lowest wear rate. However, the addition of rGO beyond 0.4 wt.% led to a deterioration in the wear properties of the composite. Figure 2.25 shows that in most studies, the addition of GN beyond an optimum content of 1 wt.% resulted in diminished tribological properties due to increased porosity and poor mechanical properties resulting from GN agglomeration. Baig et al. [39] investigated the effect of GN dispersion quality on the wear resistance of AGNCs by treating GN with two different surfactants: ethylcellulose and sodium 54 dodecyl benzenesulfonate. Composites treated with ethylcellulose (0.3 wt.% GN) and sodium dodecyl benzenesulfonate (0.5 wt.% GN) exhibited a 98% decrease in wear rate than unreinforced Al, marking the highest reported decrease in wear rate of an AGNC in the literature. Similarly, the highest reported decrease in CoF for an AGNC was 88.4%, achieved with the addition of 1.2 wt.% GN using BM to disperse GN in the Al matrix [241]. Therefore, achieving uniform dispersion of GN is crucial for improving tribological properties. Figure 2.25. A schematic depicting the GNs effects during sliding wear (a) wrapping effect, (b) bridging effect, (c) interlayer sliding, and (d) formation of rollers [35]. Adapted from Surf. Coat. Technol., 358, J. Zhang, Z. Chen, H. Wu, J. Zhao, Z. Jiang, Effect of graphene on the tribolayer of aluminum matrix composite during dry sliding wear, 907–912, Copyright (2019), with permission from Elsevier. 2.5 Chapter Summary The growing demand in the transportation industries for enhanced fuel efficiency and lower CO2 emissions led to the development of Al-based composite materials. GN has emerged as a promising nano-reinforcement for Al matrix composites due to its attractive properties, including high modulus (1 TPa), low density (1.8 to 2.2 g/cm3), high fracture strength (125 GPa), low CTE (-1 × 55 10-6K−1), self-lubricating ability, good electrical conductivity (2.02×104S/m) and thermal conductivity (4840-5300 W/mK). The literature demonstrates that AGNCs exhibit improved mechanical, tribological, thermal, and electrical properties than traditional Al alloys. However, research gaps remain since experimental measurements of properties often fall below the theoretically predicted values for AGNCs. This discrepancy is attributed to various processing challenges, including GN agglomeration and interfacial reactions between Al and GN. Various processing techniques and approaches have been researched over the past decade to overcome these challenges and exploit the full potential of GNs. The strengthening effect of GN strongly depends on its size and thickness (number of layers). Novel distribution and consolidation methods to obtain a uniform dispersion of GN, control over interfacial reactions, and reduce damage to GN during processing are key for the future adoption of these composites. In light of the current challenges in the development of AGNCs, this chapter reviewed diverse aspects of AGNCs, including processing challenges, fabrication techniques, and properties, with selected highlights summarized as follows: 1. The majority of the studies on AGNCs have focused on improving GN dispersion in Al matrices using various approaches and studied their effects on mechanical properties. So far, the dispersion of GN by BM and/or SAM (sonication or stirring), followed by consolidation using the PM route, has been the most viable option. This approach is favored for its ease of use and the ability to adjust parameters to address processing challenges. However, the PM route may not be suitable for making AGNCs on a large scale, which is essential for commercialization. In this context, the casting route is known for manufacturing large parts with intricate shapes at high production rates. Yet, limited research has been carried out on casting AGNCs. 2. Researchers addressed the processing challenges of AGNCs using various approaches in both the PM and casting routes. In the PM route, these approaches include (i) flake-PM, (ii) functionalization of GN, (iii) GN coating with Cu or Ni, (iv) using GO or rGO instead of pristine graphene, and (V) in-situ synthesis of GN. The underlying concept behind all these approaches was to achieve a uniform dispersion of GN in Al by modifying the surface characteristics of Al and GN. Similarly, in the casting route following approaches have been explored: (i) direct addition of graphene in Al melt, 56 (ii) addition of ball-milled Al-graphene powder composite, (iii) pressure infiltration, and (iv) addition of pre-densified compact (composite master alloy). To some extent, each of these approaches has aided in achieving uniform GN distribution in Al. Notably, the flake-PM, using composite master alloy and in-situ synthesis of GN approaches have demonstrated promising results, yielding relatively high-quality GN with a good interfacial bonding between Al and GN. However, scaling up the in-situ synthesis process for industrial-scale production of AGNCs while maintaining consistent quality and properties is challenging. Additionally, besides PM and casting, FSP has emerged as another technique used for producing AGNCs. Moreover, in recent years, researchers have explored novel techniques like ERHAMI, HPT, and SLM for fabricating AGNCs. 3. The outstanding improvements in the mechanical properties of Al with the incorporation of GN have been extensively reported. To date, the maximum percentage improvements achieved for AGNCs are as follows: ~242% for hardness, ~77% for elastic modulus, ~182% for yield strength, ~215% for tensile strength, ~134% for fracture strain, and ~586% for compressive strength. The majority of the researchers reported that the strengthening achieved by adding GN in PM and cast AGNCs was primarily through the load transfer and Orowan mechanism, while in FSP-processed AGNCs, grain refinement strengthening plays a significant role. 4. Literature suggests that the improvement in thermal and electrical conductivity of Al with GN addition is marginal, and in some instances, reduced conductivity in AGNCs has been reported. However, the incorporation of GN consistently enhances the tribological behavior of Al owing to its self-lubricating properties. It has been reported that the percentage improvement in specific properties of AGNCs decreases beyond an optimal content of GN due to GN agglomeration at higher concentrations. Most studies have identified an optimal GN content of ≤ 1 wt.% for achieving desired properties in AGNCs. 2.6 Problem Definition and Approach Based on the detailed literature review, it is evident that the addition of GNs to Al alloys may enhance the mechanical, tribological, thermal, and electrical properties while maintaining the 57 alloy’s lightweight characteristics. However, addressing the challenge of graphene agglomeration is crucial to achieving the expected properties in AGNCs. The flake-PM approach and using of composite master alloy in the casting route have shown promising results and possess the potential to fabricate uniformly dispersed graphene composites. Despite numerous studies focusing on the tribological behavior of AGNCs fabricated through the PM route, limited research exists regarding the tribological behavior of cast AGNCs. Specifically, there are hardly any reports found on investigating graphene's effect on the friction response, wear response and wear mechanisms of cast B319 Al alloy. The present work aims to fabricate B319-graphene composites through a novel distribution approach. This approach involves: (i) synthesis of composite master alloy through the flake-PM approach using BM followed by SPS, and (ii) introduction of the composite master alloy into B319 Al alloy melt via stir casting. The study investigates the effects of graphene addition on the microstructure, hardness, thermal conductivity, and electrical conductivity of B319 Al alloy. Additionally, the friction and wear behavior of B319-graphene composites are examined under reciprocating dry sliding conditions. In-depth characterization of the worn surface and its subsurface microstructure is conducted to understand the role of graphene in enhancing the wear resistance of B319 Al alloy and to identify the dominant wear mechanisms in these composites. 58 Chapter 3: Materials and Experimental Details This chapter outlines the experimental methods employed throughout this research. Section 3.1 describes the materials utilized for the research, while the fabrication process of composite master alloys is detailed in Section 3.2, followed by the description of the fabrication of B319-graphene composites in Section 3.3. Section 3.4 offers a detailed overview of the techniques used to characterize the B319-graphene composites and the procedures used for sample preparation, encompassing Spark OES, friction and wear testing, XRD, Raman spectroscopy, laser particle size analysis, optical microscopy, SEM, EPMA, and DTA/TGA. Additionally, the description of the techniques used to evaluate the properties of B319-graphene composites, including density, hardness, electrical conductivity, and thermal conductivity, are also presented in Section 3.4. 3.1 Materials In order to effectively incorporate graphene into the liquid B319 Al alloy melt, a composite master alloy was first fabricated via the PM route, and this master alloy was subsequently introduced into the liquid metal. To avoid significant dilution of the B319 Al alloy melt upon addition of the master alloy, a “pre-alloyed” feedstock powder material (F357 Al alloy powder) with a composition similar to that of the B319 Al alloy was used for the fabrication of the master alloy. The F357 Al alloy powder (D50: 11 µm) used in this study was obtained from AP&C (GE additive, Canada). Graphene (surface area: 351 m2/g), provided by Zentek Ltd., Canada, was used as a reinforcement. It was synthesized by a mechanical exfoliation method. A commercial B319 Al alloy ingot was used for all casting experiments. 3.2 Fabrication of Composite Master Alloys BM of F357 Al alloy powder and graphene was carried out in a high-energy planetary mill (Fritsch P-5) using a zirconia vial (250 ml) and 10 mm zirconia balls. Milling at 200 rpm with a ball-to-powder weight ratio of 5:1 in ethanol was carried out. Initially, the F357 Al alloy powder was milled for 1 hour to change its morphology from spherical to flakes. Subsequently, graphene was added to the milled F357 Al alloy powder to yield a 5 wt.% graphene concentration. Probe sonication (Branson Digital Sonifier SFX 550) for 5 min was sufficient to homogenize the blend. Additional milling was performed for 1 hour to mechanically attach the graphene to the flaked 59 F357 Al alloy powder. The milled mixture was then dried at 120 °C for 2 hours under an argon atmosphere. Subsequently, the dried milled powder was consolidated in an SPS system (Thermal Technology LLC) using 60 mm graphite tooling. The SPS was performed at a temperature of 500 °C, a pressure of 30 MPa, a heating rate of 50 °C/min, and a holding time of 30 seconds under an argon atmosphere. 3.3 Fabrication of Cast B319-Graphene Composites The SPSed composite master alloy (F357-5 wt.% graphene) was subsequently used during a metal casting operation to fabricate cast B319 Al alloy composites reinforced with 0.1, 0.2, 0.3, 0.5, and 0.7 wt.% graphene. The total weight of the material (B319 Al alloy + composite master alloy) was maintained at 500 g for all castings. Figure 3.1. A schematic showing a hybrid distribution approach used in the present study to fabricate B319-graphene composites. Three cylindrical graphite molds (inner diameter = 25.4 mm, outer diameter = 50.8 mm, length = 127 mm) equipped with a thermocouple were employed for all casting trials. The graphite molds were preheated to 350 °C along with the raw B319 Al alloy ingots. Subsequently, the preheated 60 B319 Al alloy ingots were heated to 750 °C in a melting furnace, at which point the preheated composite master alloy was introduced and mechanically stirred into the melt for 3 min without disrupting the melt surface. Thereafter, the melt was isothermally held for 5 min and then cooled to the pouring temperature of 725 °C. Figure 3.1 summarizes the fabrication approach used to prepare B319-graphene composites. 3.4 Characterization of B319-Graphene Composites 3.4.1 Spark Optical Emission Spectrometer (OES) The chemical composition of F357 Al alloy and as-cast B319-graphene composites was determined using a spark OES (Q4 TASMAN, Bruker). For all samples, an Al-Global method was used. However, for the B319-graphene composites, an iron (Fe) method was specifically employed to determine the carbon content, since the Al-Global method does not contain carbon element. Prior to measurement, the samples were polished to a mirror finish and subsequently cleaned with ethanol to remove contaminants. Table 3.1 and Table 3.2 show the chemical composition of the F357 Al alloy and the as-cast B319-graphene composites determined by spark OES, respectively. Table 3.1. Chemical composition of F357 Al alloy, in wt.%. Alloy Si Cu Fe Zn Mn Mg Ti Cr Sn C Al F357 6.8 0.01 0.08 0.01 0.001 0.46 0.06 - - - Bal. Table 3.2. Chemical composition of B319-graphene composites, in wt.%. Graphene content Si Cu Fe Zn Mn Mg Ti Cr Sn C Al 0 wt.% 8.47 2.23 0.45 0.47 0.27 0.30 0.17 0.04 0.005 - Bal. 0.1 wt.% 8.70 2.21 0.48 0.48 0.28 0.31 0.17 0.05 0.005 0.12 Bal. 0.2 wt.% 8.60 2.28 0.50 0.48 0.29 0.33 0.15 0.06 0.006 0.19 Bal. 0.3 wt.% 8.70 2.38 0.58 0.69 0.25 0.23 0.16 0.07 0.009 0.34 Bal. 0.5 wt.% 8.57 2.38 0.74 0.85 0.29 0.28 0.20 0.08 0.008 0.45 Bal. 0.7 wt.% 8.36 2.20 0.77 0.78 0.24 0.26 0.19 0.08 0.009 0.75 Bal. 61 3.4.2 Friction and Wear Measurement All the specimens for friction and wear testing were prepared by sequentially polishing as-cast coupons extracted from the cylindrical castings. Emery papers were initially used for polishing from 320 to 1000 P, followed by diamond paste (9, 6, 3, and 1 µm) and colloidal silica (0.05 µm) polishing steps. The dry-sliding friction and wear behavior of B319-graphene composites were investigated under linear reciprocating motion using a pin-on-disk tribometer (TRB3, Anton Paar) to simulate the conditions of a piston sliding in an engine block during operation, as shown in Figure 3.2. Although engine operating conditions typically involve lubricants (oils), understanding the wear behavior of this alloy and its composites in non-lubricated conditions is essential to ensure reliability and durability in a broad range of operating conditions and applications. Figure 3.2. Photograph of the pin-on-disk tribometer used in this study, showing its main parts. In this study, a 6 mm bearing steel ball (100Cr6, ∼63.8 HRC) was chosen as a counterbody due to its hardness comparable to a steel piston ring material (55-65 HRC). The steel ball was fixed in a holder, while the sample stage moved back and forth at a frequency of 5 Hz along a sliding length of 10 mm. The tests were conducted under normal loads of 5 N and 10 N with a total sliding distance of 300 m. A new ball was used for each sample during the wear test. The contact pressure was estimated to be in the range of 2.0-2.9 MPa and 3.3-3.9 MPa for 5 N and 10 N loads, 62 respectively. The nominal contact pressure was calculated by dividing the applied load by the counterbody contact area. The contact area of the counterbody was determined for each specimen by analyzing the ball-wear scar region using microscopy and Image J® software. The tangential friction force was recorded at a frequency of 50 Hz, and the CoF (µ) during sliding was simultaneously calculated according to Equation (3.1). Each sample was tested at least three times under identical experimental conditions. The CoF data was analyzed using InstrumX software (Version 10.0.11). µ =𝐹𝑓𝐹𝑛 (3.1) Where µ is the coefficient of friction, 𝐹𝑛 is the normal load, and 𝐹𝑓 is the tangential friction force. The wear volume of all samples was evaluated by two methods: (i) using the weight loss or gravimetric method, calculated with Equation (3.2), and (ii) using a 3D Laser Confocal Scanning Microscope. The weight of the samples and the counterbody before and after the wear test was measured using a high-precision analytical balance (Sartorius Entris II) having a 0.0001g resolution. Prior to checking the weight of the counterbody after the wear test, the counterbody was cleaned thoroughly with ethanol to remove any debris or material transferred from the sample during the wear test. The specific wear rate and cumulated dissipated frictional energy were calculated using Equations (3.3) and (3.4) [242], respectively. 𝑉 =𝛥𝑊⍴ (3.2) 𝑘 =𝑉𝐹𝑛 × 𝑆 (3.3) 𝐸𝑑 = µ𝐹𝑛ʋ𝑠𝑡 (3.4) 63 Where 𝑉 is the wear volume (mm3), 𝛥𝑊 is weight loss during the wear test (g), ⍴ is the density of the sample (g/mm3), 𝑘 is the specific wear rate (mm3/Nm), 𝐹𝑛 is the applied load (5 N and 10 N), 𝑆 is the sliding distance (300 m), µ is the coefficient of friction, ʋ𝑠 is the relative sliding speed (0.1 m/s) and 𝑡 is the sliding time (3000 s). In a 3D Laser Confocal Scanning Microscope (LEXT OLS5100, Olympus), first high-resolution in-focus images of the wear track from various depths were acquired via scanning the entire wear track in a raster pattern at 20x magnification using the stitching function in Olympus Data Acquisition Application; this process is called optical sectioning [243]. Subsequently, the intensity data collected at each point during the scanning process was used to construct a 3D image or z-stack of the wear track. This process aligns the sections based on their depth coordinates, creating a continuous representation of the wear track in three dimensions. Following this, the volume of the wear track was calculated by analyzing the reconstructed 3D image of the wear track using the Olympus Analysis Application. In the analysis application, the wear track region on the 2D image was accurately selected using the freeline tool. The software sectioned the 3D image of the wear track into multiple cross-sections using mathematical algorithms. Then, the software calculated the volume by summing the areas of the cross-sections at each depth and multiplying them by the distance between adjacent sections (known as slice thickness). This integration process finally yielded a volume of the wear track, which represents the material loss that occurred during the wear test. Subsequently, the specific wear rate was calculated using Equation (3.3). At least two measurements for the laser microscope method and three for the weight loss method were performed for each sample, and the average wear volume and specific wear rate values were reported along with the standard deviation. 3.4.3 Microstructural Analysis 3.4.3.1 X-ray Diffraction Analysis (XRD) The XRD analysis of graphene, ball-milled powder, composite master alloys, and cast B319-graphene composites was carried out using Bruker D8-Advance equipped with LynxEye silicon strip detector and Malvern-Panalytical Empyrean 3 equipped with a PIXcel 3D detector, both 64 employing Cu-Kα radiation. XRD data was collected in the range of 5-90° using a 0.02° step size. X'Pert HighScore Plus software was used to analyze the XRD patterns. 3.4.3.2 Raman Spectroscopy Raman spectroscopy, conducted using a Bruker SENTERRA II Raman Microscope, was performed on both the as-received graphene powder and the graphene present on the worn surface of the cast B319-graphene composite, employing a 532 nm excitation laser wavelength. Prior to Raman measurement, the as-received graphene was dispersed in ethanol using a probe sonicator (Branson Digital Sonifier SFX 550) for 2 minutes. Subsequently, a small amount of the dispersed graphene was dispensed onto the center of a clean glass slide using a transfer pipette. Then, a clean spatula was used to spread the sample into a thin layer on the glass slide, allowing it to dry at room temperature. Measurements were carried out on the dried sample using a 50x objective, 1.5cm-1 spectral resolution, integration time of 1500 ms, and 5 co-additions. The data was analyzed using OPUS 8.5 (SP1) software. 3.4.3.3 Laser Particle Size Analysis The particle size analysis of graphene and F357-graphene blend were measured using a Laser Particle Size Analyzer (PSA 1090, Anton Paar) equipped with two lasers (wavelength: 635 and 830 nm). The particle size analysis was performed in wet mode, with deionized water as the liquid medium. Prior to the measurement, the powder was dispersed in water within the sample chamber through simultaneous ultrasonication and stirring for 2 minutes. Subsequently, the data was processed using Anton Paar Kalliope Professional software (Version 2.30.4), employing the Mie scattering model to extract particle size distribution data. 3.4.3.4 Optical Microscopy For microstructural analysis of the as-cast samples, 10 mm thick samples were taken from the center of the castings, positioned at a length of 45 mm from the edge. The sectioned samples were polished in successive steps using SiC abrasive papers up to 1200 grit, followed by fine polishing with polycrystalline diamond paste (9 μm to 1 μm). Finally, the samples underwent final polishing with a 0.05 μm colloidal silica suspension. The optical micrographs were captured at 50x, 100x, and 500x with Zeiss Axio Observer Inverted Microscope. The average secondary dendrite arm 65 spacing (SDAS) and the morphological characteristics of Si were analyzed by Image J®, using at least 100 readings measured over 10 optical micrographs. Similarly, for grain size analysis, the polished samples were etched for 60 sec with Keller's reagent using the immersion technique. Keller’s reagent was prepared with 2.5 ml of nitric acid (HNO3), 1.5 ml of hydrochloric acid (HCl), 1 ml of hydrofluoric acid (HF), and 95 ml of distilled water. Grain size analysis was performed using Image J® software. At least 25 grains were considered for measurement, and the average grain size, along with the standard deviation, was reported. 3.4.3.5 Scanning Electron Microscopy (SEM) SEM analysis of B319-graphene composites was carried out using a Tescan Mira 3 XMU microscope coupled with an Oxford Instruments X-Max EDS detector. The samples for SEM analysis were polished following the standard metallographic procedure outlined in Section 3.4.3.4. For the subsurface analysis of worn surfaces, the specimens were prepared by cutting them perpendicular to the wear track with a precision diamond saw (LECO VC50), as illustrated in Figure 3.3. This was followed by cold mounting the cross-section sample in acrylic resin and polishing it to a mirror finish. Figure 3.3. Schematic representation of a wear track, showing the position of the cross-section examined in this work and sliding direction. 66 Before SEM analysis of the as-received graphene, it was dispersed in ethanol using a probe sonicator for 2 min to deagglomerate it. Subsequently, the dispersed graphene was dispensed onto double-sided carbon tape on a metal stub and allowed to dry at room temperature. Similarly, the ball-milled powders and wear debris were directly mounted onto double-sided carbon tape without ethanol dispersion. To minimize potential charging effects during SEM imaging, the cold-mounted samples and wear debris were coated with a 10 nm layer of platinum (Pt). SEM imaging was conducted at 15 kV to obtain high-resolution images, while EDS spot/area analysis was performed at 20 kV for better element detection. A working distance of 15 mm and a beam intensity of 15 were maintained for SEM-EDS analysis of all samples. Data analysis was carried out using Aztec software (Version 6.1). 3.4.3.6 Electron Probe Microanalysis (EPMA) Qualitative elemental distribution maps of as-cast microstructure and wear track cross-section were collected using an electron probe microanalyzer (EPMA, Cameca SXFiveFE) coupled with a wavelength dispersive spectrometer (WDS). For a higher detection sensitivity, the analytical conditions of 15 kV accelerating voltage, 200 nA probe current with a focused beam, dwell time of 50 ms, and step size of 1 µm were used for WDS elemental mapping. Similarly, WDS quantitative spot measurements were collected at 15 kV with a focused beam and a lower current of 20 nA to decrease the size of the electron interaction volume. All WDS quantitative spot measurements were calibrated using natural and synthetic reference standards provided by Micro-Analysis Consultants Ltd. The samples for EPMA were coated with 20 nm carbon, as it has minimal X-ray fluorescence interference with most elements of interest in EPMA analysis. Post-data processing was performed using the PeakSight 6.5 software. SEM and EPMA analyses were performed at the Fipke Laboratory for Trace Element Research (FiLTER) facility at the University of British Columbia, Okanagan. 3.4.3.7 Density Measurement The density of as-cast materials was measured using Archimedes’ principle [244], employing distilled water at room temperature as the fluid medium, and Equation (3.5). The theoretical density of the samples was calculated using the rule of mixtures, considering the densities of B319 Al alloy and graphene as 2.79 g/cm3 and 2 g/cm3, respectively. Finally, the relative density of 67 B319-graphene composites was determined using Equation (3.6). Prior to the density measurements, the samples were mechanically polished to remove any possible oxide layer formed on the surface, and they were subsequently dried. ⍴𝑠𝑎𝑚𝑝𝑙𝑒 =(𝑚𝑎)(𝑚𝑎 −𝑚𝑤)× (⍴𝐻2𝑂) (3.5) Relative Density = (⍴𝑠𝑎𝑚𝑝𝑙𝑒)(⍴𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙) (3.6) Where ⍴𝑠𝑎𝑚𝑝𝑙𝑒 is the sample density (g/cm3), 𝑚𝑎 is the sample weight in air (g), 𝑚𝑤 is the sample weight in water (g), ⍴𝐻2𝑂 is the density of water (1 g/cm3), and ⍴𝑡ℎ𝑒𝑜𝑟𝑒𝑡𝑖𝑐𝑎𝑙 is the theoretical density (g/cm3). 3.4.4 Thermal Analysis 3.4.4.1 Differential Thermal Analysis/Thermogravimetric Analysis (DTA/TGA) DTA measurements were carried out using Netzsch STA 449 F3 Jupiter. The DTA of the composite master alloy was carried out from 25°C to 750°C at a heating rate of 5 K/min and held for 10 min under a nitrogen atmosphere. The proximate analysis of graphene was conducted to determine its purity, following the method outlined in ASTM D7582-15 [245]. The sample was heated from 25°C to 120°C and held isothermally for 10 min under a nitrogen (N2) atmosphere, with the mass loss during this step considered as moisture. Subsequently, the sample was heated to 950°C and held for 15 min under an N2 atmosphere, with the mass loss during this step considered as volatiles. To determine the carbon content, oxygen (O2) was injected into the DTA chamber at 950°C and held for 45 min to burn the carbon in graphene; the mass loss during this step corresponded to the carbon content. The remaining mass of residue in the Al2O3 crucible was considered as ash. A heating rate of 10 K/min was used for the proximate analysis. DTA/TGA testing was performed using a sample mass of 15 mg and an alumina (Al2O3) crucible. The DTA and TGA data were analyzed using NETZSCH Proteus-Thermal Analysis software (Version 8.0.2). 68 3.4.4.2 Thermal Conductivity Measurement The specific heat capacity and thermal diffusivity of the B319-graphene composites (Diameter: 12.6 mm; Thickness: 2.5 mm) were measured from room temperature to 300 °C using laser flash analysis (LFA 467 HyperFlash, Netzsch). Subsequently, the thermal conductivity of the B319-graphene composites was determined using Equation (3.7). The data was recorded and processed using NETZSCH LFA analysis software (Version 8.03). At each temperature, three readings were taken, and the average value was reported with an error bar. 𝜅 = 𝛼𝐶𝑝⍴ (3.7) Where 𝜅 is the thermal conductivity (W/m.K), 𝛼 is the thermal diffusivity (m2/s), 𝐶𝑝 is the specific heat capacity (J/g.K), and ⍴ is the density (Kg/m3) of the sample. 3.4.5 Electrical Conductivity Measurement The electrical conductivity of the B319-graphene composites was evaluated via two approaches. First, the %IACS (International Annealed Copper Standard) was measured by Ether NDE SigmaCheck-2 Eddy Current tester, and the reported values represent an average of three readings. Additionally, a 4-point probe method was employed using the SBA 458 Nemesis, Netzsch equipment to measure the electrical conductivity of samples (Length: 14 mm, Width: 10 mm, Thickness: 2.5 mm) from room temperature to 300 °C. The data was recorded and processed using SBA 458 analysis software. Similar to thermal conductivity measurement, three readings were taken at each temperature, and the average value was reported with an error bar. 3.4.6 Hardness Measurement The hardness of the B319-graphene composites was measured with a Vickers hardness tester (SXHV-1000TA) with a load of 9.81 N and a dwell time of 15 s. All samples were indented at least 15 times, and the average hardness value was reported along with the error bars. Similarly, hardness measurements on the cross-section of the wear track were performed along the depth direction with a load of 0.245 N and a holding time of 15 s. A minimum gap of three times the indentation size was maintained between indentations during the hardness testing to avoid potential strain-hardening effects. 69 Chapter 4: Effect of Graphene on the Microstructure, Hardness, and Thermal Conductivity of Cast B319 Al Alloy This chapter focuses on the fabrication of B319-graphene composites using a composite master alloy, and it discusses the effect of graphene on the microstructure, hardness, and conductivity of B319 Al alloy. The microstructural analysis of the as-received graphene is provided in Section 4.1. Section 4.2 discusses the BM of F357 Al alloy-graphene powder and its morphological changes during milling. The optimization of graphene content in the composite master alloy, along with its microstructure, relative density, and hardness, is discussed in Section 4.3. Section 4.4 presents a discussion on the influence of graphene and its content on the microstructure of B319 Al alloy. Likewise, the effect of graphene on the hardness, electrical conductivity, and thermal conductivity of B319 Al alloy is explored in Section 4.5. Finally, Section 4.6 summarizes the key points derived from this chapter. 4.1 Microstructural Analysis of Graphene The as-received graphene exhibited a sheet-like morphology with a layered structure, as seen in Figure 4.1(a,b). Elemental mapping of graphene in Figure 4.1(c) and point EDS analysis of the region (A) marked in Figure 4.1(c) revealed the presence of C (97.7 at.%), O (2.1 at.%), and Si (0.2 at.%). It was reported that Si originated from the graphite precursor [246], while O was absorbed from the surrounding atmosphere during synthesis. Similarly, the proximate analysis of graphene revealed 0.6% of moisture, 3.4% of volatile matter, and 96% of carbon, as shown in Figure 4.2(a). Graphene particle size measured by Image J from SEM images showed an average particle size of 4.0 ± 0.9 µm, which was consistent with the particle size measured by the laser particle size analyzer (D50 of 4.3 ± 0.1 µm (Figure 4.2(b)). Raman spectroscopy clearly showed D, G, and 2D bands at ~1342 cm-1, ~1567 cm-1, and ~2683 cm-1, respectively (Figure 4.2(c)), which are characteristic peaks of graphene [32,247]. The D-band originates from defects or structural imperfections in the graphene lattice, where ‘D’ stands for defect or disorder. The intensity of the D-band is proportional to the amount of disorder. On the other hand, the G-band is associated with the in-plane vibration of sp2-bonded carbon atoms arranged in a hexagonal lattice, representing the fundamental graphitic or crystalline structure of graphene. 70 Figure 4.1. SEM images of graphene at (a) low, (b,c) high magnification and EDS mapping of Fig. (c), and the EDS data of region (A) marked in Fig. (c). Therefore, the ratio between the defect-induced D-band and the G-band (ID/IG) serves as a parameter for quantifying disorder in graphene. The 2D-band is related to the two-phonon process, and often, I2D/IG is used to estimate the number of layers in graphene [248,249]. In the present study, the ID/IG of the graphene used was ~0.4, indicating some disorder in the as-received graphene. Meanwhile, the I2D/IG of the graphene was ~0.37, suggesting a multilayer (5-10) of 71 graphene [249]. The multilayer of graphene was further supported by a sharp peak at 2θ = 26.35º in the XRD pattern (Figure 4.2(d)), showing a (002) plane with corresponding interlayer spacing of 0.338 nm, which was close to the interlayer spacing in perfect graphite (0.335 nm) [250]. The interlayer spacing of graphene was calculated using Bragg’s law, as shown in Equation (4.1). 𝜆 = 2𝑑𝑆𝑖𝑛(𝜃) (4.1) Where 𝜆 is the incident X-ray beam wavelength (Cu Kα: 0.154 nm), d is the interlayer spacing, and 𝜃 is Bragg’s angle (13.175º). Figure 4.2. (a) TGA curves for proximate analysis, (b) Particle size distribution curve, (c) Raman spectra, and (d) XRD pattern of raw graphene powder. 72 4.2 Microstructural Analysis of Ball Milled F357-Graphene Powder Mixture The as-received F357 powder had a spherical morphology with an average particle size of ~11 μm, as shown in Figure 4.3(a). The spherical particles were changed to flake morphology after 1-hour of milling, resulting in a mixture of spherical and flake-shaped F357 powder (Figure 4.3(b)). During BM, the F357 powder is subjected to intense impact forces caused by the collision of the milling balls and the powder itself. These impact forces led to the plastic deformation and shearing of the powder particles, ultimately resulting in the formation of flakes or thin layers [138]. The area EDS maps of the as-received F357 powder reveal the presence of Al, Si, and Mg elements, as shown in Figure 4.3(c). The EDS data of the F357 powder are consistent with the spark OES data (Table 3.1). Figure 4.3. SEM images of F357 Al alloy powder (a) As received, (b) 1-hour ball milled, and (c) EDS area mapping of Figure (a). 73 To optimize the graphene content in the composite master alloy, various amounts of graphene, including 2.5 %, 5 %, and 10 wt.%, were added to the 1-hour milled F357 powder. The SEM images and corresponding EDS maps of the 2-hour milled F357-graphene powder were presented in Figure 4.4(a-c). The EDS maps of 2-hour milled F357-graphene powder revealed the presence of C element in addition to Al, Si, and Mg observed in the as-received F357 powder. It is evident from Figure 4.4(a-c) that graphene is uniformly distributed in F357-2.5 wt.% and F357-5 wt.% graphene powder. Since the flaky powder particles were geometrically more compatible with graphene, the graphene sheets were attached to the powder flakes during the probe sonication. The absorbed graphene dispersed in Al flakes with further milling, resulting in uniformly distributed graphene in F357-2.5 and 5 wt.% graphene samples (Figure 4.4(a,b)). However, F357 with 10 wt.% graphene indicated the presence of clusters of graphene (Figure 4.4(c)), suggesting that 5 wt.% is the optimum amount of graphene for the size of F357 powder and the processing approach used in the present study. The effect of graphene content and milling on the particle size distribution curves of F357-graphene powder is depicted in Figure 4.5. It is evident that the size distribution of F357 powder widened after 1-hour milling, a phenomenon attributed to the formation of flakes observed in 1-hour milled powder (Figure 4.3(b)). The average D50 and D90 of F357 powder increased from ~11 µm to ~14 µm and ~20 µm to ~38 µm, respectively with 1-hour milling (Table 4.1). With 2-hour milling and graphene addition, the size distribution further expanded due to an increased number of flakes and their sizes resulting from the increased degree of deformation of the powders. However, all the F357-graphene powder samples showed similar values of D10 and D90, confirming that the increase in graphene content from 2.5 to 10 wt.% didn’t affect the size distribution. It was worth noting from Table 4.1 that the average D50 of the F357-2.5 and 5 wt.% graphene samples were ~19 µm, and it was slightly increased to ~22 µm with addition of 10 wt.% graphene. The increased D50 of the F357-10 wt.% graphene sample can be attributed to the agglomeration of graphene observed in SEM-EDS maps (Figure 4.4(c)). Furthermore, all the powder samples showed an average D10 of 4-6 µm, indicating that there is no fracturing phenomenon that occurred during the BM process, which can be due to the insufficient strain hardening of powder particles during 2 hours of BM. 74 Figure 4.4. SEM images and EDS area elemental mapping of ball milled powder (a) F357-2.5 wt.% graphene, (b) F357-5 wt.% graphene, and (c) F357-10 wt.% graphene. 75 Figure 4.5. Particle size distribution curves for F357 Al alloy and F357-graphene powder. Table 4.1. The average particle size of F357-graphene powder. Sample D10 (µm) D50 (µm) D90 (µm) F357-As received 5.5 (± 0.1) 11.4 (± 0.1) 20.5 (± 0.2) F357-1h milled 5.8 (± 0.1) 14.4 (± 0.2) 38.0 (± 1.2) F357-2.5 wt.% graphene-2h milled 5.6 (± 0.2) 19.0 (± 0.3) 50.8 (± 0.8) F357-5 wt.% graphene-2h milled 5.3 (± 0.3) 19.5 (± 0.4) 49.0 (± 0.6) F357-10 wt.% graphene-2h milled 4.3 (± 0.1) 21.6 (± 0.1) 50.7 (± 1.4) 4.3 Microstructural Analysis of SPSed Composite Master Alloy The optical microstructure of SPSed F357 composite master alloys, incorporating varying amounts (2.5, 5, and 10 wt.%) of graphene, is shown in Figure 4.6(a-c). The black regions in the optical 76 microstructures represent graphene, which is well distributed in the F357-2.5 wt.% (Figure 4.6(a)) and F357-5 wt.% graphene (Figure 4.6(b)) composite master alloys. However, noticeable agglomeration of graphene is observed in the 10 wt.% graphene composite master alloy (Figure 4.6(c)). Additionally, all composite master alloys exhibited similar relative densities, approximately 98-99%, as shown in Figure 4.6(d). Nonetheless, the hardness of the composite master alloys changed with increasing graphene content. The average hardness of F357-2.5 wt.% graphene was 50 HV, slightly increasing to 55 HV with the addition of 5 wt.% graphene, indicating a 10 % improvement. However, the hardness of the composite master alloy deteriorated to 38 HV at a higher graphene content (10 wt.%), attributed to the agglomeration of graphene (Figure 4.6(d)). Figure 4.6. Optical microstructure of SPSed composite master alloy: (a) F357-2.5 wt.% graphene, (b) F357-5 wt.% graphene, and (c) F357-10 wt.% graphene; (d) Relative density and hardness of composite master alloys. 77 Figure 4.7. Backscattered electron (BSE) images and corresponding EDS area maps of SPSed composite master alloys (a) F357-2.5 wt.% graphene, (b) F357-5 wt.% graphene, and (c) F357-10 wt.% graphene. 78 The BSE images and corresponding EDS area maps of SPSed composite master alloys are shown in Figure 4.7(a-c). The EDS maps reveal the presence of Al, Mg, Si, and C elements in the composite master alloy, which is consistent with the maps of powder mixtures observed in Figure 4.4. Furthermore, the C map confirms that the black regions in BSE images correspond to graphene, revealing a uniform distribution of graphene in F357-2.5 and 5 wt.% graphene composite master alloys Figure 4.7(a,b). Additionally, the agglomeration of graphene can be seen in F357-10 wt.% composite master alloy, as shown in Figure 4.7(c). It has been documented in the literature that there is always an optimum content above which agglomeration of graphene in the metal matrix occurs, irrespective of the distribution approach [31,47,124,125]. The preheated composite master alloys were further tested for their melting behavior by adding them into B319 molten metal. Interestingly, the F357-2.5 wt.% graphene and F357-5 wt.% graphene master alloys were melted fully, and the graphene was released into the molten metal. However, the composite master alloy with 10 wt.% graphene content didn’t melt completely even after stirring, as shown in Figure 4.8(a), which could be related to the thermal properties and behavior of the graphene. The presence of high thermal conductivity graphene can act as a heat conductor, potentially leading to localized cooling effects during the melting process. Additionally, graphene clusters may create regions of higher thermal resistance, further impeding heat transfer and hindering the melting process. The milled powder mixture of F357-graphene has been directly added to the B319 melt without SPS, resulting in the powder floating to the top surface and a cluster of graphene being observed, as shown in Figure 4.8(b). Generally, graphene tends to suspend on the top surface of the molten Al due to the difference in density between graphene (1.8-2.2 g/cm3) and Al (2.7 g/cm3) [31]. Mirjavadi et al. [49] used an Al-graphene composite master alloy to develop graphene-reinforced Al-10Zn-3.5Mg-2.5Cu alloy matrix composites by casting. However, in their study, the ball-milled Al-graphene powder mixture was consolidated using the conventional PM route (compaction followed by sintering). The composite with 0.7 wt.% graphene showed a 35% higher UTS compared to the unreinforced matrix. In another study, Yang et al. [169] also prepared Al-20Si reinforced graphene composite master alloy using BM followed by compaction and sintering. The composite master alloy was subsequently used to fabricate graphene reinforced Al-20Si matrix composites by PIC. Adding 1.5 wt.% graphene improved the tensile and bending strength of Al-20Si alloy by 130 and 230 %, respectively. 79 Figure 4.8. Photographs of (a) unmelted F357-10 wt.% graphene composite master alloy and (b) clusters of ball-milled F357-10 wt.% graphene powder. Figure 4.9. (a) XRD pattern of as-received F357 alloy powder and SPSed composite master alloy, and (b) DTA curves of as-received F357 alloy powder and F357-5 wt.% graphene milled powder. In the present study, considering the agglomeration of graphene and the melting behavior of the F357-10 wt.% graphene composite master alloy. The graphene content in the composite master 80 alloy was optimized to 5 wt.% and used this optimized alloy to make B319-graphene composites. The XRD pattern of SPSed F357-5 wt.% graphene composite master alloy shows characteristic peaks of Al, Si, and C while revealing that no undesirable compounds (such as Al4C3 and Al2O3) formed during the master alloy manufacture (Figure 4.9(a)). This was further confirmed with DTA analysis, which showed mainly two endothermic peaks corresponding to the melting of Al-Si eutectic and Al, as shown in Figure 4.9(b). Thus, these results confirm that the graphene used in the present study was stable up to 750°C and was successfully integrated into the composite matrix during sintering. 4.4 Microstructure and Phase Analysis of B319-Graphene Composites The optical micrographs of the B319 Al alloy in an ingot and as-cast states are shown in Figure 4.10(a) and Figure 4.10(b), respectively. The as-cast alloy exhibited a more refined dendritic structure (SDAS: 18.1 µm) compared to the ingot (SDAS: 38.0 µm) due to the higher cooling rates in the as-cast condition. The XRD patterns of the B319 Al alloy in both ingot and as-cast states show peaks corresponding to Al, Si, Al2Cu, and Fe-bearing phase (Figure 4.10(c)). Similarly, the XRD patterns of the B319-graphene (0-0.7 wt.%) composites in the as-cast state show peaks corresponding to Al, Si, Al2Cu, and Fe-bearing phases (Figure 4.11). However, none of the XRD scans exhibited statistically significant peaks corresponding to graphene since its content was below the XRD detection limit. It is well documented that defects in graphene readily act as reaction sites for carbide formation [31]. Given that the graphene used in this study contained a relatively low level of defects (ID/IG: 0.4), no Al- or Si-carbides were observed. XRD scans further suggest the absence of Al4C3 in all composites, which can be attributed to the high thermal stability of graphene or Al4C3 content below the XRD detection limit. Similar observations were also reported by other researchers [251,252]. SEM-EDS analysis (area elemental mapping) of as-cast B319 Al alloy showed the presence of α-Al dendrites, Al-Si eutectic, and Al-Cu and Fe-Mn rich phases, as shown in Figure 4.12. Furthermore, the point EDS analysis of secondary phases revealed that the stoichiometry of the Al-Cu and Fe-rich phases observed in Figure 4.12 correspond to Al2Cu, and Al15(Fe, Mn)3Si2 phases, respectively (Table 4.2). The BSE micrograph of as-cast B319 Al alloy, its corresponding 81 EPMA area elemental mapping (Figure 4.13) and point WDS analysis of phases marked in Figure 4.13 indicate the presence of Al, Al-Si eutectic, Al2Cu and Al15(Fe, Mn)3Si2 phases (Table 4.3). The EPMA results are consistent with the SEM-EDS data. Figure 4.10. Optical microstructure of B319 Al alloy in (a) ingot and (b) as-cast condition; (c) XRD patterns of B319 Al alloy in ingot and as-cast condition. Table 4.2. The average composition of secondary phases in B319 Al alloy, obtained by EDS analysis. Assigned Phase Element, at.% Al Si Fe Mn Cu Al15(Fe, Mn)3Si2 73.2 (±1.1) 11.4 (±0.5) 10.4 (±0.7) 4.9 (±0.4) - Al2Cu 71.1 (±0.3) - - - 28.4 (±3) 82 Figure 4.11. XRD pattern of as-cast B319 Al alloy-graphene (0-0.7 wt.%) composites. Figure 4.12. SEM image and EDS area elemental mapping of B319 Al alloy. 83 Figure 4.13. SEM-BSE image of as-cast B319 Al alloy and corresponding EPMA-WDS elemental mapping. Table 4.3. EPMA-WDS chemical analysis of phases in as-cast B319 Al alloy. Region in Figure 4.13 Atomic percent (at.%) Assigned phase Al Si Cu Fe Mn A 98.5 (±0.1) 0.6 (±0.1) 0.7 (±0.1) 0.02 (±0.01) 0.05 (±0.008) Al B 3.2 (±0.8) 96.5 (±0.7) 0.20 (±0.01) 0.001 (±0.003) 0.001 (±0.002) Si C 67.8 (±2.4) 2.9 (±1.1) 28.9 (±1.5) 0.16 (±0.07) 0.03 (±0.008) Al2Cu D 70.8 (±0.7) 11.9 (±0.3) 2.1 (±1.0) 10.7 (±0.5) 4.3 (±0.6) Al15(Fe, Mn)3Si2 The Fe-bearing phases were usually present as Chinese-script morphology and associated with the Al17(Fe3.2, Mn0.8)Si2 and Al19(Fe4, Mn1)Si2 stoichiometries, but commonly represented by Al15(Fe, Mn)3Si2 in the literature [253,254]. As seen in Figure 4.14, the Al2Cu was present in two morphologies: i) rod/needle and ii) globular shape. Furthermore, Al2Cu was present in blocky (Figure 4.14(a)) and eutectic (Figure 4.14(b)) forms. The possible reactions and the sequence 84 during the solidification of the B319 Al alloy are summarized schematically in Figure 4.15. Generally, B319 Al alloy solidification occurs mainly through two solidification stages: (i) nucleation of primary Al-dendrites and (ii) formation of Al-Si eutectic. However, alloying elements (Cu, Mg) and impurities (Fe, Mn) in this alloy result in a more complex solidification sequence [255,256]. Following the nucleation and growth of α-Al dendrite, interdendritic melt becomes enriched in Fe and Si, which leads to precipitation of the Al15(Fe, Mn)3Si2, or Al5FeSi phase depending on the Fe content and cooling rate. Due to the enriched Si interdendritic regions, the development of the Al-Si eutectic follows. At the end of the solidification, other secondary phases, including Al2Cu and Mg2Si, precipitates form from the remaining melt. It was observed that the Al15(Fe, Mn)3Si2 phase can also form after forming Al-Si eutectic, primarily on the Si particles. Furthermore, the intermetallic compounds (Al15(Fe, Mn)3Si2, Al5FeSi) and Si are the main nucleating sites for the precipitation of Al2Cu [257,258]. The SEM analysis at high magnification clearly shows the precipitation of blocky Al2Cu on the Al15(Fe, Mn)3Si2 phase (Figure 4.14(a)) and eutectic Al2Cu on the Si particles (Figure 4.14(b,c)). Thus, these microstructural observations of B319 Al alloy were consistent with published results [257,259]. Figure 4.14. (a) SEM image and (b) Optical microstructures of B319 Al alloy, showing the precipitation of Al2Cu on Al15(Fe, Mn)3Si2 and Si particles. 85 Figure 4.15. The possible solidification reactions and sequence in B319 Al alloy. The general microstructure of the B319-graphene composites captured at 100x is shown in Figure 4.16(a-f). The results indicate that the average SDAS of the B319 Al alloy decreased with graphene addition. Although the SDAS values of the B319-graphene composites fall within the standard deviation, the highest decrease (26 %) in SDAS from 18.1 μm to 13.4 μm, was recorded for 0.1 wt.% graphene addition. Interestingly, the SDAS of all B319 Al alloy-graphene composites was lower than that of the unreinforced B319 Al alloy, which was attributed to a higher cooling rate resulting from a higher heat extraction caused by the presence of highly conductive graphene in the molten alloy. Further, the refinement was caused by the presence of graphene between the growing dendrites of α-Al, thus contributing to grain growth restriction. Finally, adding graphene provided additional nucleation sites for α-Al, leading to a smaller SDAS and grain refinement (via heterogeneous nucleation), with a concomitant homogenization and dispersion of secondary phases throughout the microstructure. 86 Figure 4.16. Optical microstructure of cast samples captured at 100x (a) B319, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, (f) B319-0.7 wt.% graphene. Figure 4.17(a-f) shows that the average grain size of the B319 Al alloy decreased with the addition of graphene. The largest decrease of ⁓45% in grain size from 1658 to 913 μm was achieved with 0.7 wt.% graphene addition. However, the agglomeration of graphene (black regions) in the B319 Al matrix can be seen in composites with more than 0.3 wt.% graphene (Figure 4.17(e,f)). Literature shows that there is always an optimum content of graphene above which the agglomeration occurs. Alipour et al. [214] observed agglomeration of graphene along the grain boundaries in the AA7068 Al matrix with an addition of more than 0.5 wt.% graphene. It was also 87 noticed in the author’s other works that the addition of more than 0.5 wt.% graphene resulted in agglomeration of graphene in the Al matrix [29,30]. The agglomeration of graphene was mainly due to the strong van der Waals forces between the graphene sheets and the density difference between the B319 Al matrix and the graphene. Furthermore, the graphene clusters exhibit poor wettability by the B319 Al melt, resulting in the formation of pores or voids within the material. These pores contribute to a decrease in relative density [31]. It can be seen from Figure 4.18 that the relative density of the B319 Al alloy decreased from 99.82% to 97.09% with the addition of 0.7 wt.% graphene. Yang et al. [169] also reported a decrease in the relative density of Al-20Si alloy with increasing graphene content, observing a ~4% decrease in relative density at 3.5 wt.% graphene. In a study by Chen et al. [260], the relative density of Al2024 alloy decreased from 99.74% to 96.48% with 1.5 wt.% graphene addition, indicating a ~3% reduction. Figure 4.17. Macrostructure of cast samples etched with Keller's reagent: (a) B319, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene and (f) B319-0.7 wt.% graphene. 88 Figure 4.18. Relative density of B319 Al alloy-graphene composites. Figure 4.19(a-f) shows that in addition to influencing grain size and SDAS, the addition of graphene also changed the morphological characteristics of Si. The length and aspect ratio of Si decreased while its width increased, indicating the evolution of a fibrous form of Si (Figure 4.20 (a and b)). The average aspect ratio decreased by 59% with the addition of 0.1 wt.% graphene, but no significant difference in the aspect ratio was observed with higher graphene levels (Figure 4.20 (c)). Recently, Singh et al. [36] fabricated Al-Si (ADC12) alloy-rGO composites by stir casting and observed a similar trend of decreasing Si size in ADC12 alloy with the 0.7 wt.% rGO addition. The reduced Si size with graphene addition was similarly attributed to: (i) grain refinement of primary Al (Figure 4.17), (ii) heterogeneous nucleation of Si on graphene, and (iii) impeding the growth of Si, as shown in Figure 4.21 and Figure 4.22. It can be seen from Figure 4.21(a) that graphene is present on the surface of Si, which suggests heterogeneous nucleation of Si by graphene. 89 Figure 4.19. Optical microstructure of cast samples: (a) B319, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene and (f) B319-0.7 wt.% graphene. Grain refinement of primary Al reduces the interdendritic melt volume and limits Si nucleation regions, resulting in fine Si particles during eutectic transformation. The EDS elemental mapping of the B319-0.1 wt.% graphene composite reveals the presence of graphene in the immediate vicinity of the Si phases (Figure 4.21(b)). SEM and EDS area maps of B319-0.3 wt.% graphene composite also confirm the presence of graphene at the Si growth front, as shown in Figure 4.22(a,b). The presence of graphene at the Si growth front changes the Si crystal's growth direction, eventually changing the morphology of the Si phase to a more fibrous one. The presence of graphene can also inhibit the growth of Si, which is beneficial for mechanical and tribological properties [36,261]. 90 Figure 4.20. Effect of graphene on the morphological characteristics of Si: (a) Length, (b) Width, and (c) Aspect ratio. The presence of graphene can influence the thermal field (caused by a thermal conductivity gradient between the Al matrix and graphene) and the solute field (by obstructing diffusion of solute Si atoms) locally, both of which contribute to Si refinement. It is known that graphene (4840-5300 W/mK [110]) has a higher thermal conductivity than the B319 Al matrix (109 W/mK [262]). The higher thermal conductivity of graphene can quickly transfer and dissipate heat toward the casting–mold interface [214]. Consequently, the temperature of graphene-rich regions is lower than that of the surrounding matrix during solidification, contributing to grain refinement in composites fabricated in the present work. Chak and Chattopadhyay [42] also observed a ~47% reduction in the grain size of AA7075 Al alloy with the addition of 0.5 wt.% graphene. 91 Figure 4.21. (a) SEM image and (b) EDS elemental mapping of B319-0.1 wt.% graphene composite, showing the presence of graphene in the growth front and on the surface of Si. Figure 4.22. SEM image and corresponding EDS elemental maps of B319-0.3 wt.% graphene composite. 4.5 Hardness, Electrical, and Thermal Conductivity of B319-Graphene Composites The hardness of the B319-graphene composites is shown in Figure 4.23(a). The hardness increased from 77.36 to 92.78 HV with 0.1 wt.% addition of graphene, resulting in ⁓20% improvement over the base B319 Al alloy. The improved hardness in composites indicates the enhanced plastic resistance improvement, likely due to the microstructural refinement (dendrites, grains, and eutectic Si) caused by the graphene addition (Figure 4.17 and Figure 4.20). However, further increase of graphene concentration to 0.2 and 0.3 wt.% showed no significant improvement in 92 hardness, and at 0.5 and 0.7 wt.% graphene addition levels, the hardness decreased. The decrease in hardness at higher concentrations of graphene was attributed to graphene's agglomeration and increased porosity in the as-cast material (Figure 4.18). Figure 4.23. Effect of graphene content on the (a) hardness, (b) % IACS, measured by eddy current electrical conductivity meter, (c) electrical conductivity, measured by four-point probe method, and (c) thermal conductivity of B319 Al alloy. Li et al. [263] fabricated Al-0.2 wt.% graphene composites by casting route followed by rolling, and the resultant composite exhibited a ~16% improvement in hardness over the unreinforced matrix. This enhancement was mainly attributed to the load transfer strengthening mechanism. In a separate study, Askarnia et al. [264] developed Al-0.5, 1, and 2 wt.% GO composites via 93 microwave sintering and observed a ~44% increase in the hardness of Al with 0.5 wt.% GO addition. However, beyond 0.5 wt.% graphene, the hardness of the alloy deteriorated due to agglomeration and non-uniform distribution of GO. The effect of graphene content on the electrical and thermal conductivity of B319 Al alloy is presented in Figure 4.23(b and c) and Figure 4.23(d), respectively. The %IACS of the B319 Al alloy decreased with graphene addition up to 0.3 wt.% and then slightly increased with 0.5 and 0.7 wt.% graphene addition. However, all the composites showed lower %IACS than the base B319 Al alloy. A similar trend was also observed in the electrical conductivity of the B319-graphene composites from 25 °C to 300 °C evaluated by the four-point probe method. The B319-graphene composites exhibited lower thermal conductivity compared to the corresponding unreinforced alloy. The low electrical and thermal conductivity of the B319-graphene composites can be attributed to enhanced electron and phonon scattering caused by the formation of structural imperfections such as matrix-reinforcement interfaces, increased porosity, and an increase in the grain boundary area. The deterioration in conductivity of the Al matrix with graphene was also reported by other researchers [218,219]. 4.6 Chapter Summary This chapter discussed the fabrication of B319 Al alloy-graphene composites by a novel distribution approach and the effect of graphene addition on the microstructure, hardness, electrical conductivity, and thermal conductivity of the B319 Al alloy. The distribution approach used in this study involves two main steps: first, the fabrication of an optimized composite master alloy of F357 Al-5 wt.% of graphene using a flake-PM process. Secondly, the composite master alloy was introduced into the B319 Al alloy melt. As the master alloy melted within the liquid B319 Al alloy, graphene dispersed throughout the B319 melt. The molten B319 Al alloy, now containing dispersed graphene, was then stirred to ensure uniform distribution before being poured into graphite molds. The key points of this chapter are summarized below. The graphene in the present study had a carbon content of approximately 96%, measured by proximate analysis using DTA-TGA. Raman analysis of graphene revealed that the graphene consisted of 5-10 layers. The BM of the F357 powder for 1 h resulted in a morphological transformation of the powder from spherical to flake. The F357 powder with varying amounts of 94 graphene, including 2.5 wt.%, 5 wt.%, and 10 wt.%, was prepared by BM. The milled powder mixture was consolidated by SPS, and all composite master alloys have shown similar relative densities (~98-99%). A uniform distribution of graphene was observed in the F357-2.5 wt.% and F357-5 wt.% graphene composite master alloys. However, the F357-10 wt.% graphene composite master alloy exhibited severe agglomeration of graphene. Additionally, the F357-10 wt.% graphene composite master alloy did not completely melt in the molten alloy of B319. As a result, F357-5 wt.% graphene composite master alloy was selected to prepare B319-graphene composites. SEM-EDS and EPMA-WDS analysis of the as-cast microstructure of B319 Al alloy revealed the presence of Al, Si, Al2Cu, and Al15(Fe, Mn)3Si2 phases. The addition of graphene decreased the grain size of the B319 Al alloy, which was attributed to the additional nucleation sites and the alteration of the thermal field induced by the presence of graphene. Moreover, adding graphene refined the Si phase and changed its morphology to a more fibrous structure. The microstructural characteristics, including SDAS, Si aspect ratio, grain size, as well as relative density and hardness of B319-graphene composites, are summarized in Table 4.4. Table 4.4. Microstructural characteristics, relative density, and hardness of B319-graphene composites. Sample SDAS (µm) Si aspect ratio Grain size (µm) Relative density (%) Hardness (HV) B319 18 (±3) 11 (±3.2) 1660 (±400) 99.82 77 (±3) B319-0.1wt.% graphene 13 (±3) 4.5 (±2.4) 1530 (±450) 97.95 93 (±4) B319-0.2wt.% graphene 14 (±3) 4.4 (±2.9) 1390 (±370) 97.63 91 (±5) B319-0.3wt.% graphene 14 (±2) 4.4 (±2.4) 1350 (±300) 97.52 91 (±6) B319-0.5wt.% graphene 15 (±2) 4.5 (±2.5) 1210 (±240) 97.30 81 (±5) B319-0.7wt.% graphene 15 (±2) 4.5 (±2.5) 910 (±220) 97.00 80 (±6) The hardness of the B319 Al alloy increased by ~20% with the addition of 0.1 wt.% graphene, but it deteriorated with 0.5 and 0.7 wt.% graphene. The relative density of the B319 Al alloy decreased with graphene addition, indicating increased porosity. The decrease in hardness at higher concentrations of graphene was primarily due to graphene agglomeration and increased porosity. 95 Additionally, the B319-graphene composites exhibited inferior electrical and thermal conductivity compared to the base alloy, likely due to the formation of structural imperfections. 96 Chapter 5: Effect of Graphene on the Friction and Wear Behavior of B319-Graphene Composites This chapter focuses on the study of the effect of graphene on the reciprocating dry sliding friction and wear behavior of B319 Al alloy under 5 N and 10 N loads. Section 5.1 presents the frictional response of the B319-graphene composites. The wear response of the B319-graphene composites is discussed in Section 5.2, where wear track, wear debris, worn surface, subsurface microstructure, and hardness of these composites are characterized to understand the role of graphene in enhancing the wear resistance of B319 Al alloy and to identify the wear mechanisms in these novel composites. Section 5.3 provides the summary of this chapter. 5.1 Frictional Behavior of B319-Graphene Composites The evolution of linear reciprocating CoF with respect to sliding time for the B319-graphene composites tested at 5 N and 10 N is shown in Figure 5.1 and Figure 5.2, respectively. Smoothed CoF curves are depicted by the yellow line in each plot. The smoothing of CoF curves was carried out using the OriginPro software via the adjacent-averaging method (500 points). The friction curves of all samples demonstrated two distinct regimes called “I” running-in and “II” steady-state, regardless of the applied load. At the beginning of the wear test, a running-in stage with increased CoF was observed for all the tested samples. Generally, the virgin surfaces of the sample and the steel ball did not perfectly conform during the initial contact cycles, and the contact was primarily at the peaks of surface asperities. Consequently, the contact area remained relatively small, leading to a high CoF. As running-in progressed, the surface asperities underwent elastic-plastic deformation (asperity truncation) [265], with a concomitant increase in the contact area and strain hardening of the contact surface, resulting in a reduction in CoF. Removing surface asperities in the running-in period resulted in a direct contact between the B319 Al alloy and its counterbody with freshly exposed surfaces, thereby increasing the adhesive and abrasive wear modes. Concurrently, the plowing of surface asperities, generation of wear debris, and increased surface roughness lead to the observed fluctuations in CoF during the initial stages of the wear test. Thus, the fluctuations in CoF during the running-in regime can be attributed to the microtopography of the surface features, changes in the strain hardening state, initial elastic-plastic deformation of the surface asperities, and small contact areas between the surface and the counterbody [266]. 97 Figure 5.1. Evolution of CoF as a function of sliding time for the B319 Al alloy composites reinforced with varying amounts of graphene tested at 5 N: (a) 0 wt.%, (b) 0.1 wt.%, (c) 0.2 wt.%, (d) 0.3 wt.%, (e) 0.5 wt.%, and (f) 0.7 wt.%. “I” and “II” zones in the figures represent the running-in and steady-state regimes, respectively. 98 Figure 5.2. Evolution of CoF as a function of sliding time for the B319 Al alloy composites reinforced with varying amounts of graphene tested at 10 N: (a) 0 wt.%, (b) 0.1 wt.%, (c) 0.2 wt.%, (d) 0.3 wt.%, (e) 0.5 wt.%, and (f) 0.7 wt.%. “I” and “II” zones in the figures represent the running-in and steady-state regimes, respectively. 99 After a period of sliding, the aforementioned effects stabilized, leading to the attainment of a steady-state regime for the CoF (labeled as “Stage II”). The steady-state of wear (and CoF) was associated with the formation of a mechanically mixed layer (MML) and/or a tribolayer. The addition of graphene showed no significant change in the amount of time to reach a steady-state, since all the composites reached a steady state after ∼10 min or 60 m of sliding at a 5 N load, as shown in Figure 5.1(a-f). However, at a 10 N load, composites with higher concentrations (0.5 and 0.7 wt.%) of graphene reached steady-state at ∼15 min or 90 m of sliding distance (Figure 5.2(e,f)). The extended running-in regime observed in composites with 0.5 and 0.7 wt.% of graphene can be related to the higher surface roughness caused by increased porosity and agglomeration of graphene (Figure 4.18 and Table 5.1). Pores provide an inherent roughness that results in greater asperity-asperity contacts during sliding. Consequently, composites with a high porosity took longer to reach a steady-state. However, composites with a higher percentage of graphene tested at a lower load of 5 N did not take additional time to reach a steady-state. Table 5.1. Average surface roughness of polished samples as measured by a profilometer. Sample Surface roughness (Ra, µm) B319 0.04 (±0.01) B319-0.1wt.% graphene 0.05 (±0.01) B319-0.2wt.% graphene 0.06 (±0.01) B319-0.3wt.% graphene 0.06 (±0.01) B319-0.5wt.% graphene 0.19 (±0.16) B319-0.7wt.% graphene 0.21 (±0.16) At lower loads (5 N), the contact pressure between the contacting surfaces is reduced, resulting in less intimate contact or effective contact area between the mating surfaces and a lower frictional force. This can lead to a quicker stabilization of the frictional behavior and attaining steady-state conditions. Conversely, higher loads (10 N) can induce greater deformation of the material, causing pores in B319-0.3 wt.% and 0.7 wt.% graphene composites to collapse, thereby increasing the effective contact area [267]. This increase in effective contact area combined with complex topographical features, results in increased frictional forces, hindering the efficient transfer of forces between the sample and its counterbody, delaying the establishment of steady-state friction, 100 as seen in Figure 5.2(e,f). Additionally, the increased generation of wear debris at higher loads can trap the debris within the pores or voids, further hindering the establishment of steady-state friction by having the debris act as abrasive and disrupting the contact between the surfaces. The average CoF during the running-in regime of the B319 Al alloy tested at 5 N was 0.49. It decreased with graphene addition, showing a maximum reduction of 55% with 0.3 wt.% of graphene addition. However, composites with higher concentrations of graphene exhibited an increased running-in CoF (~0.46), returning to levels similar to that of the base alloy. In contrast, samples tested at a higher load of 10 N exhibited similar CoF values (ranging from 0.25-0.28) in the running-in regime, irrespective of the graphene content (Figure 5.2). This observation indicates that the graphene content did not significantly affect the friction behavior in the early stages of the wear test under a 10 N load. At both loads, the average CoF in the steady-state regime decreased with graphene addition up to 0.3 wt.% graphene and then increased at 0.5 wt.% and 0.7 wt.% graphene. Furthermore, CoF fluctuations around its mean value can be observed for all the samples in the steady-state regime regardless of the load (Figure 5.1 and Figure 5.2), which can be attributed to the stick-slip phenomenon caused by adhesion work [268]. The adhesion-induced stick-slip phenomenon is closely related to friction and the variation of static (Fs) and kinetic friction force (Fk) as a function of sliding time or distance. High adhesion work promotes the formation of strong junctions at the contact surfaces, which increase the tangential force. These conditions can lead to a period of no relative motion between the contact surfaces, called the “stick” phase; hence, the friction force reaches the magnitude of Fs. However, when the tangential force surpasses the threshold of Fs, it breaks the adhesive junctions, and the friction decreases to the magnitude of Fk, leading to a sudden increase in the sliding speed called “slip”. During the slip stage, the junctions are weak as there is no time to fully form adhesive junctions. At this point, the system decelerates to keep the sliding speed constant. This slowing down can lead to the formation of strong adhesive junctions again, and the cycle is repeated. Therefore, the tangential force undergoes periodic instability during sliding due to intermittent motion, causing severe fluctuations in CoF around the mean value (Figure 5.1 and Figure 5.2) [269,270]. Other researchers have also observed the stick-slip phenomenon in Al-Si alloys under dry sliding conditions [268,271]. 101 Figure 5.3. Average CoF (including running-in and steady-state regimes) of B319-graphene (0-0.7 wt.%) composites tested at 5N and 10N loads. The average CoF (including running-in and steady-state regimes) of the composites tested at 5 N and 10 N is presented in Figure 5.3. The B319-0.3 wt.% graphene composite exhibited the best anti-friction properties, with the average CoF decreasing from 0.42 to 0.20 at 5 N, representing a ⁓52% reduction compared to the unreinforced alloy. Similarly, the CoF decreased from 0.26 to 0.17 at 10 N load, indicating a ~42% reduction in CoF. Further addition of graphene to 0.5 wt.% and 0.7 wt.% increased the CoF to 0.40 and 0.41 at 5 N, 0.22, and 0.24 at 10 N, respectively, suggesting the best anti-friction performance at 0.3 wt.% graphene concentration. Even though the composites with 0.1 wt.% and 0.2 wt.% graphene concentrations exhibited lower CoF values compared to the base alloy, the values were still relatively higher than those of the 0.3 wt.% graphene composite. The insufficient amount of graphene at the contact surface can be the reason for the higher CoF at lower concentrations of graphene (0.1 and 0.2 wt.%). Khorshid et al. [272] also observed that the 1 wt.% graphene composite showed a lower CoF than the 0.1 wt.% graphene composite. The degradation of friction properties at higher concentrations of graphene can be 102 related to the lower hardness and higher porosity resulting from the agglomeration of graphene at the grain boundaries of the alloy. Similar observations were also reported by Hsieh et al. [273], where Al6061-0.25 wt.% and 0.5 wt.% graphene composites fabricated through high-frequency induction heat sintering followed by hot extrusion. Their findings revealed that the addition of 0.25 wt.% graphene resulted in the maximum reduction in the CoF of A6061 alloy, while further addition of graphene to 0.5 wt.% increased the CoF. This degradation in the friction response of A6061 alloy at 0.5 wt.% graphene was attributed to the agglomeration of graphene primarily along the grain boundaries. The improved friction performance of B319 Al alloy with graphene addition can be attributed to graphene’s self-lubricating ability, resulting from its unique structural and mechanical properties. It has been reported that graphene has a high in-plane strength and weak interlayer shear strength [31,274]. The interlayer bonding in graphene can easily break under the loads applied in wear tests, reducing the resistance to shear between layers and facilitating easier sliding [35]. Additionally, it was reported that the formation of rollers from the packets of graphene layers on the counter surface can act like roller bearings and reduce the CoF [275,276]. Thus, the formation of rollers and interlayer sliding in graphene during sliding can contribute to the lower CoF of B319-0.3 wt.% graphene composite. It is worth noting from Figure 5.3 that the overall magnitude of CoF (during running-in and steady-state regimes) of all the tested samples was comparatively lower at 10 N compared to 5 N. Wu et al. [201] also observed a decrease in CoF with increasing normal load in AlSi10Mg-graphene composites fabricated by SLM. This trend was attributed to the self-lubricating property of graphene, which facilitates easy interlayer sliding at a higher load, thereby reducing friction. Furthermore, the increase in load increased the contact area between the alloy at the counterbody, resulting in a more uniform distribution of contact force and a concomitant decrease in friction or resistance to sliding per unit area. Increasing the load from 5 N to 10 N was expected to increase the surface temperature of the test sample during wear testing due to frictional heating. This heat and elevated surface temperature would enhance the rate of oxide layer formation. Simultaneously, Si-rich phases were observed to fracture into fine particles, along with an increased generation of wear particles at a higher frictional load [277]. As a result, forming a thick MML consisting of aluminum oxide and fractured Si particles decreased the adhesion by effectively preventing direct 103 contact between the steel ball and B319 Al alloy, thus reducing overall CoF [278]. This is further discussed in Sections 5.2.1, 5.2.5.1, and 5.2.5.3.1, where wear track surface and subsurface analysis results are presented. 5.2 Wear Behavior of B319-Graphene Composites 5.2.1 Wear Track Analysis The SEM images of wear tracks for the B319-graphene (0–0.7 wt.%) composites tested under 5 N and 10 N loads are shown in Figure 5.4 and Figure 5.5, respectively. The wear tracks of all the tested samples showed an uneven wear pattern, suggesting the “stick-slip” phenomena described earlier. Surface profile maps of worn surfaces for the B319-graphene composites tested under 5 N and 10 N loads, scanned with a laser microscope, are shown in Figure 5.6(a-f) and Figure 5.7(a-f), respectively, with the color scale indicating the depth of the wear track. The surface profile map of the B319 Al alloy shows two to three peak regions at the surface of the wear track under both loads, and the addition of graphene showed no significant changes in terms of peak regions. At a 5 N load, the maximum wear track depth slightly decreased from 204 (±5) µm to 201 (±3) µm with 0.3 wt.% graphene addition, as shown in Table 5.2. Similarly, it can be seen from Figure 5.5 and Figure 5.7 that the wear track becomes significantly narrower and shallower with the addition of graphene up to 0.3 wt.%. The average maximum wear track depth of the B319 Al alloy was 385 (± 8) µm, which decreased to 222 (± 3) µm when 0.3 wt.% of graphene was added, as shown in Table 5.2. However, at a higher graphene concentration (0.5 and 0.7 wt.%), the maximum depth increased regardless of the load, indicating severe wear. Depth variation is often attributed to the ejection and accumulation of wear debris during sliding at the surface of the wear track. The wear track width of all tested samples was measured using Image J software by making at least 50 measurements on SEM images, and the average value was reported with a standard deviation in Table 5.2. While the wear track widths of all the tested samples overlap within the range of standard deviation, a trend of decreased average wear track width is evident with the addition of graphene. Similar to the wear track depth, the average wear track width also decreased with the addition of graphene and had the lowest width in the 0.3 wt.% graphene composite. The average wear track width reduced from 1 mm to 0.92 mm and 1.37 to 1.02 mm under 5 N and 10 N loads, respectively. However, adding graphene beyond 0.3 wt.% increased the wear track size, 104 indicating reduced wear resistance. The increase in wear track size at higher graphene contents can be attributed to increased porosity and reduced hardness caused by the agglomeration of graphene. Figure 5.4. Secondary electron SEM images of the wear track: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 5 N load. Figure 5.5. Secondary electron SEM images of the wear track: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 10 N load. 105 Figure 5.6. Surface profile maps of wear tracks captured by laser confocal microscope: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 5 N load. The color scale indicates wear track depth. Figure 5.7. Surface profile maps of wear tracks captured by laser confocal microscope: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 10 N load. The color scale indicates wear track depth. 106 Table 5.2. Effect of graphene on the average wear track depth and maximum depth of B319-graphene composites tested under 5 N and 10 N loads. Sample Wear track width (mm) Maximum depth (µm) 5 N 10 N 5 N 10 N B319 1.00 (±0.15) 1.37 (±0.26) 204 (±5) 385 (±8) B319-0.1wt.% graphene 0.97 (±0.16) 1.27 (±0.15) 219 (±9) 254 (±6) B319-0.2wt.% graphene 0.94 (±0.10) 1.10 (±0.13) 194 (±4) 260 (±2) B319-0.3wt.% graphene 0.92 (±0.14) 1.02 (±0.14) 201 (±3) 222 (±3) B319-0.5wt.% graphene 0.94 (±0.17) 1.16 (±0.19) 246 (±5) 305 (±5) B319-0.7wt.% graphene 1.03 (±0.16) 1.19 (±0.18) 237 (±7) 298 (±1) 5.2.2 Analysis of Wear Scar on Counterbody In the present study, the counterbody (100Cr6 ball) exhibited no noticeable wear, regardless of the load. This can primarily be attributed to the material transfer that occurred from the sample to the counterbody via adhesive wear and the high hardness (770 HV) of the ball. The plastic deformation of the Al alloy can also contribute to material transfer from the tested sample to the ball due to localized deformation and flow. In addition, the surface profile maps clearly showed the buildup or pileup of material at the sides and extremities of a wear track, further confirming an adhesive wear mechanism involving plastic deformation in B319-graphene composites, as seen in Figure 5.6 and Figure 5.7. The surface film that evolved on the counterbody acted as a form of lubrication or protective barrier between the virgin ball surface and the contact region, thereby reducing direct contact between the sample and the counterbody, leading to minimal ball wear [279]. The SEM images of the wear scar on the ball tested under a 5 N are presented in Figure 5.8(a-f). These images clearly show the adhesive patches on the wear scars of the balls (Figure 5.8(a)). EDS area maps of the counterbody wear scar of the B319-0.3 wt.% graphene composites tested at a 5 N load, as marked in Figure 5.8(d), reveal that the adhered patches mainly composed of Al, O, Si, Cu, and Mn elements, confirming material transfer from the Al alloy to the ball through adhesive wear (Figure 5.9(a,b)). The EDS area maps also show the presence of carbon in the adhesive patches on the ball, indicating 107 graphene transfer from the B319-0.3 wt.% composite. The presence of other elements, including Cr and Fe, is associated with the steel ball (Figure 5.9(b)). Figure 5.8. SEM images of wear scar on the counterbody (100Cr6 balls) after wear testing of (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene composites tested under 5 N load. 108 Figure 5.9. (a) SEM micrograph of wear scar on steel ball slide against B319-0.3 wt.% graphene composite under 5 N, marked region (A) in Figure 5.8(d) and (b) the corresponding EDS area map. Figure 5.10. SEM images of wear scar on the counterbody after wear testing of (a) B319-0 wt.% graphene, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (d1) high magnification image of a region (A) marked in Fig. (d), (d2) carbon map of Fig. (d1), (e) B319-0.5 wt.% graphene, and (f) B319-0.7 wt.% graphene tested under 10 N load. 109 Figure 5.11. Average wear scar area on the counterbody (100Cr6 balls) after wear testing of B319-graphene composites. Furthermore, the SEM images of wear scars on the balls under a 10 N load are shown in Figure 5.10(a-f). Similar to the EDS maps of wear scars observed at a 5 N load, the EDS maps of wear scars tested under a 10 N load also clearly show material transfer from the composite surface to the ball, as shown in Figure 5.10. Additionally, the high-magnification image of the region (A) marked with a yellow square box (Figure 5.10(d1)) and its corresponding carbon EDS map (Figure 5.10(d2)) indicate the presence of graphene on the adhered patches. Observations from Figure 5.8 and Figure 5.10 show varying sizes of wear scars on the counterbody as a function of graphene content in the B319 Al alloy, tested under loads of 5 N and 10 N, respectively. The B319-0.3 wt.% graphene composite counterbody exhibited the lowest wear scar area at both loads, as shown in Figure 5.11. At 5 N and 10 N, the B319-0.3wt.% graphene composite exhibited a wear scar area 110 of 1.69 mm2 and 2.55 mm2, showing ~30% and ~15% reduction over the wear scar area of the base alloy counterbody, respectively. These observations agree well with the results of the wear track morphology characterization. Consequently, these findings suggest that the wear of the B319 Al alloy decreased by adding 0.3 wt.% graphene, which was reflected in the size of wear tracks and wear scars on the counterbody. 5.2.3 Wear Properties The mass loss of B319-graphene composites after testing at 5 N and 10 N loads is summarized in Figure 5.12. The base alloy experienced a mass loss of 4.9 (±0.4) mg at 5 N, which escalated to 9.8 (±0.4) mg at 10 N, representing a 100% increase. Similarly, the B319-composites showed increased mass loss with the load rising from 5 N to 10 N, showing increments in the range of approximately 55% to 120%. This increase is attributed to higher loads exerting greater contact or frictional forces between the surfaces, leading to more severe mechanical interactions and greater material removal through various wear mechanisms. However, the mass loss of B319 Al alloy decreased with graphene addition up to 0.3 wt.% content at both loads. Beyond 0.3 wt.% graphene addition, the mass loss increased, but all the composites showed lower or similar mass losses compared to the unreinforced alloy. In this study, two methods (weight loss measurement and 3D laser microscope analysis), were employed to assess the wear volume and specific wear rate of the B319-graphene composites, with the aim of understanding their wear behavior. Wear properties (wear volume and specific wear rate) measured using both methods exhibited a similar trend, as seen in Figure 5.13 and Figure 5.14. Several factors can introduce errors in the weight loss method, such as oxidation of the worn surface, accuracy of the weighing scale (especially at small mass losses), moisture absorption, material transfer onto the worn surface from the steel ball, etc. In contrast, a laser microscope allows for accurate wear volume determination due to selective imaging, non-destructive testing, and high resolution [243]. As observed in CoF (Figure 5.3) and mass loss (Figure 5.12) measurements, the wear volume and specific wear rate of B319 Al alloy decreased with the addition of graphene up to 0.3 wt.% and then increased with further addition at both 5 N (Figure 5.13(a) and Figure 5.14(a)) and 10 N loads (Figure 5.13 (b) and Figure 5.14(b)). The average wear volume and specific wear rate of the 0.3 wt.% graphene composite at 5 N showed a reduction of 111 ⁓39% (measured by weight loss) and ⁓37% (measured by the laser microscope) compared to the base B319 Al alloy. Figure 5.12. Effect of graphene on the mass loss of B319 Al alloy-graphene composites tested under 5 N and 10 N loads. The average wear volume and specific wear rate of the B319-graphene composites tested under 10 N are presented in Figure 5.13(b) and Figure 5.14(b), respectively. The wear volume of the B319 Al alloy at 10 N was estimated as 3.8 ± 0.02 mm3, which decreased to 1.8 ± 0.12 mm3 with 0.3 wt.% graphene addition. Similarly, the specific wear rate of the B319 Al alloy decreased from 12.7 ± 0.08 mm3/N.m×10-4 to 6.1 ± 0.40 mm3/N.m×10-4 at 0.3 wt.% graphene addition, corresponding to a ∼52% improvement in wear resistance. Xia et al. [280] fabricated Al7075-0.3 wt.% graphene composites through BM and SPS. The incorporation of graphene reduced the wear rate of Al7075 from 3.4 mm3/N.m×10-3 to 2.7 mm3/N.m×10-3, indicating a 20% reduction. The wear volumes of the B319-graphene composites obtained at 10 N in the present study are ∼2x of 112 the wear volumes observed at 5 N, indicating that the specific wear volumes are similar at both loads and comparable wear mechanisms exist at both loads. The increased wear volume with an increase in load can be associated with increased Hertzian contact radius and contact pressure, which are calculated using Equations (5.1), (5.2), and (5.3) [281]. Both the contact radius and maximum contact pressure increased with the increasing load, as shown in Table 5.3. Higher contact pressure intensifies the wear mechanisms, such as adhesive, abrasive, oxidation, surface fatigue, and deformation-related wear, leading to greater wear. 𝑎 = (3𝐹𝑛𝑅4𝐸∗)13 (5.1) 1𝐸∗=1−𝜗12𝐸1+1−𝜗22𝐸2 (5.2) 𝑝𝑜 =3𝐹𝑛2𝜋𝑎2 (5.3) Where: 𝑎 is contact radius, 𝐹𝑛 is normal load (5, 10 N), 𝑅 is 100Cr6 ball radius (6 mm), 𝐸∗ is reduced modulus, 𝐸1 is elastic modulus of the B319 Al alloy (74 GPa), 𝐸2 elastic modulus of the 100Cr6 ball (210 GPa), 𝜗1 is Poisson's ratio of the B319 Al alloy (0.33), 𝜗2 is the Poisson's ratio of the 100Cr6 ball (0.30), and 𝑝𝑜 is the maximum Hertzian contact pressure. Figure 5.13. Effect of graphene on the wear volume loss of B319 Al alloy tested under (a) 5 N and (b) 10 N loads. The wear volume loss was calculated using the weight loss method and a laser microscope. 113 Figure 5.14. Effect of graphene on the specific wear rate of B319 Al alloy tested under (a) 5 N and (b) 10 N loads. The specific wear rate was calculated using the wear volume loss determined by the weight loss method and 3D laser microscope. Table 5.3. Effect of normal load on the Hertzian contact radius and maximum contact pressure. Normal load (N) Contact radius (a, mm) Contact pressure (𝒑𝒐, MPa) 5 0.0569 737 10 0.0716 929 The improvement in wear resistance of B319 Al alloy with 0.3 wt.% graphene addition was attributed to the higher hardness caused by microstructural refinement (grain size reduction (Figure 4.17) and Si modification (Figure 4.20)) and the formation of a graphene-rich lubricant film at the contact surface. Many researchers used hardness to correlate the tribological properties of Al-graphene composites [239,282], and it was demonstrated that Si refinement improves Al alloys' wear resistance [271,283]. The fine Si particles reduce stress concentration at the Al/Si interface and are considered less prone to cracking or decohesion during the sliding contact, thus reducing abrasive wear [284]. Also, graphene enhances wear resistance by forming a tibolayer that prevents direct contact between the sliding surfaces. Graphene exhibits self-lubricating behavior through interlayer sliding and roller formation during sliding. The strong in-plane covalent bonds and weak van der Waals interlayer forces in graphene enable the layers to slide under shear forces generated 114 during sliding. Moreover, the interlayer bonds eventually break and re-orient at the contact surfaces, generating roller sheets [35,275,276]. These two effects (interlayer sliding and roller formation) caused by graphene can reduce friction and wear. The increased CoF and wear in 0.5 wt.% and 0.7 wt.% graphene composites can be related to lower hardness, high porosity, and agglomeration of graphene. 5.2.4 Dissipated Frictional Energy The dissipated frictional energy is an important aspect of tribology and can give valuable insight into the wear behavior and damage mechanisms [285,286]. A plot of overall wear volume measured by a laser microscope and dissipated energy against the graphene content for B319-graphene composites tested at 5 N and 10 N is presented in Figure 5.15(a) and Figure 5.15(b), respectively. At both loads, the wear volume loss and dissipated frictional energy show a similar pattern of change, indicating a strong dependence of wear volume on dissipated energy and CoF (Figure 5.15). The dissipated energy at the contact surfaces either gets stored in the material or may transform into heat energy (frictional heating). The dissipated energy stored in the material enables plastic deformation, delamination, wear debris generation or ejection, and direct material loss. At the same time, the heat energy increases the temperature at the contact surface and can induce tribochemical reactions (i.e., oxidative wear) [287–289]. Figure 5.15. Average wear volume and dissipated frictional energy of B319 Al alloy-graphene composites tested at (a) 5 N and (b) 10 N loads. 115 Several researchers showed linear dependence of wear volume on frictional energy, but this is only expected when the wear mechanisms remain the same for all the tested samples [242,287,290,291]. The 0.3 wt.% graphene composite dissipated the lowest frictional energy compared to the unreinforced alloy at both loads, which was attributed to the enhanced heat dissipation capacity of the MML, thereby minimizing heat accumulation in MML and reducing surface damage [292]. The base alloy dissipated 630 J and 780 J of energy at 5 N and 10 N, respectively, which decreased to 300 J and 510 J with 0.3 wt.% graphene addition. Hence, the higher wear resistance of B319-0.3 wt.% graphene composite is reconfirmed. Maurya et al. [236] fabricated AA6061-graphene composites by FSP and observed a ~45% decrease in dissipation energy with graphene addition over unreinforced alloy. All the samples tested at 10 N (Figure 5.15(b)) exhibited a higher dissipated energy than the samples tested at 5 N, as seen in Figure 5.15(a). Frictional energy dissipation is associated with generating heat at the contact interface. Under higher loads, the sample undergoes greater deformation. This deformation can result in microstructural changes such as plastic flow, surface roughening, and even material transfer between the contacting surfaces, all of which contribute to increased frictional energy generation. Additionally, the higher levels of frictional energy led to increased heat generation, which can further exacerbate material degradation, surface roughening, and wear, contributing to even higher frictional energy dissipation. Quantifying the energy consumed by wear mechanisms can be challenging due to the complex nature of wear processes [285]. However, it was reported that 70% of energy is dissipated by heat generation [293]. The dominant energy dissipation pathways can be identified by analyzing the surface and subsurface of wear tracks, as discussed in the following Section. 5.2.5 Wear Mechanisms 5.2.5.1 Worn Surface Analysis The wear tracks of all alloy samples were examined using SEM-EDS, and the track of B319 Al alloy at 5 N is shown in Figure 5.16. The deepest region in the B319 Al alloy’s wear track, marked in Figure 5.16(a) and Figure 5.16(b), revealed a delaminated oxide-rich layer (Al2O3) or MML, suggesting wear by tribo-oxidation accompanied by delamination (Figure 5.16(c) and Figure 5.16(d)). Moreover, the EDS area map of the worn surface marked in Figure 5.16(b) reveals the 116 presence of mechanically mixed Al, Si, Cu, Fe, and Mn together with alumina, as shown in Figure 5.16(e). Both mechanical mixing and tribo-oxidation are considered to be sub-mechanisms of tribochemical reactions [294]. Rigney et al. [295–297] observed the “mechanical mixing” mechanism in ductile materials under unlubricated sliding conditions. In this mechanism, initially, the plastic deformation occurs at the local contacts and accumulation of plastic strains occur with continued sliding. Subsequently, the formation of deformation substructure occurs heterogeneously eventually leading to shear instability at the near-surface material. This, in turn, leads to the local transfer of material across the interface, mainly through the adhesion mechanism. This initiates a complex mixing process similar to what occurs in BM or mechanical alloying [279,298,299], resulting in the formation of patches or layers of mechanically mixed material (MMM). Simultaneously, the formation and fracturing of oxides will occur with further sliding. The generated oxides are then mechanically mixed with unoxidized base material [300]. Therefore, the MMM consists of components from the sample, counterbody, and/or the environment, resulting in a possible composition unique from that of the baseline material. Material transfer can occur from the sample to the counterbody or vice versa, or transfer can occur in both directions. In this study, material transfer mainly occurred from the sample to the counterbody. Wear debris is typically generated from the MMM, rather than directly from the base material. The formation of MML can be confirmed by analyzing the subsurface microstructure of the wear track, which will be discussed in Section 5.2.5.3. Oxide evolution during the dry sliding of B319 Al alloy was enabled by increasing temperature (due to frictional heating) at contact surfaces. While the bulk surface temperature may not reach very high levels during the wear test, the contact temperature may significantly increase, typically ranging from 180-400°C, depending on the applied loads, as estimated by Ashby’s model [301,302]. Consequently, these higher contact temperatures create favorable conditions for oxidation reactions to take place during wear testing. In oxidative wear, oxidation first occurs at single spots of the contacting plateaus, eventually developing into oxide islands or layers with further sliding. Simultaneously, as the oxide-rich layer forms, the compacted wear debris also undergoes oxidation due to frictional heating. Once these oxide layers grow into a critical thickness (1–3 µm) or the accumulated frictional energy in the MML attains a critical value [303,304], the 117 B319 Al alloy below the MML or oxide layer softens. The mismatch in CTE between the oxide layer and base alloy can generate thermal stresses at the interface. In this state, the softened Al can no longer support the MML, and the wear mechanism changes from tribo-oxidation to delamination. The critical thickness of the oxide layer depends on factors such as the mechanical properties and composition of the oxide layer itself and the adhesive strength between the base material and the oxide layer [305]. As a result, the MML dissipates energy through its plastic delamination, leading to wear debris generation or material transfer to the counterbody under cyclic stresses and frictional heat and subsequently exposing a fresh surface of the base alloy to oxidation [304]. The freshly exposed surface will have higher effusivity and heat dissipation capacity compared to the oxide layer and will absorb a greater amount of localized heat. Consequently, this heat will be retained within the contacting layers for an extended period, again resulting in an oxidative reaction. Therefore, oxidation wear in B319 Al alloy occurs through the following sequence: oxidation-fracturing-reoxidation. This cyclic process of oxidation wear continues with sliding in B319 Al alloy, leading to the gradual loss of the material. The formation of an oxide layer during the sliding wear occurs at the interface between the oxide and gas. It is well known that oxygen diffusion within oxide layers is slower compared to the diffusion of metal atoms (Al) within oxide layers. Therefore, the formation of an oxide layer during the wear test primarily occurs through the outward diffusion of metal atoms from the B319 Al alloy through the oxide layers. This suggests that the temperature or the available thermal energy at the contact surface, along with the duration of energy availability, governs the distance that metal atoms can travel or the growth rate of the oxide layer [306]. Thus, the thermal conductivity of the tested specimen is crucial, as it determines the temperature at the contact surfaces. Other researchers have also observed oxidative wear in Al-Si alloys [16,307–309]. In addition, the worn surface of B319 Al alloy had parallel and deep grooves along the reciprocating direction, characteristic of an abrasive wear mechanism caused by wear debris on the contact surface (Figure 5.17). The worn surface further shows evidence of abrasive wear sub-mechanisms, including micro-plowing and micro-cutting [294,310]. The wear debris of fractured oxide and the pull-out of exfoliated Si particles from the B319 Al alloy matrix trapped on mating surfaces can induce three-body abrasive wear, thereby accelerating the wear process. The worn 118 surface morphologies of the B319-0.3 wt.% graphene composite tested under a 5 N load are shown in Figure 5.18. Like the B319 Al alloy, the worn surface of the B319-0.3 wt.% graphene composite shows a broken oxide layer corresponding to the deepest wear track region, indicating tribo-oxidation wear mode (Figure 5.18(a-d)). Moreover, the EDS area mapping of the worn surface marked in Figure 5.18(c) reveals the presence of carbon-rich areas (graphene) alongside the elements observed on the worn surface of the base alloy Figure 5.18(e). Figure 5.16. Worn surface morphologies of B319 Al alloy tested at 5 N load: (a) SEM image, (b) laser microscope image, (c) SEM micrograph of the region marked in Fig. (a) and (b), (d) SEM micrograph of the region marked in Fig. (c) and corresponding EDS area map (e). 119 Figure 5.17. SEM micrograph of the worn surface of B319 Al alloy tested under a 5 N load, showing the different wear features. However, the presence of graphene appears to have increased the stability of the MML by (i) enhancing the load-bearing capacity of the MML and (ii) dissipating the frictional heat away from the contact surface. Both these effects can decrease the delamination rate of the MML, and the MML can further act as a tribolayer, thus improving the friction and wear resistance of the composite. The worn surface of the B319-0.3 wt.% graphene composite had smooth surface regions lined with graphene (Figure 5.19(a)), contributing to enhanced anti-friction and wear resistance. EDS mapping confirmed the presence of graphene on the worn surface (Figure 5.19(a-c)). Similar to the worn surface of the B319 Al alloy, the B319-0.3 wt.% graphene composite's worn surface exhibited adhesive traces (Figure 5.19(a)), micro-cutting, and plowing (Figure 5.18(b)), indicating adhesive and abrasive wear mechanisms. 120 Figure 5.18. Worn surface morphologies of B319-0.3 wt.% graphene composite tested at 5 N: (a) SEM image, (b) laser microscope image, (c) SEM micrograph of the region marked in Fig. (a) and (b), (d) SEM micrograph of the region marked in Fig. (c) and corresponding EDS area map (e). 121 Figure 5.19. (a) SEM micrograph of B319-0.3 wt.% graphene composite tested under a 5 N load, showing the different wear features; (b) SEM image of graphene marked in Figure (a) and corresponding EDS area map of carbon (c). The SEM images of the worn surface of B319-graphene composites tested under a 10 N load are shown in Figure 5.20(a-f). The area elemental maps (Figure 5.20(g)) of the base alloy’s worn surface (Figure 5.20(a)) showed the presence of oxygen-rich regions along with Al, fractured Si, and Cu, suggesting MML formation and oxidative wear. By combining the SEM image of the base alloy’s worn surface (Figure 5.20(a)) and corresponding area elemental maps (Figure 5.20(g)) with the surface profile maps of the wear tracks (Figure 5.7(a)), it is evident that the deepest regions in surface profile maps correspond to the delaminated areas of the oxide layer or MML. Furthermore, the worn surface of B319-0.1 wt.% and 0.2 wt.% graphene composites are shown in Figure 5.21(a) and Figure 5.22(a), respectively. The EDS maps of B319-0.1 wt.% graphene (Figure 5.21(b)) and B319-0.2 wt.% graphene (Figure 5.22(b)) composites also exhibited features of oxidative wear along with carbon-rich regions, confirming the presence of graphene. Although the oxidation wear was observed at both loads, it is expected that the degree of oxidation wear will be higher at the 122 10 N load compared to the 5 N load. The elevated temperatures resulting from the frictional heating at 10 N load can accelerate oxidation reactions, leading to higher rates of oxidation wear [311,312]. The delaminated regions of oxides can be seen on the base alloy's worn surface, and the relative area of delaminated regions decreased with the addition of graphene up to 0.3 wt.% (Figure 5.20(b-d)). Maurya et al. [236] also observed a reduction in oxidation wear of the Al6061 Al alloy under the dry sliding wear conditions with the incorporation of graphene. This reduction in oxidation wear was attributed to a reduction in frictional heating facilitated by the lubricating properties of graphene. Additionally, the presence of high-conductivity graphene at the contact surface can effectively dissipate the heat generated during sliding, thereby reducing localized temperature and consequently slowing down oxidation reactions. The worn surface of all the tested specimens further showed evidence of microploughing and microcutting along the sliding direction, irrespective of the load applied. These features are associated with an abrasive wear mechanism. The debris generated during sliding, which includes oxides, fractured Si particles, or work-hardened wear fragments, was expected to have a higher hardness than the B319 Al alloy. Consequently, these particles could freely slide or roll along the mating surface of the sample during sliding contact, contributing to three-body wear [310]. Lozano et al. [313] investigated the tribological response of a cast Al-Si-Cu alloy under sliding wear conditions, observing that the wear predominantly occurred via an abrasive mechanism. Furthermore, it was reported that the presence of small and uniformly dispersed Si particles reduced three-body wear, thereby improving the overall wear response of the alloy. In another study, Xu et al. [314] studied the effect of heat treatment and modification on the abrasive wear behavior of Al-Si alloy. The results showed that the modified and heat-treated alloy exhibited a superior anti-wear response over the non-treated alloy. The inferior wear resistance of the non-treated alloy was attributed to the presence of needle-like Si particles. Furthermore, the abrasive particles could penetrate deeper into the non-treated sample due to the larger free path of Si particles compared to treated alloys. In contrast, the small free path of Si particles in the treated alloy effectively supported the applied load, resulting in decreased wear. Similarly, Liu et al. [315] reported improved wear properties of Al-20 wt.% Si alloy with strontium (Sr) addition up to 0.06 wt.%, attributed to spheroidization of the Si particles. However, the wear resistance of the alloy degraded beyond 0.06 wt.% of Sr addition, which was attributed to an increase in Si size resulting 123 from over-modification. Therefore, the wear properties of the Al-Si alloys are dependent on the size, morphology, and distribution of Si particles. Figure 5.20. SEM images of worn surfaces of the B319-graphene composites tested under 10 N: (a) B319 Al alloy, (b) B319-0.1 wt.% graphene, (c) B319-0.2 wt.% graphene, (d) B319-0.3 wt.% graphene, (e) B319-0.5 wt.% graphene, (f) B319-0.7 wt.% graphene and (g) EDS area elemental mapping of B319 Al alloy (a). 124 Figure 5.21. (a) SEM image of the worn surface of the B319-0.1 wt.% graphene composite tested under a 10 N load and its corresponding EDS area maps (b). The bright white particles on the worn surfaces (marked in Figure 5.20(a)) correspond to debris from the oxide layer. The presence of deep and long grooves on the worn surface of the base alloy can be seen clearly in Figure 5.23(a). The number of grooves significantly decreased with 0.3 wt.% graphene addition (Figure 5.20(d)), and the addition of graphene resulted in the evolution of 125 smooth regions (Figure 5.23(b)). Furthermore, the worn surfaces of composites containing lower concentrations of graphene exhibited comparatively smoother regions than the unreinforced B319 Al alloy, indicating decreased magnitude of abrasive and delamination wear (Figure 5.20, Figure 5.21, and Figure 5.22). Zhang et al. [35] similarly observed the formation of relatively shallower and narrower wear grooves on the surface of the A355 Al alloy-1 vol.% graphene composite compared to the worn surface of the base A355 Al alloy. Figure 5.22. (a) SEM image of the worn surface of the B319-0.2 wt.% graphene composite under a 10 N load and its corresponding EDS area maps (b). 126 Figure 5.23. SEM image of worn surfaces of (a) B319 Al alloy at 100x magnification and (b) B319-0.3 wt.% graphene composite at 500x magnification. Figure 5.24 and Figure 5.25 show the worn surfaces of B319-0.5 wt.% graphene and B319-0.7 wt.% graphene composites, respectively. The worn surfaces of composites with higher percentages of graphene (0.5 and 0.7 wt.%) showed an increased magnitude of abrasive and delamination wear compared to the worn surface of the B319-0.3 wt.% graphene composite. This increase can be primarily attributed to higher porosity and reduced hardness caused by the agglomeration of graphene in the bulk of the B319 Al alloy composite, as shown in Figure 5.24 and Figure 5.25. The regions with graphene agglomeration are considered weak; they act as stress concentrators under the load, consequently accelerating wear. Raj et al. [316] observed improvement in the wear resistance of AA7075 Al alloy with graphene addition up to 0.75 wt.%. However, further addition of graphene showed deterioration in wear properties, which was mainly attributed to pore formation caused by the agglomeration of graphene. Similarly, Ghazaly et al. [238] reported an optimum content of 3 wt.% graphene for improving the wear resistance of AA2124 Al alloy, with inferior wear resistance observed in 5 wt.% graphene composite. Furthermore, the worn surface of 5 wt.% graphene composite exhibited severe abrasive and delamination wear, whereas the worn surface of the 3 wt.% graphene composite showed the presence of smooth regions associated with a graphene-containing tribolayer. 127 In the present study, the change in the topography of the worn surface with 0.3 wt.% graphene addition can be attributed to increased hardness (Figure 4.23(a)) and the presence of a graphene-containing tribolayer on the worn surface. Figure 5.26(a,g,h) shows the worn surface of the B319-0.3 wt.% graphene composite, revealing the presence of graphene on the worn surface. A high-magnification SEM image of graphene (region A) marked in Figure 5.26(a) and the corresponding carbon EDS map Figure 5.26(c)) further confirmed the presence of graphene on the worn surface. Furthermore, the point EDS analysis of graphene (region C) marked with a red square box in Figure 5.26(d) showed the presence of carbon along with O, Al, Si, Cu, and Fe (Figure 5.26(i)). The elements other than C in the EDS spectrum of graphene can be attributed to EDS-matrix interaction [317]. Figure 5.24. SEM images of worn surfaces and their corresponding EDS area maps for B319-0.5 wt.% graphene composite. Figure 5.25. SEM images of worn surfaces and their corresponding EDS area maps for B319-0.7 wt.% graphene composite. 128 Figure 5.26. (a) SEM image of a worn surface in B319-0.3 wt.% graphene composite, showing the presence of graphene; High magnification images of graphene marked in Figure (a): (b) region A, (c) carbon map of Figure (b), (d) region B, and (e) optical image of B319-0.3 wt.% graphene composite’s worn surface captured by Raman spectrometer, (f) Raman spectra of graphene (region D) marked in Figure (e), (g,h) SEM images of worn surfaces, showing tribo-morphological changes of graphene, (i) EDS elemental spectrum of a region (red square box, C) marked in Figure (d) and its quantification data. 129 Figure 5.26(f) shows the Raman spectroscopy analysis of B319-0.3 wt.% graphene composite’s worn surface marked with a red square box (region D) in Figure 5.26(e). The Raman spectra showed Raman fingerprints of graphene structure, corroborating SEM-EDS results. It is worth noting from Figure 5.26(f) that the average ID/IG ratio of the graphene increased from 0.4 to 1 as a result of the wear test, indicating amorphization or increased disorder in the graphene. Similarly, graphene's average I2D/IG increased from 0.37 to 0.83, suggesting a decrease in graphene layers from multi-layer (5−10) to a bi-layer range [249]. Due to weak interlayer shear rigidity in multi-layer graphene, graphene layers slide over each other even under moderate shear loadings [318]. Then, eventually, the interlayer shear deformation occurs, leading to the separation of individual graphene layers from the bulk graphene, a process known as mechanical exfoliation. The exfoliation of graphene layers can result in new interfaces, edges within the graphene, and the generation of graphene debris (Figure 5.23(b)), all of which increase the disorder. Previous reports have also reported similar observations of increasing disorder in graphene after the wear test [319–321]. SEM analysis of graphene on worn surfaces further suggests that many tribo-morphological phenomena, such as warping [322], bending [323], buckling [324], folding [325], etc., could occur in combinations or simultaneously during the wear test (Figure 5.26(b,d,g,h)), which can play a crucial role in improving wear resistance of B319 Al alloy. 5.2.5.2 Wear Debris Analysis Figure 5.27(a-f) and Figure 5.28(a-f) show the SEM images of the wear debris of B319-graphene composites tested under 5 N and 10 N loads, respectively. The wear debris from the base alloy contained mostly coarse flakes with an average size of 120 ± 35 µm regardless of the load, as shown in Figure 5.27(a) and Figure 5.28(a). These coarse flakes have been generated from the delamination wear. The mixture of flakes and fine debris can be seen in wear debris of the B319-0.1 wt.% graphene composite. At both loads, the fraction of fine debris increased with increasing graphene content, and 0.3 wt.% graphene composite predominantly showed fine equiaxed particles (3–15 µm) with fewer coarse flakes, as shown in Figure 5.27(d) and Figure 5.28(d). However, the composites with higher concentrations of graphene (0.5 and 0.7 wt.%) showed decreased fine debris and increased coarse flakes, attributed to the increased high porosity and agglomeration of graphene in the as-cast B319-composites (Figure 5.27(e,f) and Figure 5.28(e,f)). 130 The generation of very fine wear debris in the B319-0.3 wt.% graphene composite may be due to the high stability of the MML as a result of reduced friction and enhanced layer hardness, both of which would decrease the delamination wear. The fine wear particles (bright phase in SEM micrographs) were predominantly oxides. Thus, it can be inferred from these findings that delamination was dominant in the unreinforced alloy, but the extent of delamination decreased, and MML breakage accompanied by oxidative and adhesion wear became the prominent rate-controlling mechanism in the B319-0.3 wt.% graphene composite. Moreover, the changes in the wear debris morphology of B319-graphene composites are consistent with their corresponding worn surface topographies in Figure 5.16, Figure 5.18, and Figure 5.20. Wang et al. [326] fabricated Al-Cu3.7-Mg1.3/in-situ TiC composites by HP and extrusion. The wear debris of unreinforced Al-Cu3.7-Mg1.3 alloy comprised of large flake-like wear debris, and the wear debris of the composite possessed a mixture of fine and flake-like particles. The decreased wear debris size in composite was ascribed to enhanced mechanical properties, which was seen to inhibit delamination wear. These results regarding wear debris are consistent with the findings of the present study. Dwivedi et al. [308] studied the effect of alloying elements on the wear behavior of Al-17%Si alloy. It was observed that the addition of alloying elements improved the wear resistance of the Al-17%Si alloy, attributed to enhanced mechanical properties. This improvement was also reflected in the size of wear debris and fine wear debris generated during the sliding of multi-component alloy, indicating oxidative wear. In contrast, large wear debris was generated during the sliding of an Al-17%Si alloy, indicating severe wear by delamination. In another study, Wang et al. [327] investigated the effect of processing routes such as spray-deposition and conventional casting on the wear behavior of Al-20Si alloy. It was observed that spray-deposited alloy showed superior wear resistance compared to the alloy fabricated by conventional casting. The improved wear resistance of the spray-deposited alloy was attributed mainly to its refined microstructure and small Si particles. The worn surface of the conventional casting alloy exhibited the presence of fractured Si particles, large dimples, and delaminated regions, generating coarse debris. In contrast, the worn surface of the spray-deposited alloy showed small dimples and decreased delaminated regions, generating fine debris. Therefore, the wear in the cast alloy and spray-deposited alloy occurred mainly by delamination and oxidative mechanisms, respectively. 131 Figure 5.27. SEM images of wear debris from B319-graphene composites with varying amounts of graphene tested under a 5 N load: (a) 0 wt.% graphene, (b) 0.1 wt.% graphene, (c) 0.2 wt.% graphene, (d) 0.3 wt.% graphene, (e) 0.5 wt.% graphene, and (f) 0.7 wt.% graphene. 132 Figure 5.28. SEM images of wear debris from B319-graphene composites with varying amounts of graphene tested under a 10 N load: (a) 0 wt.% graphene, (b) 0.1 wt.% graphene, (c) 0.2 wt.% graphene, (d) 0.3 wt.% graphene, (e) 0.5 wt.% graphene, and (f) 0.7 wt.% graphene. 133 5.2.5.3 Subsurface Analysis of Wear Track 5.2.5.3.1 Subsurface Microstructure of Wear Track Figure 5.29(a-c) and Figure 5.30(a-c) show the subsurface microstructure evolution in the B319 Al alloy during dry sliding contact conducted under a 5 N and 10 N loads, respectively. The subsurface microstructure of the B319 Al alloy revealed the presence of three distinct zones designated as: MML (Zone 1), plastically deformed material (Zone 2), and undeformed bulk material (Zone 3). Similar zones were also observed in the subsurface microstructure of B319-0.3 wt.% graphene composites tested under 5 N (Figure 5.31(a) and Figure 5.32(a)) and 10 N loads (Figure 5.33(a,b)). EDS area elemental maps of the topmost layer (marked as “Zone 1”) of the B319 Al alloy’s cross-section revealed the presence of mechanically mixed Al, fractured Si particles, and oxide fragments, confirming the formation of MML during sliding (Figure 5.29(c), and Figure 5.30(c)). MML of B319-0.3 wt.% graphene composite showed carbon-rich regions in addition to elements observed in the unreinforced alloy, confirming the presence of graphene (Figure 5.31(b), Figure 5.32(b), and Figure 5.33(c)). The formation of MML usually occurs layer-by-layer as follows: (i) generation of wear particles, (ii) comminution or fragmentation of wear particles, (iii) mixing and agglomeration, followed by (iv) compaction of the mechanically mixed wear particles [328]. When two surfaces come into contact during a wear test, an initial transfer of material occurs between them due to adhesion and/or abrasion. As the sliding continues, frictional forces and mechanical loads cause material removal, generating wear debris. These debris undergo simultaneous mixing, agglomeration, and comminution. Abrasive grooves or delaminated regions seen in Figure 5.16 and Figure 5.20 often serve as preferential sites for debris agglomeration. The sliding motion of the contact surfaces results in relative movements among the particles that are trapped between contact surfaces. These particles are subjected to further mechanical action, such as deformation and fragmentation, facilitating their further mechanical mixing with evolving oxides. When the thickness of accumulated wear debris in grooves or delaminated regions exceeds the level of the sample’s contact surface, the pressing and flattening effects of the counterbody lead to local compaction or welding of the debris. These compacted layers adhere firmly to the base material below, providing a foundation for MML formation [328]. Over time, the repeated interaction between the wear debris and the contacting surfaces results in the formation of a thick 134 MML with a sandwich structure along the sliding direction [301]. The formation of MML has been widely reported in Al alloys under dry sliding conditions [300,301,329–332]. The MML exhibits an altered microstructure and mechanical properties of the bulk material. After the MML formation, it can serve as a protective layer before critical conditions are reached, leading to the detachment or delamination of loose debris from the MML. Gui et al. [333] also observed the formation of MML in Al-6Cu-Mn alloy during dry sliding wear, attributing the enhanced wear response of this alloy primarily to MML formation. Additionally, the MML was found to exhibit similar features to the wear debris, suggesting that the wear debris was primarily generated from the MML in their study. Therefore, the wear rate depends on the rate of formation and the subsequent delamination or detachment of the MML. The MML of B319 Al alloy exhibited non-uniform thickness across the wear track, showing a wavy pattern with a layered structure in the subsurface of the wear tracks of both materials (B319 Al alloy and its composites) (Figure 5.29(a,b), Figure 5.32(a) and Figure 5.33(a,b)). The uneven thickness and wavy pattern of the MML on the worn surface suggest the simultaneous occurrence of different stages of MML formation in different locations, including wear debris formation, mechanical mixing, compaction, and welding [328]. The layered structure of the MML on the worn surface indicates a layer-by-layer formation through the compaction of particulate aggregates followed by partial welding of layers. The distinct sublayers of the MML are evident in both the base alloy and its composite, as shown in Figure 5.29(b) and Figure 5.33(b). The maximum thickness of the MML on the worn surface of B319 Al alloy tested under a 5 N load and 10 N load was ~27 µm and ~6 µm, respectively. At higher loads, there’s typically more severe wear, leading to an increased generation of wear debris. This can result in severe breakage or delamination of the MML, contributing to its reduced thickness. Zhan and Zhang [328] also observed an increase in the thickness or coverage rate of MML in Cu-SiC composites with increasing load up to 100 N, but a further increase in load led to a decrease in MML thickness. Similar observations were also made by Kim et al. [334], where the MML thickness in A6061 alloy increased up to 150 N, but then decreased with a further increase in load. In another study, Venkataraman and Sundararajan [330] observed the formation of MML at 30 N in pure Al, but MML was absent at 52 N and 122 N. Thus, there exists a critical load for a given material above which the rate of MML delamination exceeds its rate of MML formation, leading to the absence 135 of MML formation [335]. It is also evident from Figure 5.30 that in most areas on the worn surface of B319 Al alloy tested at 10 N load, the MML is absent, indicating complete delamination of the MML, leading to higher wear, as seen in Figure 5.13 and Figure 5.14. Figure 5.29. Cross-sectional SEM-BSE micrographs of the worn surface of B319 Al alloy tested under 5 N load: (a) at low magnification, (b) at high magnification, and (c) EDS area elemental maps of Figure (b), showing three different zones. 136 Figure 5.30. Cross-sectional SEM-BSE micrographs of the worn surface of B319 Al alloy tested under a 10 N load: (a) at low magnification, (b) at high magnification, and (c) EDS area elemental maps of Figure (b), showing three different zones. 137 Figure 5.31. (a) Cross-sectional SEM-BSE micrograph of the worn surface of B319-0.3 wt.% graphene composite tested under a 5 N load at low magnification and corresponding EDS area maps (b), showing three different zones. 138 Figure 5.32. (a) Cross-sectional SEM-BSE micrograph of the worn surface of B319-0.3 wt.% graphene composite tested under a 5 N load at high magnification and corresponding EDS area elemental maps (b), showing three different zones and stable MML. 139 Figure 5.33. Cross-sectional SEM-BSE micrographs of the worn surface of B319-0.3 wt.% graphene composite tested under a 10 N load, (a) at low magnification, (b) at high magnification, and (c) EDS area elemental mapping of Figure (b), showing stable MML and its sublayers. 140 The MML of the B319-0.3 wt.% graphene composite was different from that of the base alloy, with a thickness of ~14 µm at 5 N (Figure 5.32(a)) and ~20 µm at 10 N (Figure 5.33(a,b)) load, respectively. The reduced thickness of MML in the B319-0.3 wt.% graphene composite at 5 N can be related to the enhanced load-bearing capacity and thermal stability caused by the presence of graphene, limiting the depth of deformation of the substrate. However, at higher loads (10 N), the contact stresses were sufficient to cause more deformation and plastic flow on the near surface, leading to an increased MML formation rate or MML thickness. The higher thickness of the composite’s MML in comparison to the MML of base alloy can be related to the high stability of MML in the composite, minimizing delamination wear. It was also reported that the presence of hard, stable, and thin MML provides the best wear resistance [330]. Furthermore, WDS (Figure 5.35(b)) area elemental mapping analyses indicate that the MML was composed of Al, O, and fractured Si particles, indicating the aluminum oxide formation and extensive intermixing of secondary phases at the contact surface. These WDS results agree well with the EDS data (Figure 5.32(b)). It was also evident from area elemental mappings that there was no material transfer from the counter body (100Cr6 ball) to the wear track since the elemental maps did not indicate the presence of Cr. The Fe map coincided with the Mn map, suggesting that the Fe in the MML mainly originated from the Al15 (Fe, Mn)3Si2 phase of the B319 Al alloy, and these observations further confirm that the ball did not wear out during sliding at both 5 N and 10 N loads. The wear track cross-section micrograph of the base B319 Al alloy showed the presence of subsurface cracks at the interface between MML (Zone 1) and the deformed layer (Zone 2), regardless of the load, as shown in Figure 5.34(a,b) and Figure 5.35(a). These cracks were associated with surface fatigue and its sub-mechanism delamination wear [294]. According to the delamination theory proposed by Suh [336,337], cyclic plastic deformation of the worn surface occurs due to repeated loading in sliding contact. This results in the generation and pile-up of dislocations at harder particles (e.g., Si), eventually subsurface cracks are initiated through the formation and coalescence of voids at high-stress concentration regions (e.g., Al-Si interfaces, defects, etc.) under cyclic loading, as shown in Figure 5.34(a,b) and Figure 5.35(a). These cracks propagate almost parallel to the wear surface and turn upwards when the crack reaches a critical length with continued sliding, resulting in delamination of the MML through the generation of 141 sheet-like wear debris, as seen in Figure 5.27(a) and Figure 5.28(a) [338]. The wear debris ready to delaminate can be seen on the worn surface of B319 Al alloy, as shown in Figure 5.30(a), Figure 5.34(a,b), and Figure 5.35(a). These results strongly suggest that the sheet-like particles generated by delamination accelerated the micro-plowing by acting as an abrasive, resulting in a synergistic effect where the wear rate was higher than what would be achieved by either mechanism alone [294]. This may be the reason for higher wear rates and volume loss in B319 Al alloy, as seen in Figure 5.13 and Figure 5.14. Cracks propagating parallel to the wear surface and wear debris with ~100 µm length can be seen in Figure 5.35(a). It was also evident from the cross-section of the B319 Al alloy that the MML was thin in most areas, and pristine Al was nearly exposed at the contact surface in some places, which was the result of severe wear caused by MML delamination. Unlike the subsurface cracks observed in the base B319 Al alloy, the wear-track cross-section micrograph of the B319-0.3 wt.% graphene composite showed no subsurface cracks, as shown in Figure 5.36(a)). Moreover, the MML in the composite appeared more continuous and denser, indicating reduced delamination wear and higher stability of the MML, which was also in agreement with the wear debris size as seen in Figure 5.27(d) and Figure 5.28(d). Similar observations were also reported by Wu et al. [201] in AlSi10Mg-graphene composites fabricated by SLM. In their work, the tribological behavior of the resultant composites was investigated under dry sliding conditions. The composite exhibited better friction and wear resistance over the base alloy at all the sliding conditions tested. This improvement was mainly attributed to the increased hardness and stable MML in the composite. The subsurface microstructure of the monolithic AlSi10Mg alloy showed the presence of microcracks, and the MML was severely damaged. In contrast, the subsurface microstructure of the AlSi10Mg-graphene composite showed compact MML with a smaller number of cracks, indicating the graphene’s addition limited subsurface crack nucleation and propagation. The subsurface cracks in the B319 Al alloy can also be related to the Si morphology; the sharp corners of the Si with needle-like morphology in the B319 Al alloy act as stress raisers in crack initiation and propagation, which would further enable the MML delamination. However, the refined (blunted) or fibrous Si in the B319-0.3 wt.% graphene composite likely decreased the magnitude of concentration, thereby reducing the localized deformation or debonding at the Al/Si interface [314,339]. 142 Figure 5.34. Cross-sectional SEM-BSE micrographs of the worn surface of B319 Al alloy tested under a 5 N load: (a) at low magnification, (b) at high magnification, and (c) EDS area elemental maps of Figure (b), showing the presence of cracks and delamination of MML. 143 Figure 5.35. (a) Cross-sectional EPMA-BSE micrograph of the worn surface of B319 Al alloy tested under a 10 N load, and (b) WDS area elemental mapping of the region (red square box) marked in Figure (a), showing the presence of cracks and delamination of MML. 144 Figure 5.36. (a) EPMA-BSE cross-sectional micrograph of the worn surface of B319-0.3 wt.% graphene composite tested under a 10 N load, and (b) WDS area elemental mapping of the region (red square box) marked in Figure (a). Just beneath the MML, a plastically deformed layer (Zone 2) of ~15 µm and ~10 µm can be observed in the cross-section microstructure of the B319 Al alloy (Figure 5.29(a,b)) and B319-0.3 wt.% graphene composite (Figure 5.32(a)), respectively. Similarly, at 10 N load, the B319 Al alloy (Figure 5.30(a,b)) and B319-0.3 wt.% graphene composite (Figure 5.33(a,b)) showed ~17 µm and ~8 µm layer thickness, respectively. The lower thickness of the plastically deformed layer in the composite may be associated with higher hardness and the presence of graphene, which can limit frictional plastic deformation by hindering dislocation motion [340]. The Zone 2 region of the B319 Al alloy and B319-0.3 wt.% graphene composite shows features of plastic flow, fractured Si, and Al2Cu, suggesting severe plastic deformation. 145 Figures 5.37(a-c) and 5.38(a-c) show the cross-section microstructure of the worn surface of B319-0.5 wt.% graphene and B319-0.7 wt.% graphene composites, respectively. Similar to the observations in the B319 Al alloy and the B319-0.3 wt.% graphene composite, comparable subsurface zones were also observed in the cross-section microstructure of the B319-0.5 wt.% graphene (Figure 5.37(b)) and B319-0.7 wt.% graphene (Figure 5.38(a)) composites. In contrast to the thick MML in the B319-0.3 wt.% graphene composite, the composite with high graphene content (0.5 and 0.7 wt.%) showed decreased MML thickness. In fact, in most places, the MML was delaminated, and fresh Al matrix was exposed on the contact surface, as observed in the unreinforced B319 Al alloy. Additionally, the subsurface microstructure of these composites revealed the presence of subsurface cracks, pores, and agglomeration of graphene, as shown in Figures 5.37(a) and Figure 5.38(a,b). It is evident from Figure 5.37(b) and Figure 5.38(b) that the subsurface cracks were nucleated at pores and the matrix/graphene interface. As mentioned earlier, the pores create stress concentrations and can initiate the formation of microcracks around them, eventually leading to the delamination of the MML [341]. Similarly, the agglomerated graphene may not have good metallurgical bonding with the surrounding matrix, leading to a weak interface. The weak interface diminishes the effectiveness of load transfer from the Al matrix to graphene, making these interfaces more susceptible to cracking under reciprocal sliding. The cracks propagate and create a network of cracks connecting different subsurface pores or weak Al/graphene interfaces (Figure 5.37(b)), eventually extending to the worn surface and generating sheet-like wear debris. The presence of graphene along the crack can be seen in Figure 5.38(b,c). These results make it evident that the presence of pores and agglomeration of graphene in B319-0.5 wt.% and B319-0.7 wt.% graphene composites contributed to the diminished tribological behavior in these materials, as seen in Figure 5.13 and Figure 5.14. Zeng et al. [342] fabricated hybrid composites of Al-7 wt.% Si reinforced with 10 wt.% SiC and varying amounts of rGO (0.3, 0.5, and 0.7 wt.%) using a conventional powder metallurgy route followed by hot extrusion. Among the composites, one with 0.5 wt.% rGO exhibited superior tribological behavior, and further addition of rGO to 0.7 wt.% rGO degraded the tribological behavior of the Al-7Si alloy. The degradation of tribological properties at 0.7 wt.% rGO concentration was attributed to the formation of micro-cracks in the subsurface region resulting from the agglomeration of the rGO. 146 Figure 5.37. Cross-sectional SEM-BSE micrograph of the worn surface of B319-0.5 wt.% graphene composite tested under a 10 N load: (a) at low magnification, (b) high magnification image of the region marked in Figure (a) and corresponding EDS area elemental maps (c), showing the presence of cracks and graphene agglomeration. 147 Figure 5.38. Cross-sectional SEM-BSE micrograph of the worn surface of B319-0.7 wt.% graphene composite tested under a 10 N load: (a) at low magnification, (b) high magnification image of the region marked in Figure (a) and corresponding EDS area elemental maps (c), showing the presence of cracks and delamination of MML. 148 5.2.5.3.2 Subsurface Hardness of Wear Track Figure 5.39(a) and Figure 5.39(b) shows the variation in cross-sectional hardness of the B319 Al alloy and B319-0.3 wt.% graphene composite as a function of depth from the worn surface tested under 5 N and 10 N loads, respectively. Both specimens exhibited a higher cross-sectional hardness near the worn surface, regardless of the applied load. At 5 N load, the top layer of the B319 Al alloy and B319-0.3 wt.% graphene composite showed a hardness of 90 HV and 110 HV, respectively. These hardness values indicate ~17% and ~21% increments compared to their respective bulk hardness of the matrix materials. Similarly, at 10 N, the top layer of the B319-0.3 wt.% graphene composite's hardness was 124 HV, indicating a 16% and 36% increase with respect to the unreinforced alloy’s top layer (107 HV) and base composite’s hardness (91 HV). All the tested samples exhibited a decreasing trend in hardness with increasing depth from the worn surface, eventually reaching the baseline hardness of the matrix material. Furthermore, the top surface of B319 Al alloy and B319-0.3 wt.% graphene composite tested under a 10 N load showed a ~19% and ~13% higher hardness over respective samples tested under a 5 N load. The increase in the hardness below the worn surface is attributed to strain hardening resulting from severe plastic deformation occurring under frictional stress during the wear test. The increase in load results in increased plastic deformation in the material near the worn surface, leading to increased hardness. Figure 5.39. Cross-sectional microhardness of subsurface layers as a function of distance from worn surface for B319 Al alloy and B319-0.3 wt.% graphene composite tested under loads: (a) 5 N and (b) 10 N. 149 It is known that the effect of strain hardening is the greatest near the wear track and decreases with increasing depth into the matrix material. The higher near-surface hardness of B319-0.3 wt.% graphene composite over base alloy may be attributed to the increased dislocation density, which is associated with dislocation-graphene interactions [31]. Graphene can effectively hinder the dislocation movement through pinning, facilitated by its high surface area and sheet morphology. This obstruction leads to the accumulation of more dislocations at the Al/graphene interface, thereby contributing to an enhanced strain hardening capability and increased hardness in the composite [273,343]. Wu et al. [201] also observed higher cross-sectional hardness of subsurface layers in AlSi10Mg-graphene composite over unreinforced AlSi10Mg alloy. However, the strengthening effect via pinning is less pronounced when graphene is present in an agglomerated form, as observed in the B319-0.5 wt.% and 0.7 wt.% graphene composites of this study. Figure 5.40(a-d) and (e-h) presents the schematic illustration of the rate-controlling wear mechanisms in the B319 Al alloy and the B319-0.3 wt.% graphene composite, respectively. The wear in the B319 Al alloy occurred mainly by MML delamination through subsurface crack propagation (Figure 5.40(c,d)), which was responsible for severe wear and generating coarse sheet-like wear debris (Figure 5.40(d)), as directly observed in Figure 5.27(a) and Figure 5.28(a). In the B319-0.3 wt.% graphene composite, wear occurred via MML breakage through crack propagation within MML (Figure 5.40(g)), producing equiaxed wear debris (Figure 5.40(h)), as seen in Figure 5.27(d) and Figure 5.28(d). The interfaces between sublayers within the MML are considered weak regions since the MML forms in a layer-by-layer mode (Figure 5.40(b,d)). The formed MML protected the surface until it reached a critical thickness, and then, crack initiation occurred at the interfaces between the MML sublayers or oxide regions in MML (Figure 5.40(g)), ultimately leading to MML breakage (Figure 5.40(h)) [328,344]. Figure 5.40(i-l) illustrates the schematic of the wear process in the B319-0.5, 0.7 wt.% graphene composites. The wear process predominantly involved through MML delamination in composites with higher concentrations of graphene (0.5 and 0.7 wt.%), similar to the wear process observed in the base alloy. Subsurface cracks originated at pores and the regions of graphene agglomeration, as seen in Figure 5.37(b) and Figure 5.38(a). Pores and agglomerated regions of graphene serve as stress concentrators, leading to localized stress under cyclic loading during sliding. This leads to elevated stresses in their proximity, potentially exceeding the material’s strength and initiating 150 crack formation Figure 5.40(k). These cracks then propagated, typically following the interface between the MML and the matrix. As they extended, they propagated parallel to the wear surface. Eventually, the cracks turned upwards towards the wear track (Figure 5.40(l)). Shearing forces during sliding contact can facilitate the delamination of the MML, generating sheet-like coarse wear particles, as seen in Figure 5.27(e,f) and Figure 5.28(e,f). Figure 5.40. Wear process in the B319 Al alloy: (a) initial microstructure with coarse as-cast grain structure, (b) formation of a thick MML and a deformed layer in B319 Al alloy, (c) subsurface crack initiation, (d) MML delamination, generating coarse wear debris; Wear process in the B319-0.3 wt.% graphene composite: (e) initial microstructure with refined as-cast grain structure, (f) formation of compact and graphene-containing MML or tribolayer, (g) crack initiation at the interface of sublayers within MML, and (h) MML fragmentation and generation of equiaxed wear debris; Wear process in the B319-0.5, 0.7 wt.% graphene composites: (i) initial microstructure, showing the presence of pores and graphene agglomeration, (j) formation of MML, (k) crack 151 initiation at pores and regions of graphene agglomeration, (l) MML delamination, generating sheet-like coarse wear particles. The high hardness of the composite and the effective heat dissipation away from the worn surface associated with the presence of graphene can decrease the transfer of shear force to the base material underneath, thereby contributing to a reduced wear rate for the composite. Additionally, the decreased stress concentration at the Si/Al matrix interface caused by finer Si particles in the composite can decrease the tendency for crack nucleation, thereby lessening the possibility of subsurface delamination [284]. The formation of a graphene-containing tribolayer or stable MML on the worn surface of B319-0.3 wt.% graphene composite contributes to enhancing the tribological properties of the B319 Al alloy. These results demonstrate the dominant wear mechanisms or energy dissipation pathways in both B319 Al alloy and B319-0.3 wt.% graphene composite, including adhesive, abrasive, oxidation, and surface fatigue wear mechanisms. However, graphene addition limited the magnitude of abrasive and delamination (surface fatigue) wear mechanisms, thereby improving the overall wear resistance of the B319 Al alloy. 5.3 Chapter Summary The present chapter discussed the friction and wear behavior of B319-graphene (0-0.7 wt.%) composites under dry sliding conditions. A detailed microstructural analysis of the wear tracks, wear debris, worn surfaces, and their subsurface was carried out to elucidate the dominant wear mechanisms in these composites. The salient points extracted from this chapter are summarized below. All tested specimens exhibited two distinct stages in their CoF curves, identified as the running-in and steady states. The running-in regime showed fluctuations in CoF, and this was attributed to the microtopography (surface texture and roughness) of the surface features, changes in the strain hardening state, initial elastic-plastic deformation of the surface asperities, and small contact areas. All the composites reached a steady state after 10 min of sliding at 5 N. However, B319-0.5 wt.% and 0.7 wt.% graphene composites reached a steady state after 15 min of sliding at a higher load of 10 N, possibly due to the contact pressure between surfaces. Additionally, CoF fluctuations around the mean value were observed for all samples in the steady-state regime regardless of the load, which was attributed to the stick-slip phenomenon caused by adhesion work. The average 152 CoF of B319 Al alloy decreased with graphene addition up to 0.3 wt.% graphene, then increased at 0.5 wt.% and 0.7 wt.% graphene concentrations. The B319-0.3 wt.% graphene composite exhibited ~52% and ~42% reduction in CoF at 5 N and 10 N loads, respectively, compared to the base alloy. Table 5.4 presents the average CoF of the running-in and steady-state regimes of B319-graphene composites. Similarly, Table 5.5 shows the average overall CoF (running-in and steady-state regime) of B319-graphene composites. Table 5.4. Average running-in and steady-state CoF of B319-graphene (0-0.7 wt.%) composites. Sample Running-in state Steady-state 5 N 10 N 5 N 10 N B319 0.49 (±0.08) 0.27 (±0.08) 0.40 (±0.07) 0.26 (±0.05) B319-0.1 wt.% graphene 0.31 (±0.11) 0.27 (±0.10) 0.26 (±0.12) 0.24 (±0.07) B319-0.2 wt.% graphene 0.26 (±0.08) 0.27 (±0.09) 0.24 (±0.09) 0.16 (±0.08) B319-0.3 wt.% graphene 0.22 (±0.07) 0.25 (±0.09) 0.20 (±0.06) 0.15 (±0.07) B319-0.5 wt.% graphene 0.46 (±0.07) 0.25 (±0.08) 0.38 (±0.03) 0.22 (±0.06) B319-0.7 wt.% graphene 0.46 (±0.08) 0.28 (±0.08) 0.39 (±0.05) 0.24 (±0.06) Table 5.5. Average overall (running-in and steady-state regime) CoF of B319-graphene (0-0.7 wt.%) composites. Sample Overall 5 N 10 N B319 0.42 (±0.08) 0.26 (±0.06) B319-0.1 wt.% graphene 0.27 (±0.1) 0.25 (±0.07) B319-0.2 wt.% graphene 0.24 (±0.09) 0.18 (±0.09) B319-0.3 wt.% graphene 0.20 (±0.06) 0.17 (±0.09) B319-0.5 wt.% graphene 0.40 (±0.05) 0.23 (±0.07) B319-0.7 wt.% graphene 0.41 (±0.06) 0.25 (±0.07) Furthermore, the average CoF (during running-in and steady-state regimes) of all tested samples was lower at 10 N compared to a 5 N load. The decreased CoF with increasing load from 5 N to 153 10 N load may be attributed to easier interlayer sliding of graphene at higher loads, resulting in a reduction in CoF. The improved friction performance of B319 Al alloy with graphene addition can be attributed to its self-lubricating ability, resulting from its unique structural and mechanical properties. Additionally, the enhanced heat dissipation capacity of the MML with graphene addition contributed to the reduction in the CoF. The degradation of friction properties at higher concentrations of graphene can be related to the lower hardness and higher porosity resulting from graphene agglomeration. The wear tracks of both B319 Al alloy and its composites reinforced with graphene exhibited an uneven wear pattern, further confirming the occurrence of the “stick-slip” phenomenon. The wear track became significantly narrower and shallower with the addition of graphene up to 0.3 wt.%. However, at higher graphene concentrations (0.5 and 0.7 wt.%), the maximum depth increased regardless of the load, indicating more severe wear. In this study, the counterbody (100Cr6 ball) exhibited no noticeable wear, irrespective of the load. This can primarily be attributed to the formation of a surface film on the ball’s surface via adhesive wear, which acted as a protective layer to minimize ball wear. Additionally, the high hardness of the ball prevented its wear. EDS area maps of the adhesive patches on the counterbody wear scar revealed the presence of Al, O, Si, Cu, and Mn elements, confirming material transfer from the substrate to the counterbody. In this study, two methods were employed to assess the wear volume and specific wear rate of B319-graphene composites: weight loss and 3D laser microscope analysis. Both methods revealed similar trends in wear properties; however, laser microscope analysis offered more accurate wear property measurement due to its selective imaging and high-resolution capabilities. As observed in CoF, mass loss, wear volume, and specific wear rate, the addition of graphene up to 0.3 wt.% resulted in a decrease in wear for B319 Al alloy. However, with further graphene addition, there was an increase in wear at both 5 N and 10 N loads. The B319-0.3 wt.% graphene composite showed ~39% and ~52% reduction in wear volume at 5 N and 10 N loads, respectively, compared to the unreinforced alloy. Table 5.6 and Table 5.7 summarizes the wear volume and specific wear rate of B319-graphene composites, respectively. The worn surface of B319 Al alloy exhibited features of adhesion, abrasive, oxidation, and delamination wear. However, the intensity of these mechanisms decreased with graphene addition up to 0.3 wt.%. Conversely, the worn surface of the B319-0.5 wt.% and 0.7 wt.% graphene 154 composites displayed an increased degree of abrasive and delamination wear compared to the worn surface of the B319-0.3 wt.% graphene composite. This increase can be primarily attributed to higher porosity and reduced hardness from the agglomeration of graphene in the bulk of the B319 Al alloy composite. The areas of graphene agglomeration were clearly evident on the worn surfaces of B319-0.5 wt.% and 0.7 wt.% graphene composites. Table 5.6. Average wear volume (mm3) of B319-graphene composites evaluated by weight loss and laser microscope method. Sample 5 N 10 N Weight Loss Laser Microscope Weight Loss Laser Microscope B319 1.8 (±0.13) 1.6 (±0.005) 3.5 (±0.05) 3.8 (±0.02) B319-0.1 wt.% graphene 1.5 (±0.07) 1.5 (±0.005) 3.3 (±0.18) 3.2 (±0.06) B319-0.2 wt.% graphene 1.3 (±0.03) 1.3 (±0.01) 2.3 (±0.11) 2.4 (±0.07) B319-0.3 wt.% graphene 1.1 (±0.06) 1.0 (±0.06) 1.9 (±0.02) 1.8 (±0.12) B319-0.5 wt.% graphene 1.4 (±0.12) 1.4 (±0.22) 2.6 (±0.02) 2.8 (±0.01) B319-0.7 wt.% graphene 1.8 (±0.16) 1.7 (±0.20) 2.8 (±0.02) 2.9 (±0.04) Table 5.7. Average specific wear rate (mm3/N.m×10-4) of B319-graphene composites evaluated by weight loss and laser microscope method. Sample 5 N 10 N Weight Loss Laser Microscope Weight Loss Laser Microscope B319 11.8 (±0.13) 10.3 (±0.005) 11.8 (±0.18) 12.7 (±0.08) B319-0.1 wt.% graphene 10.1 (±0.07) 9.9 (±0.005) 11 (±0.06) 10.7 (±0.21) B319-0.2 wt.% graphene 8.6 (±0.03) 8.8 (±0.01) 7.7 (±0.37) 8.1 (±0.23) B319-0.3 wt.% graphene 7.5 (±0.06) 7.0 (±0.06) 6.3 (±0.06) 6.1 (±0.39) B319-0.5 wt.% graphene 9.6 (±0.12) 9.1 (±0.22) 8.7 (±0.06) 9.5 (±0.03) B319-0.7 wt.% graphene 12.1 (±0.16) 11 (±0.20) 9.4 (±0.06) 9.7 (±0.14) 155 The wear debris from the base alloy contained mostly sheet-like coarse flakes regardless of the load, primarily generated from delamination wear. At both loads, the fraction of fine debris increased with increasing graphene content, with the 0.3 wt.% graphene composite predominantly exhibiting fine equiaxed particles with fewer coarse flakes, indicative of reduced wear. However, composites with higher graphene concentrations (0.5 and 0.7 wt.%) showed decreased fine debris and increased coarse flakes, confirming severe wear in the B319-0.5, 0.7 wt.% graphene composites. The generation of very fine wear debris in the B319-0.3 wt.% graphene composite may be due to the high stability of the MML resulting from reduced friction and enhanced hardness, which can hamper delamination wear. The subsurface microstructure of all tested samples revealed three zones: MML, plastically deformed material, and undeformed bulk material. The wear track cross-section micrograph of the B319 Al alloy showed subsurface cracks, while that of the B319-0.3 wt.% graphene composite exhibited compact MML with no subsurface cracks. Additionally, the subsurface microstructure of B319-0.3 wt.% graphene composite showed compact MML with no subsurface cracks. Finally, the subsurface microstructure of the B319-0.5, 0.7 wt.% graphene composites exhibited subsurface cracks, pores, and graphene agglomeration. These results suggest that wear in B319 Al alloy mainly occurred through MML delamination via subsurface crack propagation, whereas in the B319-0.3 wt.% graphene composite, wear primarily resulted from MML breakage. The improved wear resistance of the B319-0.3 wt.% graphene composite in the present study was a direct consequence of its refined microstructure (grains and Si particles), increased hardness, and formation of stable MML or graphene-containing tribolayer at the contact surfaces. All the microstructural changes mentioned above and improved hardness with graphene addition limited the delamination wear mechanism, enhancing the overall wear resistance of the B319 Al alloy. However, in composites with higher graphene content (0.5 and 0.7 wt.%), the presence of pores and agglomerated graphene regions contributed to subsurface crack formation, exacerbating the wear of the composite. 156 Chapter 6: Conclusions and Future Work The main objective of this dissertation was: 1) to fabricate B319 Al alloy matrix composites with uniformly distributed graphene, and 2) to study the effect of graphene on the microstructure, friction, and wear behavior of the B319 Al alloy composites. To achieve this, the B319 Al alloy composites reinforced with varying amounts of graphene (0-0.7 wt.%) were fabricated through a novel distribution approach. Initially, a composite master alloy of F357-5 wt.% graphene was developed using a flake-PM process involving BM and SPS. Subsequently, the composite master alloy was introduced into the molten B319 Al alloy to fabricate B319-graphene composites. The obtained composites underwent thorough characterization, assessing the effect of graphene on the microstructure, hardness, electrical, and thermal conductivity of the new composite alloy. Moreover, the dry reciprocating sliding friction and wear response of B319-graphene composites were studied under 5 N and 10 N loads. The worn surface and its subsurface microstructure were characterized to elucidate the role of graphene in enhancing the tribological behavior of B319 Al alloy and to identify dominant wear mechanisms in these novel composites. The results of this research demonstrated that the B319-0.3 wt.% graphene composite exhibited optimal anti-friction and wear resistance. This improvement was attributed to the microstructural refinement, high stability of the MML, increased hardness, and the formation of graphene-containing tribolayer at the contact surface. In Section 6.1, the major findings from this study are presented, while Section 6.2 summarizes challenges faced during this research. The research limitations are highlighted in Section 6.3, and directions for future research are suggested in Section 6.4. 6.1 Conclusions 6.1.1 Composite Master Alloy 1. The composite master alloy of F357-5 wt.% graphene, fabricated through a flake-PM approach involving BM and SPS, demonstrated a uniform distribution of graphene. This indicates the effectiveness of this approach in dispersing graphene within the B319 Al alloy matrix. 157 2. The introduction of ball-milled F357-graphene powder mixture into the B319 Al alloy melt without SPS resulted in powder floating on the melt’s surface and the formation of clusters. This suggests that the addition of graphene into the B319 Al alloy melt in powder form was not effective. 6.1.2 Microstructure, Hardness, Electrical, and Thermal Conductivity of B319-Graphene Composites 1. The addition of graphene reduced the SDAS of the B319 Al alloy. This reduction was primarily attributed to a higher cooling rate induced by the increased heat extraction facilitated by the presence of graphene between the growing dendrites of α-Al. Furthermore, graphene provided additional nucleation sites, which contributed to restricting grain growth in the alloy. 2. The grain size of the B319 Al alloy decreased with graphene addition, showing the most significant grain refinement of ⁓45% with 0.7 wt.% graphene addition compared to the base alloy. However, agglomeration of graphene was observed in composites containing more than 0.3 wt.% graphene and was the result of strong van der Waals forces between the graphene sheets and a density mismatch between the B319 Al alloy and graphene. 3. The aspect ratio of Si particles decreased with graphene addition, indicating the transformation into a more fibrous form of Si. The reduced size of Si particles with graphene addition can be due to the grain refinement of primary Al, heterogeneous nucleation of Si on graphene, and the presence of graphene at the Si growth front, which impedes Si growth. 4. The hardness of B319 Al alloy increased from 77 to 92 HV with 0.1 wt.% graphene addition, an ∼18% improvement over the unreinforced matrix. However, further addition of graphene to 0.2 and 0.3 wt.% showed no significant change in hardness and further graphene addition decreased the hardness. The decrease in hardness at high concentrations of graphene was related to the agglomeration of graphene and higher porosity. 5. The addition of graphene did not yield significant improvements in the electrical conductivity and thermal conductivity of B319 Al alloy. This trend can be attributed to enhanced electron and phonon scattering resulting from the formation of structural imperfections, increased porosity, and an expansion in the grain boundary area. 158 6.1.3 Friction Behavior of B319-Graphene Composites 1. The friction curves of all tested samples exhibited two distinct regimes called “running-in” and “steady-state”, regardless of the load. 2. All the samples showed fluctuations in CoF around the mean value in the steady-state regime, irrespective of the load, which was attributed to the stick-slip phenomenon caused by adhesion work. 3. At both loads, the steady-state CoF of B319 Al alloy decreased with graphene addition up to 0.3 wt.% graphene, then increased at 0.5 wt.% and 0.7 wt.% graphene. This improvement was attributed to graphene’s self-lubricating ability through the formation of rollers and interlayer sliding, resulting from graphene's unique structural and mechanical properties. The degradation of friction properties at higher concentrations of graphene was related to the lower hardness and higher porosity resulting from graphene's agglomeration. 4. The overall magnitude of the CoF of all the samples was comparatively lower at 10 N compared to 5 N. This trend was associated with graphene's self-lubricating property, which facilitates easy interlayer sliding at higher loads, thereby reducing friction. Additionally, the more uniform distribution of contact force and formation of thick MML at higher loads are caused by elevated temperatures resulting from increased frictional heating. 6.1.4 Wear Behavior of B319-Graphene Composites 1. The wear tracks of all the tested samples showed uneven wear patterns, confirming the “stick-slip” phenomenon. 2. The wear tracks of the B319 Al alloy become shallower and narrower with graphene addition up to 0.3 wt.%, indicating enhanced wear resistance. 3. The wear resistance of the B319 Al alloy improved with graphene addition, and an optimum was reached at 0.3 wt.% graphene concentration, showing ∼39% and ∼52% improvement over base alloy at 5 N and 10 N loads, respectively. However, the addition of graphene beyond 0.3 wt.% resulted in an increased wear rate and volume loss mainly due to graphene agglomeration and increased porosity. 4. The cross-sectional microstructure of the base alloy and composite showed three distinct Zones designated as MML, plastically deformed layer, and undeformed structure. The 159 MML in the graphene-containing composite appeared more continuous and denser, suggesting improved stability of the MML and reduced delamination wear. 5. Microscopic examination of the worn surfaces, wear debris, and subsurface revealed that the wear of the unreinforced alloy was dominated by delamination caused by propagation of subsurface cracks followed by adhesion, oxidation, and abrasion mechanisms. Once graphene was added, the extent of delamination wear decreased, and the rate-controlling mechanism in 0.3 wt.% graphene composite shifted to MML breakage accompanied by adhesion, oxidative, and abrasion mechanisms. 6. The cross-sectional microhardness of subsurface layers for the B319 Al alloy and its composite showed increased hardness after the wear test, which confirmed strain hardening of subsurface layers during wear testing. Moreover, the top layer of the composite showed ∼22% and ∼16% higher hardness than that of the base alloy, which was attributed to dislocation and load strengthening contributed by graphene. 7. The B319-0.3 wt.% graphene composite showed the best anti-friction and wear performance compared to the base alloy. The improved friction and wear response of the B319 Al alloy with graphene addition was attributed to the tribo-morphological transformation of graphene (e.g., warping, bending, buckling, folding, etc.), the formation of a graphene-containing tribolayer at the contact surface, and high stability of MML contributed by the microstructural refinement, and higher hardness. 6.2 Challenges Faced During the Research The main experimental challenges encountered during this research are summarized below: 1. Agglomeration of graphene: Achieving a uniform dispersion of graphene in the B319 Al alloy is challenging, especially in the casting route used in this work. Various distribution approaches were explored, including (i) direct addition of graphene into molten alloy, (ii) addition of ball milled Al alloy-graphene powder, and (iii) addition of SPSed composite master alloy. It was found that the fabrication of composite master alloy via a flake-PM approach involving BM and SPS yielded the best dispersion. 2. Dilution of the B319 Al alloy: Initially, pure Al powder was used to fabricate the composite master alloys, resulting in a significant dilution of the B319 Al alloy. This dilution could 160 have a had a substantial effect on the properties of the resultant composites. Later, this challenge was overcome by using a pre-alloyed F357 Al powder, which has a similar composition to the B319 Al alloy. 3. Non-melting of composite master alloy: During the initial stage of this project, the SPSed composite master alloy did not melt completely in the molten B319 Al alloy. Later, by optimizing the graphene content in the composite master alloy and preheating the composite master alloy prior to adding it into the B319 Al alloy melt, this issue was resolved. 6.3 Limitations of the Research The following are the limitations of the generated results in this research: 1. In this study, the number of layers in graphene was estimated based on the Raman spectroscopy results, which may not always provide an absolute confirmatory measure. Therefore, to be certain of the precise structure of the graphene used in this study, confirming the number of layers of the graphene via atomic force microscopy or TEM would be recommended. 2. In the present study, the B319-graphene composites were fabricated at a laboratory scale, and the results generated may vary if industrial-scale processes are used. 3. The friction and wear tests in this study were carried out at room temperature under dry sliding conditions and a relatively high loading. These conditions differ from practical engine operating conditions, which typically involve lubrication, lower loads, and high temperatures. Therefore, the results obtained in this study may not fully represent the tribological behavior under actual engine conditions. 6.4 Recommendations for Future Work From the results of this dissertation, it is evident that graphene enhances the friction and wear response of the B319 Al alloy. Despite the lower conductivity observed, the B319-graphene composites demonstrate balanced strength, friction, and wear resistance, making them promising candidates for automotive parts where wear resistance is a critical design parameter. However, further exploration of the B319-graphene composites is necessary to address specific research gaps 161 and ensure their suitability for industrial applications. This section outlines recommended key research directions to advance the findings presented in this dissertation and the utilization of the B319-graphene composites for diverse automotive applications. 1. Study the friction and wear behavior of B319-graphene composites at an ultra-low wear regime under lubricated conditions to simulate engine operating conditions. 2. Investigate the friction and wear behavior of the B319-graphene composites against different counterbodies, such as ceramics, oxides, etc., to understand wear mechanisms against various materials. 3. Examine the friction and wear behavior of the B319-graphene composites at elevated temperatures to identify any unique wear mechanisms operating at higher temperatures. 4. 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Thesis/Dissertation
eng
Mechanical Engineering
Vancouver : University of British Columbia Library
Attribution-NonCommercial-NoDerivatives 4.0 International
Graduate
Lukas, Bichler
Fabrication of graphene-based master alloys for the development of B319 aluminum alloy-graphene composites with enhanced tribological properties
Text