© 2011 Daniel W. Maiorano All Rights Reserved

Daniel W. Maiorano

THE EVOLUTION AND IMPLICATION OF BORON CARBIDE

MICROSTRUCTURAL VARIATIONS AND TRANSFORMATIONS DURING

POWDER PROCESSING

DANIEL W. MAIORANO

A Dissertation submitted to the

Graduate School-New Brunswick

Rutgers, The State University of New Jersey

in partial fulfillment of the requirements

for the degree of

Doctor of Philosophy

Graduate Program in Materials Science and Engineering

written under the direction of

Professor Richard A. Haber

and approved by

______

New Brunswick, New Jersey

October, 2011

ABSTRACT OF THE DISSERTATION

The Evolution and Implication of Boron Carbide Microstructural Variations and

Transformations During Powder Processing

By DANIEL W. MAIORANO

Dissertation Advisor: Prof. Richard A. Haber

Boron carbide is a material of key interest in structural ceramic applications due to its extremely favorable physical properties including low theoretical density, high hardness, high strength, and excellent wear resistance. However, boron carbide microstructures often exhibit the presence of secondary phases, most typically composed of carbon. Under very high stress and strain rate conditions, boron carbide has been shown to have reduction in compressive strength which appears to be correlated to a possible pressure-induced creation of ~3 nm wide amorphous bands. As these variations have only been noted to date following application of high stress upon the boron carbide with no examination of powders prior to high stress, this dissertation attempts to determine whether these variations are inherent to boron carbide powders or introduced during typical powder processing.

Most boron carbides are produced by a large scale carbothermic reduction process within an electric arc furnace. While it remains the most cost effective production

method at this time, carbothermic reduction suffers from uneven heating rates across the reactive melt which makes it difficult to maintain a chemically uniform product. To determine the extent of preexisting variations, commercially available powders were compared with segments taken directly from the carbothermic reduction ingot via X-ray

Diffraction (XRD), Fourier Transform Infrared Spectroscopy (FTIR), and Raman

microspectroscopy. All powders were then subjected to various levels of comminution to

examine the level of pressure induced by common powder processing.

Analyses showed that due to the nature of the carbothermic reduction process,

significant stoichiometry variations of over 2 at% carbon are present across the length of

segments classified as “good” boron carbide visually. The comminution process

impacted the relative exposed carbon content of the boron carbides, resulting in powders

that were coated with carbonaceous materials. At high comminution energies in a

laboratory scale jet mill, resulting surfaces exhibited evidence of amorphized boron

carbide. Attempts at reversing the processing induced variations proved ineffective and

suggest inherent difficulties in altering the structure of boron carbide produced via

carbothermic reduction.

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Acknowledgements

My first thank you goes, of course, to my advisor, Prof. Haber. I still remember

the first day I met you, when Cari had just hired me as a lab tech and you wanted to show

me how “cool” corn starch could be if you added just a little bit of water. Just like my

hands were stuck in the corn starch for hours afterwards, my journey through graduate school has been filled with experiences that sound cool, yet sometimes took a little bit longer than expected! I must thank you for the unexpected levels of guidance, support, and mentorship over the years- without them, I would not be where I am today.

Secondly, thank you to my dissertation committee, Prof. Niesz, Prof. Mayo, Dr.

Domnich, Dr. Behler, and Dr. Rafaniello. Prof. Niesz has been an invaluable guide and mentor from the earliest days of class as an undergrad through the completion of my

Ph.D. Prof. Mayo has gone out of his way to help my endeavors with the use of his own equipment and offers of advice and guidance throughout the years. Dr. Domnich is incredible in his understandings of vibrational spectroscopies and is always willing to lend an ear. Dr. Behler has been so quick and helpful in responses to questions, and really helped shed light on real world problems I was encountering with analysis. Dr.

Rafaniello has invested amazing amounts of his own time and energy to contributing to my development as a researcher. You have all been of so much help in shaping this dissertation, rising above and beyond what is typically asked of a committee of your own free will. This would not have been possible without each of you.

To Laura Chirichillo and Michelle Sole- you two ladies have been responsible for keeping Prof. Haber’s group, Happy Valley, running smoothly and we flat out could not

exist without you two.

I’d like to thank all the professors in the MSE department that have all influenced

me over the years. In particular, thank you to Prof. Klein for keeping the graduate

program running, Prof. Greenhut for sharing his wealth of microscopy and general

knowledge, and Prof. Lehman, Prof. Matthewson, Prof. Chhowalla, and the rest of the department, thank you all. A very special thank you goes to the staff in the department who helped make my time here just a little bit easier, Claudia Kuchinow, Phyllis Cassell,

Phil Grill, and especially John Yaniero. John keeps literally everything operating in the department, from doors to microscopes and anything you can imagine, and I honestly

don’t know what I would have done without his aid on more occasions than I can count

without taking my shoes off.

I’d like to thank all the members of Happy Valley, present and past, for being

there with me during this trip and making it especially memorable: Steve Miller, Mihaela,

Sara, Vlad, Shawn, Cari, Navin, Scot, Laura Reynolds, Ray, Volkan, Chris, Steve

Mercurio (I’ll always remember the days we spent learning the hot press inside out trying

to fix it), Andrew, Steve “Stevie Poo” Bottiglieri, Doug, Fatih, Rob, Minh, Nick, and

Vince. Thanks also to my friends in other groups around the department who made times just a little more fun. Thanks to the many undergrads who’ve worked as lab techs for

Happy Valley. Rich, Sean, and Dan, you guys get the standout mentions for working directly for me and putting up with my unique way of thinking, but thanks also to

everyone else.

My friends deserve thanks for supporting me through all of graduate school- I

know it’s been tough at times! There are too many of you to list by name, but you have all made getting through the tough parts just a little bit easier.

Thank you to my family, especially my parents, for supporting me during the last six years, through all the glories and trials of graduate education, for not snapping at me nearly as often as I deserved, and for having a very reasonably priced apartment for rent when I needed it most. Matt, you’ve gone through four homes and seventeen cars (or so it seems) in the time I’ve been in grad school- an awful lot has happened in that time, and

I know you’ll still be there if I ever need you. Becky, you being off in Houston has been tougher than I’d ever admit to you in person, and I can’t wait to see you, Greg, and little

Nate again this summer. Finally, thank you to my wonderful girlfriend, Patty. One of the best parts of this whole journey has been meeting you, and the impact you’ve had in my life.

Table of Contents

Abstract of the Dissertation ...... ii

Acknowledgements ...... iv

Table of Contents ...... vii

List of Tables ...... xi

List of Figures ...... xii

1. Introduction ...... 1

2. Background ...... 3

2.1 Boron Carbide Historical Survey ...... 3

2.2 Production of Boron Carbide ...... 4

2.2.1. Melt Processing ...... 4

2.2.2 Chemical Vapor Deposition ...... 8

2.2.3 Single Crystal Production ...... 9

2.2.4 Rapid Carbothermal Reduction ...... 9

2.2.5 Other Production Routes ...... 10

2.3 Phase Diagram and Crystal Structure ...... 11

2.4 Properties and Applications ...... 15

2.4.1. Chemical Properties ...... 16

2.4.2 Density ...... 16

2.4.3 Elastic Properties ...... 17

2.4.4 Mechanical Properties ...... 18

2.4.5 Electronic Structure and Properties ...... 22

2.4.6 Nuclear Properties ...... 23

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2.4.7 Applications ...... 24

2.5 Common Inhomogeneities in Boron Carbide ...... 24

2.6 Amorphization of Boron Carbide under High Stress Conditions ...... 26

2.6.1 Ballistic Impact ...... 26

2.6.2 Nanoindentation ...... 30

2.6.3 Exploration of Amorphization ...... 34

2.7 Phase Transformations Induced by Comminution and Mechanical Working 39

2.7.1 Boron ...... 39

2.7.2 Carbon ...... 40

2.7.3 Boron Nitride ...... 41

2.7.4 Metals ...... 43

2.7.5 Carbides ...... 45

2.8 Comminution ...... 49

2.9 Analytical Techniques ...... 57

2.9.1 Vibrational Spectroscopy ...... 57

2.9.2 X-ray Diffraction ...... 65

3. Method of Attack ...... 71

3.1 Objective 1: Comminution of Boron Carbide ...... 71

3.2 Objective 2: Characterization of Comminuted Powders ...... 73

3.3 Objective 3: Elimination of Impurities and Reversibility of Transformations 74

4. Experimental Procedures ...... 76

4.1 Comminution ...... 76

4.2 Particle Analysis of Powders ...... 79

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4.2.1 Light Scattering Particle Size Analysis ...... 79

4.2.2 Scanning Electron Microscopy ...... 80

4.3 Vibrational Spectroscopy ...... 82

4.3.1 Raman Microspectroscopy ...... 82

4.3.2 Fourier Transform InfraRed Spectroscopy ...... 84

4.4 X-ray Diffraction ...... 86

4.5 Excess Carbon Removal ...... 88

4.6 Extrusion of Boron Carbide Rods ...... 90

5. Results and Discussion ...... 93

5.1 Comminution ...... 93

5.1.1 Low Energy Comminution: Vibratory Milling ...... 93

5.1.2 High Energy Comminution: Jet Milling ...... 97

5.2 Analysis ...... 101

5.2.1 X-ray Diffraction ...... 101

5.2.2 Fourier Transform InfraRed Spectroscopy ...... 111

5.2.2.1 FTIR for Carbon Content and Location ...... 111

5.2.2.2 FTIR Analysis for Presence of Amorphization ...... 113

5.2.3 Raman Microspectroscopy ...... 120

5.2.3.1 Raman for Carbon Content ...... 120

5.2.3.2 Raman Analysis for Presence of Amorphization ...... 124

5.2.4 Summation of Induced Variations ...... 131

5.2.4.1 Emergence of Pre-existing Compositional Variations .... 131

5.2.4.2 Evolution of Carbon ...... 132

5.2.4.3 Amorphization of Boron Carbide ...... 138

5.3 Elimination of Induced Variation ...... 139

5.3.1 Compositional Variations ...... 140

5.3.2 Removal of Excess Carbon ...... 141

5.3.2.1 Thermal Attack of Carbon ...... 142

5.3.2.2 Chemical Oxidation of Carbon ...... 147

5.3.3 Addressing the Amorphization ...... 150

5.3.3.1 Orientation Dependence of Amorphization ...... 150

5.3.3.2 Low Pressure Jet Milling ...... 155

5.3.3.3 Reversibility of Amorphization ...... 161

6. Conclusions ...... 164

7. Future Work ...... 168

Appendix: X-ray Diffraction Analysis of Boron Carbide ...... 174

References ...... 181

Curriculum Vita ...... 192

List of Tables

Table 2.1. Elastic properties for boron carbide, calculated for single crystal and observed carbon content dependence...... 18

Table 2.2. Anisotropic elastic constants for several boron carbides...... 18

Table 2.3. Comparison of high strength materials per their strength/density ratio...... 20

Table 2.4. Peak broadening of x-ray diffraction lines indicative of structural changes in boron carbide and densities of deformation mechanisms. “Sample” designates hot pressed body at 94+% theoretical density utilizing stated powder...... 49

Table 5.1. Vibratory milled boron carbides and select data obtained from analysis of the x-ray diffraction patterns...... 103

Table 5.2. Lattice parameters and resultant stoichiometric carbon content calculated from x-ray diffraction patterns for iteratively jet milled boron carbide...... 108

Table 5.3. Locations of the 480 cm-1 and 535 cm-1 peaks in samples with varying stoichiometric carbon and for the vibratory milled unprocessed boron carbide. Separation between the peaks suggests a decreasing stoichiometric carbon content from section C to section P, with section S having a much higher stoichiometric carbon content again. All values in cm-1...... 124

Table 5.5. Expected and observed intensity values and TC values for several planes of boron carbide extrudates. Both samples show relatively strong orientation along the (101) plane...... 152

List of Figures

Figure 2.1. Ingot of boron carbide just over two meters in diameter. Sections (a) are likely unreacted material while (b) are located where the electrodes sat and are likely well-reacted crystalline boron carbide with some possible graphitization, and will be visually sorted as such during initial classification...... 7

Figure 2.2. Phase diagram after Ekbom and Amundin, depicting B13C2 as the stoichiometrically stable phase of boron carbide and presuming the presence of several low temperature phases...... 12

Figure 2.3. Phase diagram after Beauvy, depicting the more widely accepted B4C as the stoichiometrically stable boron carbide, with solid solutions with boron and carbon on each respective side...... 13

Figure 2.4. Left: Boron carbide rhombohedral unit cell. Right: Equivalent hexagonal indexing depicting stacking sequencing...... 14

Figure 2.5. Calculated band structure for boron carbide. Left axis gives energies relative to valence band, right axis gives energies relative to conduction band...... 23

Figure 2.6. Dense body microstructures of boron carbide. Excess carbon impurities are readily visible as the bright white locations throughout the microstructure...... 26

Figure 2.7. Ballistic impact data from independent experiments conducted at the Army Research Laboratory and Sandia National Laboratory, depicting the drastic decrease in strength above the HEL of 20 GPa...... 29

Figure 2.8. Left: Optical image of ballistically impacted boron carbide. Right: HREM image taken close to the point of impact, showing amorphous band running through the sample...... 29

Figure 2.9. TEM of a 100 mN Berkovich nanoindent showing amorphization occurring along the (113) bands of the boron carbide crystal. a) Plan view; b) magnified image; c) and d) HREM of the corresponding boxes in b...... 32

Figure 2.10. Raman spectra of (a) pristine single crystal B4.3C; (b) indented single crystal; (c) indented hot-pressed polycrystalline material; (d) scratch debris of single crystal and (e) annealed scratch debris by using an argon ion laser with excitation wavelength of 514.5 nm...... 33

Figure 2.11. Raman spectra of deposited boron carbide films showing broad humps in amorphous materials as opposed to distinguishable peaks in crystalline materials resulting from deposition temperatures above 940°C...... 35

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Figure 2.12. FTIR spectra of boron carbides displaying the formation of the broad amorphous boron peak under dynamic indentation...... 37

Figure 2.13. Raman spectra taken during nonhydrostatic unloading of boron carbide, indicating rise of amorphous peak at ~1810 cm-1at approximately 16 GPa...... 38

Figure 2.14. X-ray diffraction patterns of bulk boron a, b) prior to comminution, c) after milling in iron mortar, and d) following leaching of mortar materials, displaying a strong trend for boron to amorphize under comminution...... 40

Figure 2.15. TEM showing crystal structure of pure synthetic graphite before (top) and after (bottom) vibratory milling, with amorphous structures dominating the post-milling material...... 41

Figure 2.16. HREM of h-BN ball milled for 180 hours, showing distinctly amorphous microstructure. Diffraction pattern insert shows only amorphous rings...... 42

Figure 2.17. HREM of mechanical alloyed a-BCN material. The mostly amorphous microstructure contains only short-range order of BCN material, evidenced by insert diffraction pattern displaying amorphous rings with minimal crystalline spot patterns. 43

Figure 2.18. XRD intensity versus planetary milling time for iron and silicon metals. Lowering of intensity combined with broadening of the peaks indicates a transformation to amorphous material as the metals were comminuted...... 44

Figure 2.19. TEM micrographs portraying areas of extreme strain (arrowhead, top) and dislocation buildup (vectors a and b, bottom) in ball milled copper. These deformation mechanisms were shown to be the precursors to amorphization...... 45

Figure 2.20. TEM showing post-polishing TiC as (a) microstructure with smoothed edges near polishing defects; (b and c) diffraction spot patterns displaying extraneous spots from the cubic TiC pattern...... 46

Figure 2.21. Volume change versus applied shock pressure for silicon carbide. Bars Y1 and Y2 are indicative of transformation to rock salt structure. PL1 and PL2 indicate compression of stable phases of silicon carbide, with PT1 and PT2 indicating onset and completion, respectively, of phase transformation...... 47

Figure 2.22. Approximate feed and product size ranges for common comminution devices...... 52

Figure 2.23. Compression during impact of two spheres...... 54

Figure 2.24. The electromagnetic spectrum by various units, and the analytical techniques operating in each region...... 59

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Figure 2.25. Relations between rotational, vibrational, and electronic transition energies. Figure is not to scale, as rotational transitions are typically much smaller than shown, and electronic transitions are typically much larger than shown...... 59

Figure 2.26. Expected Raman and IR spectra for boron carbide...... 61

Figure 2.27. Peak location of the boron carbide Raman active peaks at 480 and 535 cm-1 as a function of carbon content...... 64

Figure 2.28. Three-regime correlation between stoichiometry and lattice parameters as proposed by Conde et al...... 67

Figure 2.29. Correlation between lattice parameters and stoichiometry as observed by Aselage et al., depicting a linear relationship between 13 at% and 20 at% carbon...... 68

Figure 2.30. Generalized dependency of sinθ upon θ, leading to inaccurate calculations of d-spacings in lattice parameter calculations...... 69

Figure 2.31. Typical XRD pattern for boron carbide extended to 120° 2θ, with peaks for boron carbide and graphite delineated...... 70

Figure 3.1. Raman spectra of boron carbide before (a) and after (b) nanoindentation, displaying the propensity of particular wavelength lasers to excite various vibrations. 74

Figure 4.1. a) Raw ingot of boron carbide prior to separation and classification. b) Quarter-meter long segment that has been separated by hand...... 77

Figure 4.2. Benchtop Jaw Crusher...... 78

Figure 4.3. Spex SamplePrep 8000m Mixer/Mill...... 78

Figure 4.4 Sturtevant 2 in. Micronizer Laboratory scale jet mill...... 79

Figure 4.5. Malvern Mastersizer 2000 Light Scattering Particle Sizer with Scirocco Dry Powder Feeder (front left) and Hydro S Small Volume Wet Dispersion (front right) accessories...... 80

Figure 4.6. Zeiss Σigma field emission scanning electron microscope...... 81

Figure 4.7. Renishaw InVia Confocal micro-Raman Spectrometer...... 83

Figure 4.8. Mattson Galaxy Series 5000 FTIR Spectrometer...... 84

Figure 4.9. Sample FTIR spectra taken with ATR setup. No pattern is evident...... 85

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Figure 4.10. Panalytical X’Pert MPD X-Ray Diffractometer...... 87

Figure 4.11. Perkin Elmer TGA 7 apparatus...... 89

Figure 4.12. Malvern Rosand RH2000 Capillary Rheometer...... 90

Figure 5.1. Optical micrographs of commercial boron carbide powders from ESK (left) and UK Abrasives (right), prior to (top) and following (bottom) 30 minutes of low energy comminution via vibratory milling...... 94

Figure 5.2. Optical images of 0.25 m long segment of raw, unprocessed boron carbide obtained from Washington Mills. Full segment (a) was sorted by visual appearance into four sectors comprising of carbonaceous “C” (b); well reacted center band “CB” (c); large porosity “P” (d); and solid, well reacted “S” (e) sectors...... 96

Figure 5.3. SEM micrographs of sectors “C” (a), “CB” (b), “P” (c), and “S” (d) following 30 minutes of vibratory milling. Sectors “C” and “S” displayed a higher tendency for large aggregates than sectors “CB” and “P.” ...... 97

Figure 5.4. Particle size distributions of test boron carbide powders before (top) and after (bottom) jet milling at 100 psi and 5 grams per minute feed rate...... 98

Figure 5.5. Particle size distributions from the fines filter bag (top) and product collection bin (bottom) of the Micronizer. Although the distributions are remarkably similar, the fact that classification occurred and the difference in the light obscuration factors of the two chambers implies a marked contrast in composition of the powders. 99

Figure 5.6. SEM micrographs of unprocessed boron carbide following (a) one, (b) two, and (c) three rounds of jet milling, and (d) commercial boron carbide after one round of jet milling. All powders display marked presence of nanoscale material coating the larger micrometer scale particles...... 101

Figure 5.7. Typical XRD pattern for vibratory milled sample “C.” ...... 103

Figure 5.8. Typical XRD pattern for vibratory milled sample “CB.” ...... 104

Figure 5.9. Typical XRD pattern for vibratory milled sample “P.” This sample was unique in the high degree of variability of calculated graphite weight percent, averaging at 20 weight percent carbon...... 105

Figure 5.10. Typical XRD pattern of vibratory milled sample “S.” ...... 106

Figure 5.11. Typical XRD pattern of boron carbide after 1 round of jet milling...... 109

Figure 5.12. Typical XRD pattern of boron carbide after 2 rounds of jet milling...... 110

Figure 5.13. Typical XRD pattern of boron carbide after 3 rounds of jet milling...... 111

Figure 5.14. FTIR spectra of vibratory milled boron carbides. The shifts in location and intensity of the two bands at 1560 cm-1 and 1080 cm-1 can be used to roughly identify the nature and location of carbon in the boron carbide...... 113

Figure 5.15. FTIR spectra of jet milled unprocessed boron carbide powders, (a) full view and (b) expanded view. Jet milled powders were characterized as undergoing amorphization by evidence of broad, low intensity peaks near 800 cm-1...... 115

Figure 5.16. Observed (dashed lines) and calculated (solid lines) FTIR absorption spectra for several boron carbide isomorphs. The change in icosahedron structure results in a shift in an absorption band between 750 cm-1 and 800 cm-1...... 116

Figure 5.17. FTIR spectra of iteratively jet milled unprocessed boron carbide powders, (a) full view and (b) expanded view. The broad, low intensity peaks below 800 cm-1 do not increase in intensity significantly with iterative milling...... 118

Figure 5.18. FTIR spectra of commercial boron carbide powders before and after jet milling, (a) full view and (b) expanded view. After jet milling broad, low intensity peaks appear below 800 cm-1...... 119

Figure 5.19. Raman spectra for boron carbide powders from (left) ESK and (right) UK Abrasives after (a) no vibratory milling, (b) 10 minutes milling, (c) 20 minutes milling, and (d) 30 minutes milling. Both commercial powders exhibited marked increase in “D” and “G” intensity as a function of vibratory milling...... 120

Figure 5.20. Raman spectra of unprocessed boron carbide (left) before and (right) after 30 minutes of vibratory milling for sections (a) C, (b) CB, (c) P, and (d) S. Relative intensities of “D” and “G” peaks were not as expected...... 122-123

Figure 5.21. Typical Raman spectrum of unprocessed boron carbide following jet milling, clearly showing the characteristic amorphization peak at 1800 cm-1...... 125

Figure 5.22. Optical view of boron carbide particle possessing the characteristic amorphization peak used for mapping to determine concentration effect of the characteristic amorphization peak. (Left) 20x objective, (right) 50x objective...... 126

Figure 5.23. Raman spectra taken at 1 μm intervals across the sample from Figure 5.22, proceeding from left to right (a) – (l). Only select spectra exhibit the characteristic amorphization peak at 1800 cm-1, indicating the frequency of observation in jet milled unprocessed boron carbides is not a concentration effect...... 127-128

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Figure 5.24. Raman spectra for unprocessed boron carbide following iterative rounds of jet milling. The overwhelming intensity of the “D” and “G” bands of the carbonaceous materials present dampened the intensity of the boron carbide spectra to the point where it became impossible to determine if the characteristic amorphization peak was becoming pronounced or not...... 130

Figure 5.25. Raman spectra of commercial boron carbide powders before and after jet milling. The characteristic amorphization peak at 1800 cm-1 is very prominent after comminution...... 131

Figure 5.26. a) Ellingham Diagram portraying the free energies of several oxides, expanded in b) to portray the lines for reactions of carbon with oxygen as calculated from equations for Gibbs free energy. Of note is the intersection of the 2C + O2 = 2CO and 2CO + O2 = 2CO2 lines near 700°C...... 134-135

Figure 5.27. Graphitic inclusions in carbide armor materials, appearing as plate-like structures in both silicon carbide (left) and boron carbide (right)...... 136

Figure 5.28. SEM micrographs, in increasing magnification from (a) to (f), portraying clusters of nanoscale, equi-axed carbonaceous material coating the larger shard like boron carbide particles...... 137

Figure 5.29. Dependency of Vickers Hardness upon stoichiometry for boron carbide. Maxima is reached at stoichiometric B4C...... 141

Figure 5.30. TGA curves of raw and commercial boron carbides, fired to 900°C in air and nitrogen atmospheres. All samples showed significant weight gain during firing, rather than weight loss...... 142

Figure 5.31. SEM micrographs of boron carbide powders before (top) and after (bottom) firing in air to 900°C. The powders have become completely enveloped by the growth of oxide whiskers...... 144

Figure 5.32. Raman spectra for boron carbide powders during initial examination of feasibility of firing in air for removal of carbon. Note the creation of new peaks after firing in air, and the remaining carbon peaks after processing...... 145

Figure 5.33. Characteristic Raman spectrum of boric acid...... 145

Figure 5.34. Raman spectra of boron carbide powders subjected to various firing atmospheres. All atmospheres resulted in no significant selective oxidation of carbon for carbon removal...... 147

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Figure 5.35. FTIR spectra of boron carbide before and after chemical oxidation in Nazarchuk’s solution. Increase of intensity of bands near 1030 cm-1 and 2800-2900 cm-1 are indicative of growth of boron and carbon oxides, respectively...... 149

Figure 5.36. Raman spectra of several boron carbide powders subjected to chemical oxidation in a Nazarchuk solution consisting of potassium bichromate, perchloric acid, sulfuric acid, and nitric acid. The solution was ineffective at preferentially oxidizing carbonaceous materials, as evidenced by the continued presence of “D” and “G” peaks and the appearance of a characteristic boron oxide peak at 880 cm-1...... 150

Figure 5.37. Raman spectra of several polysaccharides taken with a 785 nm laser, including (a) microcrystalline cellulose, (b) methylcellulose, (c) carboxymethylcellulose, (d) hydroxypropyl cellulose, and (e) hydroxypropyl methylcellulose. Methylcellulose (b) was used as a binder in the boron carbide extrudates...... 153

Figure 5.38. Raman spectra of extrudates formed from (a) commercial 10 μm platelet shaped powders and (b) jet milled powders that had previously exhibited the characteristic 1800 cm-1 peak for amorphization. Spectra were taken at the indicated number of degrees of rotation about the axis of symmetry...... 154

Figure 5.39. SEM micrographs of boron carbide powders milled at 60 psi (left) and 80 psi (right) grinding pressure. Comminution decreases sharply with decreased grinding pressure, although faceted particle faces indicate brittle fracture still occurs at lower grinding pressures...... 157

Figure 5.40. Raman spectra of boron carbide powders milled at (a) 60 psi and (b) 80 psi grinding pressure. Both powders, despite the lowered milling pressure, still exhibited the characteristic peak at 1800 cm-1 indicative of induced amorphization...... 159

Figure 5.41. FTIR spectra taken on samples jet milled at 60 psi (a) and 80 psi (b) grinding pressure. Both samples exhibit evidence of a low peak at 800 cm-1, indicative of amorphization...... 160

Figure 5.42. Raman spectra of jet milled boron carbide powders before (a) and after (b) annealing at 600°C. The characteristic peak at 1800 cm-1 has been effectively eliminated from the spectra, and the relative intensity of the “D” and “G” peaks are markedly increased...... 163

Figure A1. Typical boron carbide diffraction pattern (top) with several relevant ICSD PDF cards showing the propensity for patterns to resemble B13C2 over B4C...... 176

Figure A2. PDF card notes for B4C card file 97-065-4971...... 177

Figure A3. PDF card notes for B4C card file 97-002-9093...... 178

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Figure A4. PDF card notes for B13C2 card file 97-000-0446. ……………………….. 179

Figure A5. PDF card notes for B13C2 card file 97-061-2568. ……………………… 180

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1. Introduction

In the course of commercial synthesis of a ceramic powder, there is inevitably a long and complicated sequence of operations to produce and refine the product from crude precursor materials to high purity, fine grained, uniform powders. This powder synthesis is known to induce structural variability and transformations in a multitude of materials; some are desirable and easily controllable, while others are detrimental and difficult to regulate. The determination of the effect of a structural transformation is dependent upon its impact upon the powder’s material properties and the bulk properties of bodies fabricated from these powders.

Boron carbide is a material that is utilized for a multitude of applications dependent upon highly favorable bulk properties, including low theoretical density, high hardness, excellent wear and abrasion resistance, and large neutron capture cross-section.

These properties have led to applications in nuclear shielding, spray nozzles, and as part of ballistic armor systems. Highly covalently bonded, boron carbide is a relatively stable structure, and that stability is paramount in several of the aforementioned applications.

Yet there is a great deal of variability in the performance of boron carbide bodies under extreme conditions which, when combined with thermodynamic and electronic considerations, cast suspicion upon the true stability of the structure. Previous works examining boron carbide have shown the two most significant structural variations present in cases of anomalous performance have been secondary free carbon and bands of amorphous material, the true origins of which have not yet been fully identified. The aforementioned suspicion has resulted in finite limits to the application of boron carbide,

2 and concern over the viability of boron carbide as a high performance structural ceramic in the future.

Commercial-scale quantities of boron carbide are often produced via a melt process for abrasive-grade applications, and as such are subject to conditions that can impart various inhomogeneities to the powder as a first source of variability. This synthesis method provides little flexibility in tuning the properties of the powder for other applications. As a result, powders derived from this method can be highly variable and when subsequently used in the processing of dense ceramics, can be the source of anomalous variability in bulk properties. This dissertation will examine whether powder variability can be the cause of the anomalous bulk ceramic property behavior.

Additionally, transformations that have been observed in boron carbide that experiences high stress conditions may be caused by the powder synthesis method as well as the subsequent processing and refining of a powder. The processing and refinement includes comminution, and in the case of a covalently bonded hard ceramic such as boron carbide, high energy and pressure comminution. It will be shown that this processing can lead to a second source of variability, a pressure-induced processing transformation. This dissertation will attempt to examine the early stages of powder processing on commercial melt-synthesized boron carbide, the extent of structural variability present before and after processing in an effort to differentiate between variability introduced via production and that introduced via processing, and the feasibility of reversing or accounting for these variations and transformations. The analytical tools of x-ray diffraction, Fourier

Transform Infrared Spectroscopy, and Raman microspectroscopy will be critical in evaluating these variations and transformations.

2. Background

This section will present background information on boron carbide, the most intriguing of observed variances in boron carbide in relation to mechanical properties, and techniques to investigate these variances. This information will be used to form a research approach in later sections.

2.1 Boron Carbide Historical Survey

Boron carbide was first noted by Wöhler1 in the mid 19th century during his work with pure boron, although purposeful synthesis of the material is credited to Joly in 18832

3 and Moissan in 1894 as each identified stoichiometric phases of B3C and B6C, respectively. Research into boron carbide was relatively limited, with Ridgway’s

4 identification of stoichiometric B4C coming in 1934. At this time interest in boron carbide intensified, particularly by Russian scientists, who began investigating and publishing research on many different boron-carbon compounds in an attempt to exactly determine the true phase diagram of the complex boron-carbon system.5, 6 In recent years, however, many of the “stable” phases identified during this period have been questioned as to their true existence as stable compounds, or whether they were merely solid solutions. In general, a solid solution phase homogeneity range is accepted as extending at least from B4.0C – B10.5C today. Starting in the 1950’s, extensive work was done into exploring the properties of boron carbide in relation to nuclear shielding and mechanical properties, and it began to be utilized in the field of structural ceramics.7-11

2.2 Production of Boron Carbide

2.2.1 Melt Processing

Critical in understanding the variability in boron carbide microstructures is an understanding of how the material is produced, as differing raw materials and a multitude of production techniques have led to a wide variety of boron carbide compositions over the years. As a general rule, boron carbide is difficult to produce. While many boron containing materials have been studied as a precursor for boron carbide, including pure boron, most commonly commercial-scale boron carbide is produced by the carbothermic reduction of boron oxide or boric acid in an electric arc furnace with a carbon source of either petroleum coke or anthracite coal.7, 8, 11-17 Typically, the arc furnace for large-scale commercial production is a modified Higgins18 arrangement. In the use of boron oxide as a precursor material, the reaction follows the general equation:

2B2O3 + 7C → B4C + 6CO (2.1)

with two stages involved:

B2O3 + 3CO → 2B + 3CO 2 (2.2)

4B + C → B4C (2.3)

The process is highly endothermic, with a ΔH of 1812 kJ/mol.7 There is also the formation of carbon monoxide to consider. It is a known fact that boron carbide produces

almost 3 kg of carbon monoxide per kilogram of product more than any other carbide in

common production.8 Compounding the difficulties of sequestering the carbon monoxide

released is the interplay of water vapor released when utilizing boric acid, H3BO3, which

can detrimentally react with the carbon present by either:

CO + H 2O → CO 2 + H 2 (2.4)

or:

C + H 2O → CO 2 + H 2 (2.5)

Not only does water vapor react with carbon, limiting the overall reaction to create boron

carbide, but boric acid is volatile in the presence of water vapor. This requires great care

to be taken when formulating furnace loads, but because of the relative ease of acquiring

boric acid it is now the most common boron containing material utilized in the production

of boron carbide.7, 8

The arc furnace for a typical boron carbide production run must be held at around

2500°C for the reaction to take place. However, it is possible to lower the required

temperature considerably by introducing magnesium to the reaction and replacing the

carbon source with either pure carbon black or a highly soluble hydrocarbon.7 The new reaction, which follows:

2B2O3 + 6Mg + C → B4C + 6MgO (2.6)

is in fact quite exothermic and does not require nearly as high a temperature to sustain;

either a hold of 1200°C or a point ignition is all that is required.8 However, the

magnesium-enhanced process has many more by-products than without magnesium, such

as magnesio-borides, magnesia, magnesium metal, and other unreacted raw materials

which require much more processing to eliminate. It is often considered too limiting to

include magnesium as a production aid.7, 8

The melt obtained during carbothermal reduction is highly viscous, and as

mentioned above, a large degree of gaseous by-product is produced. To reduce the risk of potentially dangerous slow bubbling within the viscous melt, it is common for producers to only partially fill the arc furnace at first, adding more precursor material as gases evolve and are released from the melt. Once the furnace is full and sufficient time has passed allowing for the melt to react and form the proper chemistry, the graphite electrodes of the arc furnace are removed and the melt allowed to cool into an ingot. As shown in Figure 2.1, this ingot can be several meters in diameter and consists of several zones. Due to the high viscosity and thermal properties of the melt, not all of the precursor materials will have reacted completely. At the edges of the ingot and to a degree in zones between each electrode, this unreacted material will be found, often visually resembling the coke which is its main constituent. Closer to the electrodes is well reacted material, fully formed boron carbide, with a high degree of crystallinity.

Abutting the electrodes, especially at the tips where material has been exposed to temperature for the longest time, may be a section of apparently over-reacted material that has become graphitized by leaching the electrodes.19 After removing the ingot from

the furnace, it is cracked open and the material visually sorted between boron carbide that

can then be further refined and material that needs to be recycled into the process again.

Typically, an ingot will yield less than 10% boron carbide from the precursor materials

put into the process.

bb aa aa

bb bb aa

Once the ingot has been broken and visually classified, the consumer grade boron carbide must be comminuted and cleared of as many processing impurities as possible.

Initial comminution is achieved by jaw crushing chunks from the ingot to a manageable particle size for subsequent processing. In a magnesium-assisted melt, the magnesium

containing impurities must be removed first by washing with either sulfuric acid or

hydrochloric acid, followed by a hot water wash. With or without magnesium, a heat

treatment of at least 1600°C for several hours under vacuum can be used to eliminate

residual free carbon and nitrogen, as well as any remaining magnesium if present. The

powder can then be attrition or jet milled to the final desired particle size, typically

somewhere between 0.5 – 10 µm. Milling will introduce steel impurities typically, the largest of which can be removed with magnetic separation. For final purification and removal of steel impurities, a last wash in hydrochloric acid is used.

2.2.2 Chemical Vapor Deposition

Although the vast majority of boron carbide production is performed utilizing melt processing as described above, laboratory-scale quantities of boron carbide can be produced using a variety of other methods. The most common of these laboratory methods is the chemical vapor deposition (CVD) of thin films.7, 8, 11, 20-22 A multitude of

reactions have been studied for the necessary reactions to create B4C, with the most

common formulations being either:

4BCl3(BBr3) + 2CH4 + 8H2 = 2B4C + 12HCl + 6HBr2 (2.7)

or:

4BCl3 +CCl4 + 8H2 = B4C + 16HCl (2.8)

Reaction (2.8) has been promoted as it has less by products, making it slightly

easier. Both reactions are tailorable to enable control over the stoichiometry of the

deposited film. This control results in highly predictable deposited structures, from pure

β-boron across the phase diagram through tetragonal boron-rich mixtures, rhombohedral

boron carbide, clear to pure graphite. With precise control, multiple phases can be co- deposited, and a multitude of assistive methods exist to attain precise stoichiometry control, including laser CVD, plasma enhanced CVD, and hot-filament CVD.7, 11

2.2.3 Single Crystal Production

Large single crystals of boron carbide several millimeters in length are often

prepared through extended CVD within the homogeneity range for rhombohedral boron

carbide. However, they can also be produced by zone refining of boron carbide rods,

where a rod is passed slowly through a thin furnace. The furnace will melt a small

section of the rod at a time, segregating impurities, refining to a eutectic melt

composition, and with a seed crystal located at the tip of the rod, solidifying material will

continue growth of the seed crystal resulting in a directionally grown single crystal.7

2.2.4 Rapid Carbothermal Reduction

Essentially a variant of the carbothermal reduction melt process outlined in

Section 2.2.1, Rapid Carbothermal Reduction (RCR) processes vary from standard melt processes in the heating rate of their precursor materials. While standard carbothermal reduction heats the precursor materials at rates below 100 °C/s, rapid carbothermal reduction effects heating rates of up to, and in some cases, over 105 °C/s.15, 17 This rapid

heating results in a much different reaction for the formation of boron carbide. Standard

carbothermal reduction relies upon a nucleation-growth mechanism in which the carbon

precursor reacts with boron oxide precursors as the boron oxide melts to a liquid form.

RCR kinetics were proposed to differ by a reliance on a reaction with the gas phase of the

boron oxide precursors as they are melted and evaporated too quickly for the carbon

precursors to react with the liquid form. It was then proposed that the product boron

carbide likely coalesced from the gas phase reaction.17 These RCR products were of note

in their fine and very uniform particle size centering on 1 μm, and a composition of

precise stoichiometric B4C with essentially no residual carbon. In a continuous process,

RCR should have relatively easy scalability to industrial production levels.

2.2.5 Other Production Routes

There are numerous other methods of obtaining boron carbide found in literature, many limited to small-scale laboratory quantities. Among these are methods including the reaction of polymeric precursors, which often allows for much lower production temperatures and times than a typical commercial melt process, possibly opening routes for controllable nanostructures of boron carbide, but also can result in large quantities of residual boron or carbon.23-26 Mechanical alloying has seen some investigation as a

preparative route, in which pure precursors are milled for extended lengths of time to

create an active surface, then the mixed powders are heated at a relatively low

temperature.27, 28 The extensive milling can introduce a number of impurities, however.

Liquid precursor processing, a process involving the mixture of solutions that are

naturally liquid at room temperature being autoclaved at temperatures lower than 600°C,

has shown some promise for obtaining very fine (<80 nm) particles.11, 25, 27, 28

2.3 Phase Diagram and Crystal Structure

Boron carbide’s phase diagram has been debated for many years now, with the

most recent debate sparked over what the “true” stoichiometrically stable phase is.

Generally accepted is that there is a wide range of solid solubility for carbon in the stable

phase, with a homogeneous range extending from around 10 atomic percent (at%) carbon

up to roughly 21 at% carbon. Beyond 21 at% carbon a mixture of stable phase boron

carbide and carbon is encountered, which has a eutectic point at roughly 30 at% carbon of

2350 °C, though this has also been debated to be as low as 2240 °C. Low carbon content

phases, below 10 at% carbon, are generally agreed to be a solid solution of stable phase

boron carbide and pure boron.7, 8, 11, 29, 30

As shown in Figure 2.2, a common argument is for the stable phase of boron carbide to be B13C2. While not as widely accepted as the phase diagram depicting B4C

shown in Figure 2.3, there is still a large contingent that is trying to prove B13C2 is the

true stable phase of boron carbide, while B4C is just a generalization of the overall

composition of solid solution. Samsonov5 accounted for both phases as stable in his

diagram, which then influenced Shuravlov,6 who is considered to have one of the most

complete phase diagrams assembled for the boron-carbon system. Many recent works

have also tended towards identifying B4.3C as more stoichiometrically accurate, pointing

specifically towards the work of Schwetz and Karduck in 199131 which claimed to

disprove the existence of the higher carbon content boron carbides.7 In most

communications, however, and for the purposes of this dissertation, B4C is assumed to be the true stoichiometrically stable phase.

Figure 2.2. Phase diagram after Ekbom and Amundin, depicting B13C2 as the stoichiometrically stable phase of boron carbide and presuming the presence of several low temperature phases.32

Figure 2.3. Phase diagram after Beauvy, depicting the more widely accepted B4C as the stoichiometrically stable boron carbide, with solid solutions with boron and carbon on each respective side.33

The argument over what truly is the stable phase of boron carbide can at least

partially be attributed to the crystal structure of the material. Boron carbide possesses a

covalently bonded rhombohedral crystal structure, with lattice structure belonging to the

_ 7 space group D3d5, or R3m, depending upon notation. As shown in Figure 2.4, the unit cell contains 15 atoms. The atoms are generally arranged into 12 atom cages of a near- perfect icosahedron, with 3 atoms extending in a chain from one of the equatorial positions on the icosahedron. In the case of B4C, one of the more accepted structures is

that the icosahedron is regularly filled entirely by boron atoms, with three carbon atoms

comprising the intericosahedral chain. However, due to the similarities between the

boron and carbon atoms in size and charge, as well as the inherent flexibility of the

structure to atomic impurities and substitutions,34 this atomic configuration has not been

proven to be the true configuration. A second expected arrangement is when a boron

atom substitutes into the middle site on the intericosahedral chain, with the carbon atom

then moving to an icosahedral position, with care taken to differentiate between an

equatorial site and a polar site.7, 34 The substituting boron atom can be located on any of

the chain positions however, and often a second boron atom can substitute onto the chain

as well in a terminal position. To account for this variability, the unit cell is often

denoted by writing the composition of the structure in terms of the icosahedron, followed

by the intericosahedral chain in parentheses. For any carbon atoms in the icosahedron,

the carbon is given a subscript e or p to denote its location in an equatorial or polar site on the icosahedron, respectively. The intericosahedral chain is written in order of stretching away from the icosahedron. Combined, the naming convention leads to written formulae of B12(CCC), or B11Cp(CBC).

Figure 2.4. Left: Boron carbide rhombohedral unit cell.7 Right: Equivalent hexagonal indexing depicting stacking sequencing.35

This high degree of variability allows for the wide range solid solution of boron

carbides in the homogeneous phase range, as a boron-rich or carbon-rich boron carbide

can merely have an atom replaced locally in the crystal structure without altering the

overall structure. The rhombohedral lattice parameters are a = 5.19Å andα = 66.18° .

Converted to the more easily-worked with hexagonal lattice parameters, boron carbide

7 has values of a0 = 5.61Å, c0 = 12.07Å, and an axial ratio of c0/a0 = 2.151. Due to the

difference in atomic radius between carbon and boron, boron-rich boron carbides will

have a slightly expanded lattice, and lattice parameters have been shown to vary directly

with stoichiometry of the boron carbide.10, 36-41 The true nature of the expansion process is somewhat debated. It is generally accepted that boron will substitute for carbon atoms entirely in either the icosahedra36, 40, 41 or the intericosahedral chains,37 eventually

resulting in transformation to four-atom boron clusters between icosahedra,42 before substituting into the other location. Icosahedral substitution is supported by DFT calculations, but the substitution into the chains is supported by electronic structure of boron carbide, elastic constants, and the hypotheses behind peaks in the vibrational spectra of boron carbide, and so neither theory is universally endorsed. With either substitution configuration however, the change in lattice parameters can be used to help determine boron-carbon ratio. This technique will be addressed in more detail in Section

2.9.2.

2.4 Properties and Applications

Boron carbide is most well known for its high hardness, low density, high neutron capture cross-section, and excellent resistance to heat and chemical attack, most of which

16 can be directly attributed to the highly covalent bonding. Most physical properties exhibit some degree of correlation to carbon content.

2.4.1 Chemical Properties

Boron carbide is resistant to most low temperature reagents, although hot oxidizing acids and fused salts have relative success in oxidizing boron carbide.7, 8

Chlorine reacts with boron carbide above 600°C, and bromine above 800°C, resulting in boron trihalides. Most metals will react with boron carbide above 1000°C, and high temperature interaction with metal oxides will promote the formation of metal borides and carbon monoxide.7 Aluminum and silicon both will readily react with boron carbide and are used either as sintering aids or, in higher concentrations, to form complex binary structures. Finally, boron carbide is a relatively slow oxidizer, with fine powders requiring extensive exposure to moist air to noticeably oxidize. Even at moderately elevated temperatures, up to 1000°C, oxidation rates are well below 0.002 mol/m2s, with significant increase in oxidation rates only seen above 1400°C.43

2.4.2 Density

B4C has been both observed and calculated to have a density of approximately

2.52 g/cm3. This value changes relative to overall carbon content according to:

ρ = 2.422 + 0.0048[C] (2.9)

where [C] is the carbon content in atomic percent. This linear relationship holds true

only within the homogeneous phase range of solid solution, however.7, 8

2.4.3 Elastic Properties

The nature of the rhombohedral crystal for boron carbide results in extremely

anisotropic elastic properties. However, most reported values are averaged for a bulk

polycrystalline material as tests are often conducted on sintered or hot pressed bodies.

The Young’s Modulus, E, of boron carbide can have a maximum to minimum ratio of

well over 8:1, with reported values ranging as high as 523 GPa on c = [0001] and as low

as 63 GPa in orthogonal directions. For isotropic determination on B4C, Young’s

Modulus is routinely given as approximately 470 GPa. Shear Modulus, G, is

approximately 197 GPa, and bulk modulus, K, has values around 243 GPa. Poisson’s

ratio is fairly low, even for a ceramic, with values ranging from 0.14 – 0.18, typically

being credited as 0.18.7, 8, 44-52 All elastic properties do have some dependency upon carbon, as shown in Table 2.1. In addition, elastic properties have shown strong anisotropy, as shown in Table 2.2.

Table 2.1. Elastic properties for boron carbide, calculated for single crystal and observed carbon content dependence.7, 44 Sample E (GPa) G (GPa) K (GPa) ν

Single Crystal 460.07 195.56 236.84 0.176

20.0 at% C 471 200 243 0.18

18.2 at% C 465 197 245 0.18

15.4 at% C 466 197 243 0.18

13.3 at% C 450 189 241 0.19

11.5 at% C 351 150 178 0.17

10.0 at% C 348 150 170 0.16

Table 2.2. Anisotropic elastic constants for several boron carbides. 44 51 51 Cij (GPa) B5.6C B6.5C B4C

11 542.8 500.4 561.8

33 534.5 430.2 517.7

44 164.8

12 130.6 125.3 123.6

13 63.5 73.9 69.6

14 7.7 17.8

2.4.4 Mechanical Properties

As one of the so-called “hard ceramics,” boron carbide ranks as the third overall hardest material known, behind diamond and cubic boron nitride. This impressive hardness ranks as one of boron carbide’s most important material properties and, in conjunction with its low density, is the deciding factor in many of its applications. This high hardness can be detrimental at times, as it makes processing of boron carbide extremely difficult. Processing boron carbide powders requires large amounts of energy

input, which may lead to non-uniform processing and variability in bulk ceramic properties. Still, despite being well known for having high hardness, there is a great deal of variability in reported hardness values due in part to the difficulties in comparing values obtained through any of the different accepted hardness tests. Generally, Knoop hardness is used as a reference for boron carbide, with tests under 100 gram loading

2 resulting in a value of HK100 between 2900 – 3100 kg/mm . Hardness has also shown a dependence on the stoichiometry of the boron carbide, with a maximum attained at

53, 54 stoichiometric B4C. Impressively, the hardness values experience very little change

in regards to elevated temperatures even in excess of 1500°C, making it one of the best

structural ceramics at high temperatures.7, 8, 45, 47, 49, 50, 52, 54-67

Bending or flexural strength of a material follows a direct relation with hardness,

and thus boron carbide is one of the higher bending strength materials with values of σ

commonly reported between 350 MPa and 450 MPa, and sometimes in excess of 450

MPa.8, 47, 49, 59, 68-73 Similar to hardness, bending strength maintains well with elevated

temperatures, with little noticeable reduction up to 1200°C. At extreme temperatures up

to 2000°C boron carbide still experiences little loss of bending strength and outperforms

similar carbides. Surface oxidation can be a problem above 1200°C in air. The

combination of strength and density makes boron carbide one of the most attractive

structural materials known, as shown in the comparison in Table 2.3.8

Table 2.3. Comparison of high strength materials per their strength/density ratio.8, 74-77 Material Bending Strength Density (g/cm3) Strength/Density (kg/mm2) Ratio

Boron Carbide 40 2.5 15.9

WC-Co 94/6 170 14.9 11.4

Silicon Carbide 35 3.2 10.9

Aluminum Nitride 33 3.3 10.0

Aluminum Oxide 35 3.9 9.0

MgAl2O4 Spinel 31 3.6 8.5

AlON 31 3.65 8.4

Steel 60 7.8 7.7

Titanium Carbide 35 4.9 7.2

Of note in both hardness and strength of boron carbide is the anecdotally observed

occasional lack of strict adherence to the Hall-Petch relation in published data. Hall-

Petch predicts that as a material’s grain size decreases, there will be a corresponding

increase in yield strength and hardness.78 In general, this is due to dislocation buildup effects. When a material is strained, dislocations slip to accommodate the strain.

However, grain boundaries require more energy for a dislocation to propagate across than solid grains as the dislocation must change direction in order to follow the orientation of the new grain. This increased energy means higher stresses are required for dislocation movement. When grain size is reduced, the ratio of grain boundary to grain increases, resulting in dislocations encountering grain boundaries more often. The smaller grains then require more stress for dislocation movement and yielding, which is mathematically described in the Hall-Petch Equation:

k y σ y = σ 0 + (2.10) d

where σy is the yield stress, σ0 is the material property stress required to initiate

dislocation movement, ky is the material property strengthening coefficient, and d is the

diameter of the grains. Although Hall-Petch predicts a material with infinitely small grains will have infinite strength, below a critical grain size, on the order of 100 nm, strength and hardness begin decreasing in an effect that has come to be known as Inverse

Hall-Petch behavior. Hall-Petch behavior has been proven in nearly every single material known, making it as near to a proven law as is possible. However, boron carbide has, on occasion, shown little to no correspondence with the Hall-Petch equation; hardness and yield strength values show variation well below expected dependencies over grain sizes ranging from 5 – 1000 µm. This is extremely anomalous in the world of material science and implies that there may be a fracture mechanism in play that is prohibiting boron carbide from following fracture behaviors generally observed in other materials. This may arise from the variability seen in the microstructures, particularly in the overwhelming presence of secondary carbonaceous materials.

Boron carbide has a relatively low fracture toughness. Values of KIC for boron

carbide are given as varying between 3.1 and 4.1 MPa m1/2,7, 57, 60, 64-66, 69, 72, 73, 79 with

higher values possible dependent upon synthesis method, up to 5.0 MPa m1/2.53 Fracture

toughness also remains fairly constant at elevated temperatures. Experiments up to

1200°C resulted in minimal drop in fracture toughness despite surface flaws created by

slow oxidation.8

Boron carbide possesses exceptional wear resistance, and ranks second behind

diamond, ahead of silicon carbide, zirconium oxide, and aluminum oxide.67, 80-85

Arbitrary unit values of wear resistance list diamond as 0.613, boron carbide as 0.422,

and silicon carbide as 0.314. Wear resistance and coefficient of friction do decrease for

boron carbide at elevated temperatures as boron oxide and boric acid form on the

surface.8

Mechanical properties can be expected to have direct relationships with

composition of the boron carbide. By calculations, boron carbide is expected to have

increasing hardness and strength as carbon content increases through the phase

homogeneity range.53 Beyond the solid solution range towards higher carbon content,

secondary graphite exists in the structure, lowering hardness and strength. Fracture toughness increases with carbon content through the phase homogeneity region, and lowers with increasing carbon as free graphite serves as stress concentrators and crack nucleators.

2.4.5 Electronic Structure and Properties

The electronic structure and properties of boron carbide have been slow to be fully understood, despite the strong relation to bonding and other properties. Part of the difficulty lay in the fact that although boron carbide was experimentally determined to be a p-type semiconductor, calculations showed that B13C2 should be the most stable phase

and metallically bonded.34, 86-90 The coexistence of the many isomorphs of boron carbide

helps to explain this phenomenon. The isomorphs also serve to increase the difficulty in

pinpointing transition energies and assigning them to specific bonds. However, by

combining results from experiments in x-ray Raman scattering (wherein high energy x- rays inelastically scatter from a material, exciting electronic transitions) and

photoluminescence with previous work in optical absorption and high temperature

electrical conductivity, and using ab-initio modeling to fill in the missing details, the

energy band structure displayed in Figure 2.5 was determined. This structure details a

band gap in boron carbide of 2.09 eV.86 The electrical conductivity of boron carbide was

found to be less than 1 S cm-1.73, 91

Figure 2.5. Calculated band structure for boron carbide. Left axis gives energies relative to valence band, right axis gives energies relative to conduction band.86

2.4.6 Nuclear Properties

Boron carbide is highly noted for its high neutron capture cross section. The

isotope 10B has an extraordinarily large cross section of roughly 4000 barns.7, 8, 11 A barn

is defined as the effective cross sectional area an atom poses to a fast thermal neutron,

such as those released from nuclear reactors and warheads. Possessing true units of 10-24

cm2, the barn is essentially a measure of the probability a thermal neutron will encounter

an atom and be slowed. There are very few elements with a higher capture cross section,

24 and most of those are of the heavy rare earth series of elements. Boron carbide, while 80 at% boron, is only about 14.7 at% 10B, resulting in a total neutron capture cross section for boron carbide of 600 – 700 barns. By comparison, lead has been shown to have a neutron capture cross section of only 0.01 – 0.1 barns.92, 93 If desired, boron carbide can be enriched to have up to 90 at% 10B, theoretically.7, 8

2.4.7 Applications

One of the largest markets for boron carbide is in nuclear ceramics; the high neutron cross capture cross section, low theoretical density, and excellent high temperature properties allow for use in nuclear reactors as a shielding material. The most prevalent application of boron carbide purely for its mechanical properties is as an abrasive material. The high hardness makes boron carbide ideal for use in lapping pastes and polishing wheels as it can polish most hard materials and is significantly cheaper than diamond. After abrasives, the next leading market is in wear resistant parts such as spray nozzles for sand blasting or water jet cutting. The exceptional hardness and low density also make boron carbide an extremely attractive material for use as an armor ceramic, and is used for personnel armor and in helicopter seats. The chemical reactions of boron carbide are also used to boronize metals, resulting in a harder surface, by depositing a thin film of metal borides on the surface.

2.5 Common Inhomogeneities in Boron Carbide

Boron carbide can be subject to multiple inhomogeneities in the form of impurities, generally introduced through the melt process depicted in Section 2.2.1 and

subsequent post-processing to reduce the ingot to fine grained powder. These impurities

are usually limited to three main materials in quantities detectable by bulk

characterization such as x-ray diffraction or flame chromatography. Iron based

impurities from milling are usually present at less than 0.2 weight percent (wt%) of the powder. Introduced during milling of the ingot, these impurities are almost entirely eliminated by magnetic separation and acid leaching.19 Also present will be boron oxide,

present partly due to the natural slow oxidation of boron carbide mentioned in Section

2.4.1, and partly because of remnant unreacted precursor material that has not been fully

separated from the processed boron carbide.7, 8, 19 Typically, boron oxide will constitute

0.5 – 2 wt% of the powder. Finally, there will always be around 2 – 3 wt% excess free

carbon present in boron carbide powders. It is commonly assumed that this carbon is

solely the result of unreacted precursor material not being sorted out of the salable boron

carbide, similar to boron oxide impurities. These carbon impurities do play a noticeable

role in bulk microstructures, where even more carbon is introduced in the form of sintering aids for densification. The total excess free carbon often results in so-called

“dirty” microstructures, as shown in Figure 2.6. The true nature and impact of these carbon impurities have yet to be fully examined, however.

Figure 2.6. Dense body microstructures of boron carbide. Excess carbon impurities are readily visible as the bright white locations throughout the microstructure.69

2.6 Amorphization of Boron Carbide under High Stress Conditions

From Section 2.4 it can be understood that nearly all of boron carbide’s

commercial applications have been determined, at least in part, by its exceptional

mechanical properties. However, there has been evidence that under extreme operational

conditions, boron carbide can undergo a phase transition with drastic effects upon the

mechanical properties. This section will examine two of those operational conditions and

recent work done to shed light on this phenomenon.

2.6.1 Ballistic Impact

As stated earlier, one of the main applications of boron carbide is as an armor

ceramic component to a complete armor package. The low density and excellent

hardness, strength, and wear resistance make boron carbide one of the most ideal

materials for ballistic protection in terms of penetrator defeat. An armor system is quite

complex and the processes that occur during ballistic impact have been shown through

in-depth studies to be many and multi-scaled as a result of the high stress and strain rate

involved. However, a brief note on penetrator defeat is appropriate at this time.

When a penetrator strikes the face of an armor ceramic, the kinetic energy of the

penetrator results in significant stress being placed upon both the ceramic and the

penetrator. During this stage, known as dwell, the penetrator deforms under the applied

stress and begins to flow laterally across the face of the ceramic, losing significant

amounts of kinetic energy and beginning to erode. Dwell, also known as interface defeat,

will last as long as the ceramic is confined in place and the penetrator cannot move

beyond the surface. However, the ceramic is also under stress, which travels as a shock

wave through the ceramic, initiating several types of cracking nearly simultaneously.

The impact creates a bending moment in the ceramic, resulting in tensile stresses on the

back surface that initiate radial cracks that travel from the back surface to the front. Part

of the shock wave is reflected from the back surface as a tensile wave, which, once it

reaches the front surface of the ceramic, will initiate Hertzian cracking at flaws near the

point of impact. As the cracks propagate, they form a conoid of damage extending out

from the impact site and widening as it extends to the back surface. This damage

separates the ceramic from the backing plate, allowing the ceramic to deform plastically.

Cracks coalesce, forming fragments of the ceramic of varying size, and the deformation

causes the fragments to move. The penetrator can then slide past the fragments, and dwell has ended. As the penetrator moves deeper into the ceramic, the fragments are

pushed around the penetrator and out of the path of its travel, abrading it and causing

great erosion, until either the penetrator is eroded completely or the back face of the

ceramic has been penetrated.94-96

Boron carbide then is a good candidate for penetrator defeat as its high fracture

strength should allow for a relatively long dwell time. More importantly, the exceptional

hardness serves well to erode the penetrator without excessive damage to the ceramic

itself. The low theoretical density also makes boron carbide attractive for reducing the

armor system weight when compared with traditional metal based armors. However, in

practice, boron carbide begins to follow anomalous fracture behavior at stresses

approaching the Hugoniot Elastic Limit (HEL), the point at which a stiff material begins

to deform plastically.70 This value is near 20 GPa for boron carbide, and near this value,

boron carbide does indeed experience wide spread dislocation formation, movement of

stacking faults, and microtwinning.71, 97 However, once the HEL is breached, rather than

plastically deform, boron carbide experiences brittle fracture of a nature akin to that of a

glass fracturing, as depicted in Figure 2.7.98 High resolution electron microscopy

(HREM) has shown a tendency for amorphous bands to be present in the microstructure near locations of ballistic impact, as shown in Figure 2.8.99 At the time, it was thought

that these amorphous bands were being created by the ballistic shock wave. The bands

coincide with the (113) plane of the unit cell, which correlates with the plane containing the intericosahedral chains.99 The occurrence of these amorphous bands was not very

common, and they were difficult to locate in ballistic rubble. While the inherent

difficulties in collecting and analyzing ballistic rubble may be responsible for the rate of

observation of the amorphous bands, the phenomenon ignited a debate over whether

29 amorphization could be an influential factor in the anomalous high strain rate ballistic performance.

Figure 2.8. Left: Optical image of ballistically impacted boron carbide. Right: HREM image taken close to the point of impact, showing amorphous band running through the sample.99

2.6.2 Nanoindentation

Ballistic impact is not the only high stress condition in which there has been

notable evidence of an amorphous transition in boron carbide. Nanoindentation studies

have repeatedly shown a phase transition when examining the sample before and after

testing. In nanoindentation, oftentimes either a spherical or Berkovich diamond indenter

is utilized to apply loads between 10 – 300 mN. Contact areas with the indenter, and thus

applied pressures, are directly dependent upon the hardness of the material being tested.

In the case of boron carbide, the high hardness results in applied pressures being very

large, on the order of 40 GPa, well above the stress observed through ballistic testing to

induce a change in performance. The recorded applied load and displacement into the

sample allow calculation of several mechanical properties, including hardness and

fracture toughness. When crystallite size is large enough, nanoindentation also allows for

examination of single crystal properties within a polycrystalline bulk, as the tested area is

significantly smaller than the grain size.

Interestingly, load-displacement curves for boron carbide do not show any

discontinuities that would be indicative of a phase transition. The altered volume and

mechanical properties arising from a switch in atomic bonding should have a noticeable

effect upon the ability of the indenter to displace material, resulting in a sudden shift in

the load-displacement curve.100 The lack of such evidence does not necessarily mean the phase transition does not occur; rather, it could mean that the volumetric change associated with the transformation is negligible, as may be expected with an amorphous band of 3 nm or less. The lack of a discontinuity in the load-displacement curve could

also arise if the amorphization occurs during the initial loading or final unloading of the

31 sample, when the change in applied pressure is high and discontinuities in the load- displacement curve are difficult to track at best. When characterized through TEM and

Raman microspectroscopy, among other techniques, however, it becomes evident that structural changes are indeed occurring. First noted was the formation of deformation shear bands along the <113> direction, implying stress relief mechanisms occur initially within the same plane as noted during ballistic impact. Further examination of indents formed under 100 mN display the formation of amorphous bands along (113) and (003) bands, as shown in Figure 2.9.101 The (113) amorphous bands are prevalent under indentation locations, supporting the theory that pressure-induced amorphous transformations are occurring in boron carbide. The amorphous bands then can be postulated to lead to premature failure in boron carbide, as Askenazy et al. showed amorphization to be the onset of spall in materials under high pressure.102

Further proof was supplied by examining the Raman spectra of boron carbide,

which are characterized by a series of bands extending from 200 to 1200 cm-1 (Figure

2.10), assigned in the literature to the vibrations of the principal structural elements in boron carbide, the icosahedra and the three-atom linear chains.103 These bands and the

Raman spectra of boron carbide will be explored in greater length in Section 2.9.1.

Examination of the Raman spectra obtained from various surfaces affected by contact

loading (Figure 2.10b-d) reveals significant structural changes that occur in boron carbide

during indentation or scratching.101 The most notable is the appearance of several new

broad bands at higher frequencies, in particular the most prominent band centered around

1330 cm-1. The fact that the Raman spectra of indentations made in single crystals and

polycrystalline material, as well as the spectra of scratches and scratch debris, have

similar high-frequency bands suggests the similarity of the microstructural changes in the

material under various contact loading situations (Figure 2.10). It should be noted also

that the spectra shown in Figure 2.10b-d have been acquired under low intensity of the incident laser beam, in order to avoid artifacts due to laser heating. When the intense laser beam was used for investigation of the scratch debris, increased temperature on the analyzed surfaces resulted in the appearance of characteristic features of graphitic/amorphous/disordered sp2 carbon, evidenced by the Raman bands at 1350 and

1590 cm-1 (the so-called D and G bands), as shown in Figure 2.10e.

(a) - pristine boron carbide (b) - indent on single crystal

(d)

Intensity (a.u.) Intensity (c) 1330

1520 (b) 1085 478 1000 720 530 830 (a)

200 400 600 800 1000 1200 1400 1600 1800 Raman Shift (cm-1)

2.6.3 Exploration of Amorphization

The observed localized amorphization occurring in boron carbide bodies placed

under high stress conditions has led to an increased effort to explain just how and why

such a transformation is taking place. The result has been an increase in publications and

significant research with several key revelations that have begun to illuminate the origin

of amorphization. This section will begin to address these revelations.

Boron carbide was deposited by DC-magnetron sputtering of hot pressed targets

in an attempt to form both crystalline and amorphous films and compare the structure and

properties resulting from sputtering. It was noted that the amorphous boron carbide films

that were deposited possessed a Raman spectra that was indicative of a true amorphous

material, with one large broad hump across the majority of the spectra, as shown in

Figure 2.11.50 This contrasts sharply with the stress-induced amorphous materials, which had been noted as showing only local effects in the D and G bands. While it can be assumed that the amorphization in stressed materials is so locally confined as to not be able to contribute significantly to an amorphous hump above the background noise in

Raman spectra, the sometimes drastic changes in intensity of D and G bands suggests that there is a marked transformation occurring that is not represented by fully amorphous boron carbide. From these results, a first hypothesis can be formed that the stress- induced amorphization may not be identical to that displayed by a truly amorphous boron carbide material.

In response to this puzzling revelation, work was conducted to try to determine

exactly what was occurring in the amorphization of stressed materials.104 Molecular

dynamics simulations were run on several different possible formations of boron carbide,

ranging from the assumed B11C(C-B-C) polycrystalline material from experimental

mechanical testing and the amorphous boron carbides from deposition experiments. The

models predicted that stress induced amorphization would occur by breaking the

intericosahedral chains, resulting in disordered but intact icosahedra bonded by a boron-

carbon amorphous film. This film was also shown to have contained small chunks of free

carbon resulting from the broken chains bonding together.105 When boron carbides were dynamically indented at strain rates approaching 103 s-1, the observed intensifying of the

D and G peaks was markedly more pronounced than in statically indented materials at similar load values.

Raman spectra taken of the dynamically indented samples demonstrated a trend towards creating amorphous carbon clusters. Fourier transform infrared spectroscopy

(FTIR) was utilized to complement the Raman spectra. The spectra, shown in Figure

2.12, displayed a trend in dynamically indented boron carbide for the formation of amorphous boron, as shown by the broad peak at 760-830 cm-1. This peak, when sharp, is attributed to B12 icosahedra, and when broadened is assumed to have been induced by amorphized boron.106 These results combine to form a second hypothesis that amorphization in stressed boron carbide occurs locally in the formation of amorphous carbon and boron clusters. This combines with the first hypothesis to create a possibility of two distinct “phases” of amorphous boron carbide, wherein one contains clusters of boron icosahedra and carbon rings randomly associated, while the other is a pure disordered random arrangement of boron and carbon atoms.

Figure 2.12. FTIR spectra of boron carbides displaying the formation of the broad amorphous boron peak under dynamic indentation.106

Recent work in loading and unloading of boron carbide materials under high

stress has led to a significant third hypothesis concerning the nature of amorphization in

boron carbide. These loading and unloading tests utilized both hydrostatic stress, where a

low strength pressure transmitting medium surrounds the sample and is pressed on

uniaxially, transmitting compressive stress across all surfaces of the sample, and

nonhydrostatic stress, where a high strength pressure transmitting medium surrounds the

sample and the applied uniaxial stress is transmitted as a uniaxial compressive stress to

the sample. Boron carbide was subjected to hydrostatic loading under pressures up to 50

GPa and subsequent unloading, with Raman spectra being recorded throughout the

testing via the use of a diamond anvil cell. The resultant Raman spectra indicated a near

complete lack of localized amorphization during the entirety of the test. Nonhydrostatic

tests, however, displayed the appearance and increase in signal of amorphous boron

carbide peaks upon unloading, at stresses approximating 16 GPa, as shown in Figure

2.13.107 This lack of amorphization under hydrostatic conditions and appearance only under unloading in nonhydrostatic conditions leads to the somewhat puzzling conclusion that the crystal structure is not actually breaking down under loading, but rather during pressure unloading.

Figure 2.13. Raman spectra taken during nonhydrostatic unloading of boron carbide, indicating rise of amorphous peak at ~1810 cm-1at approximately 16 GPa.107

These hypotheses have yet to be tested in great detail and are by no means assumed to be completely correct. One important note is that modeling of compression of boron carbide has previously predicted phase transformations at pressures lower than those experienced by experimental samples.104 However, none of these samples have yet shown a propensity for the transformations predicted. It can thus be inferred that even with these further hypotheses on the nature of localized amorphization in boron carbide, not all aspects of the process have been identified as of yet.

2.7 Phase Transformations induced by Comminution and Mechanical Working

Boron carbide is not alone in undergoing a phase transformation under high stress

conditions. In fact, many materials have been shown to experience the phase transition

even at stresses experienced during mechanical comminution, and are discussed in this

section.

2.7.1 Boron

Elemental boron and its derivative compounds have long been investigated as

promising structural and electronic material with some of the most intriguing properties

known.9 In work with mechanical comminution of boron, it was shown that structural changes were readily occurring. An in-depth x-ray diffraction study was conducted to determine the nature of the comminution-induced structural transformations. As seen in

Figure 2.14, the x-ray diffraction patterns indicate a clear shift from narrow, identifiable peaks prior to comminution to one large, broad peak after comminution.108 This change

is clearly indicative of large-scale amorphization of the boron, and serves to display that

bulk boron readily undergoes a phase transformation under comminution.

2.7.2 Carbon

As with boron, carbon is a material that has been studied extensively in multiple fields for its remarkable properties. Several of these studies have shown that carbon has a great multitude of polymorphs, with amorphous carbon occurring as a fairly common polymorph. In studies on the milling conditions required for amorphization to occur in several non-metallic species, it was found that graphite, in both natural and synthetic forms, displayed a tendency to amorphize under high energy multi-axis vibratory milling.109 Pure natural graphite was shown to have minimal amorphization under the constraints of the experiment, but a synthetic graphite showed a marked transition from crystalline structure prior to milling to a mostly amorphous structure after milling, as shown in Figure 2.15. Kosmac and Courtney were suspicious of this behavior, and suspected it to be resultant of the impurities in the natural graphite.109 To corroborate, they repeated the experiments with high-purity natural and synthetic graphite, which then both displayed a tendency to amorphize. Of note is also that the graphites used in the

41 second series of experiments reacted with the steel milling media strongly enough to result in easily detectable levels of iron carbide by x-ray diffraction.

Figure 2.15. TEM showing crystal structure of pure synthetic graphite before (top) and after (bottom) vibratory milling, with amorphous structures dominating the post-milling material.109

2.7.3 Boron Nitride

Compounds of boron have also been shown to exhibit amorphization under mechanical comminution. Specifically, hexagonal boron nitride has displayed a remarkable propensity for amorphization when ball milled, either for comminution or for attempts at mechanical alloying, a process which provides energy for reaction of two materials through co-milling of the materials. On its own, hexagonal boron nitride (h-BN) was milled by Huang et al. for increasing amounts of time in an attempt to study the microstructural effects of comminution.110 By 180 hours of ball milling, the h-BN had devolved to an entirely amorphous microstructure, as illustrated by the electron

microscopy images shown in Figure 2.16. When this work was expanded to examine the

mechanical alloying of boron nitride with carbon in an attempt to form new alloys of the boron-carbon-nitrogen (BCN) system, it was found that after only 60 hours of ball milling the structure shown in Figure 2.17 appeared, one that was almost entirely amorphous.111 The minimal crystalline material that formed was essentially just a few

basal planes forming of hexagonal product, but even these “nanocrystals” were rare.

Thus, a low-energy ball milling of time scale not atypical to industrial processes resulted

in amorphization of boron nitride.

Figure 2.16. HREM of h-BN ball milled for 180 hours, showing distinctly amorphous microstructure. Diffraction pattern insert shows only amorphous rings.110

2.7.4 Metals

While comminution induced amorphization has been shown to occur in boron

carbide’s pure constituents, the phenomenon has been most extensively studied in the

realm of metallurgy. Over the years, several metals have been shown to have extensive

amorphization due to processing induced damage, both from comminution and from cold

and hot working. In work by Stevulova et al., iron and silicon comminution and

structural changes were examined as the initial stages of mechanical alloying.112 Figure

2.18 shows the x-ray diffraction peak intensity of the iron and silicon phases as planetary

milling time increases, displaying a marked drop in intensity, which, combined with the

associated peak broadening, indicates amorphization of the metals. Similarly, Huang et

al. studied several metals under conditions of mechanical alloying, including extensive

microstructural evaluations of pure copper during comminution.113 Figure 2.19 displays the TEM micrographs of this milled copper and the resultant dislocations and twins that begin to appear as the energy imparted by ball milling induces deformation mechanisms

44 in the metal. This deformation was shown to be part of the onset of amorphization in the comminuted metal.

2.7.5 Carbides

As with the materials discussed so far, the carbide family of materials is not immune to damage induced by mechanical processing. However, while many examples exist of amorphous carbides and phase transformed carbides, and comminution of carbides, there is little literature on comminution of carbides leading to amorphous or phase transformed materials. Titanium carbide (TiC) has been shown to undergo impressive microstructural transformations when polished.114 Figure 2.20 shows the microstructure of TiC after having been diamond polished to a mirror finish. The actual micrograph displays little obvious microstructural damage from polishing, though the

edges of the scratches are unnaturally smooth. The true indication of a microstructural

change is shown in the diffraction spot patterns, which exhibit miscellaneous spots not

indicative of the cubic TiC crystal. These spots were shown to be representative of the

6H polytype of TiC, and suggest that the induced damage from mechanically working the

TiC created a build-up of dislocations and twins such that the resultant surface

microstructure was essentially a thin phase transformation.

Silicon carbide has also been shown to exhibit interesting deformation mechanisms and phase transformations under conditions of mechanical work and high pressure. Early work has shown that silicon carbide reacts to conditions of stress with formation of twins and dislocations preferentially and is difficult to induce phase transformations.115 However, under conditions of high shock pressures, as shown by

Sekine et al.,116 in Figure 2.21, the volume change induced during shock up to 160 GPa is

indicative of a severe phase transformation. The first calculated values for the

transformation, represented by bars C1 and C2 in Figure 2.21, are a bit below the

observed values of Y1 and Y2. Also, while the Y1 and Y2 values are for an observed 3C

SiC, the detailed strain curve containing PL1, PT1, PT2, and PL2 are for 6H SiC, and are indicative of compression of stable 6H SiC at PL1, onset and completion of transformation at PT1 and PT2, respectively, and compression of the new, rock-salt like structure during PL2. While not an amorphization event, this shows that even strongly bonded carbides have been observed to undergo phase transformation under conditions of shock pressurization, similar in essence to comminution.

Finally, there has been minimal research done in the past into the effect of

comminution upon the microstructure and resultant mechanical properties of boron

carbide. A study by Tkachenko et al.65 has shown that commercial boron carbide

powders can be significantly affected by planetary milling and explosive shock wave

treatment, a process designed to activate pressure-activated flaws by imparting an

explosive pressure well above the HEL through a plastic medium to particles suspended

in the plastic medium. The medium serves to prevent densification as well as impart

pressure. While not examining the pure occurrence of the amorphization event,

Tkachenko’s work does investigate the formation of twins and dislocations following the

grinding and shock treatment. The powders, as shown in Table 2.4, have markedly

increased dislocation and twin densities when studied through TEM versus the pristine

powders. Even the samples hot pressed from pure powder and shocked powders display

this same trend, indicating that the damage induced by shock treatment may not be

reversible under typical sintering conditions for this boron carbide powder. While this is a notable study, it does not examine the realm of pressures between planetary milling and explosive shock treatment, such as that induced by jet milling. Also, the work admits it lacks true mechanical property results due to highly variable densities in the powders, indicating a need yet for the mechanical effect of these process induced defects to be fully explored.

2.8 Comminution

Comminution is defined as the process of reducing objects into fine, smaller particles, which in the materials world is achieved by energetically impacting them in various mills. The impact material and mechanism of driving impact can vary greatly between the different forms of comminution, resulting in highly variable energy requirements and product sizing. The main principle and mathematics, however, are the same throughout. The common milling equipment discussed in this dissertation includes jaw crushers, ball mills, shaker mills, attrition mills, and fluid energy mills.

Jaw crushers reduce large rock sized particles to finer feed stock for mills by pinning the particles between two plates and compressing the particles by impact until they break apart into 5 mm or less particles. Ball mills are often called rotating mills, and are based on hollow cylinders rotating about a central axis, with charge of hard, wear resistant media, which come in a variety of shapes and sizes, usually in cylindrical or spherical orientation, and coarse powders. The rotating action results in a cascade effect

of the material inside the mill cylinder, dropping powder and media onto the material

lower down the side of the mill. Shaker mills are a variation on the ball mill design,

where a similar loading pattern is used, but rather than the mill cylinder being rotated, in

shaker mills the mill cylinder is energetically shaken about multiple axes, resulting in

higher energy impacts between materials. Attrition mills suspend the powders and media in a liquid medium, then agitate the entire mixture with rotating arms that stir the mixture at a high velocity. The highest energy mill is the fluid energy mill, commonly known as

a jet mill.

Jet mills operate with no media, but rather accelerate relatively fine feed powders

on a high pressure jet of gas, commonly dry, clean air. The accelerated powders then

enter a round chamber where gas jets, flowing either opposite to or in tandem with the

flow of the feed air, create a vortex where the powder particles impact one another at

high velocities. Cyclonic action is then used to classify out the much finer product.

Because of the high energies involved and relative cleanliness of a process that involves

no media and little interaction with the grinding chamber, thus limiting chances of

contamination, jet milling is commonly used to reduce powders to micron and sub-

micron particles sizes.

In all cases, the product size is dependent upon the energy of the impact between

particles, as the kinetic energy of the particles is directly used for breaking of the powders

into smaller pieces. In general, the performance of the mill based on this energy is then

given by the somewhat vague relationship:

Particles Generated Media Collisions Particle impacts Particles = ∗ ∗ (2.11) Time Time Collision impact

The actual efficiency can then be described generally by:

 1 1  U T = Ac  −  (2.12)  m m  a a0 

where UT is the total energy required to produce one unit of product, Ac is the efficiency

constant for each particular mill, a and a0 are the product and initial particle sizes, and m

is the fracture constant for the material being milled. This can all combine to roughly

create Figure 2.22, which shows the comparative mean feed and product sizes for some

of the various industrial comminution equipment, illustrating quite well the range of

processes necessary to reduce a large ingot to micrometer or finer powders.117-122

Figure 2.22. Approximate feed and product size ranges for common comminution devices.117

This is only a range of product sizes, however, and does not show the true energy

involved in each milling scenario. To truly compare comminution to high pressure

situations such as ballistic impact and nanoindentation, the energy of impact must first be

calculated. In general, as stated previously in Equation (2.11), it is assumed the full

kinetic energy of the impacting particles is expended towards fracture; implying impact energy Ei is equal to the kinetic energy Ek of the impacting particle,

1 E = E = mv 2 (2.13) i k 2

In typical rotational ball mills, the rotational velocity of the media can be

calculated and applied to Equation (2.13). However, in a jet mill, the scenario becomes a

bit more complicated due to the nature of the cyclonic action in the mill. The discussion on amorphization in boron carbide from Section 2.6 notes that pressure is of critical interest in determining amorphization. For the purposes of this dissertation, then, it is more convenient to address the contact pressure of impact in a jet mill, assuming the following:

2.1) There are no losses of energy due to friction in the system.

2.2) Air flow in the jet mill is laminar.

2.3) Particles in the jet mill are spherical and uniform in size.

2.4) Particles in the jet mill are only acted upon by the air pressure within the jet

mill prior to impact with another particle.

2.5) At the moment of impact, the particles are on a level plane traveling at

velocities of equal and opposite magnitude due to equivalent air pressure

acting on each sphere.

2.6) During impact, both spheres compress an equal amount.

2.7) Shear forces are unaccounted for in the simplified model.

The contact pressure can then be calculated from models for the impact between two spheres, as calculated by Puttock and Thwaite.123 This model accounts for elastic

compression of the spheres, with impact creating a compressed surface on each sphere, as

shown in Figure 2.23.

Figure 2.23. Compression during impact of two spheres.123

The combined compression of both spheres is a length designated α. When considering the spheres’ materials properties of Poisson’s ratio, υ, and Young’s Modulus, E, with the force being applied externally on the spheres, F, the compression α between two spheres of the same material of diameter D1 and D2, respectively, is then found by

1 2 1 2 2 3   3   3  9  1−ν 3 1 1 α =   *  * F * +  (2.14)  2   E   D1 D2 

If the first sphere has radius r1, then the right triangle in Figure 2.23 can be formed with hypotenuse r1 and legs r2 and r3. r2 and r3 are calculated by

α r = r − (2.15) 2 1 2

and, by Pythagorean Theorem,

2 2 2 r3 = r1 − r2 (2.16)

Substituting Equation 2.15 into Equation 2.16 gives

1 r 2 = αr − α 2 (2.17) 3 1 4

The contact area between the two spheres then is a circle with radius r3. Impact pressure

Pi is then

FAir Pi = (2.18) AContact

and the force applied by the air pressure, FAir, can be calculated from

FAir PAir = (2.19) ACrossSection

where ACrossSection is the cross sectional area of the particle being acted upon by the air pressure, PAir, with diameter D1. Combining Equations 2.19, 2.18, and 2.17,

PAir * ACrossSection Pi = (2.20) AContact

P *πr 2 P = Air 1 (2.21) i π απr 2 − α 2 1 4

Table 2.5 displays calculated impact pressures between two spheres of boron carbide of equivalent diameters of 45 μm with 100 psi of air pressure driving the particles. However, particles are very rarely perfect spheres, and often have surface roughness that creates an effective geometry much closer to point on plane impact rather than sphere on sphere impact. To approach this geometry, Table 2.5 also shows calculated impact pressures for situations in which a surface protrusion exists upon the impacting particle. The force of impact does not change as the air pressure is still being exerted upon a particle of equivalent diameter to the impacted particle. Only the contact pressure increases as the effective contact area decreases. Of note is the quick increase in contact pressure with reduced equivalent geometry, well above the threshold limit of 15 –

20 GPa mentioned in Section 2.6 as required for inducing amorphization. The levels of contact pressure calculated well exceed the fracture strength of boron carbide, however these calculations are for instantaneous pressures in an idealized system and thus are useful as a reference for maximum possible contact pressure.

Table 2.5. Calculated impact pressures for various effective geometries during impact of two spheres of boron carbide. D2 is the diameter of the impacted particle and D1 is the effective diameter of a surface protrusion on the impacting particle. D1 D2 Contact Area Contact Pressure (μm) (μm) (m2) (GPa) 45 45 7.11E-13 1.54 5 45 1.35E-13 8.13 0.5 45 2.72E-14 40.31 0.05 45 1.38E-15 791.97

2.9 Analytical Techniques

Over the course of this dissertation there must be techniques that can interpret the proposed structural changes and determine whether or not they are occurring during typical powder processing. Two main techniques of vibrational spectroscopies and x-ray diffraction are discussed below.

2.9.1 Vibrational Spectroscopy

Vibrational spectroscopy is essentially the process of using an incident photon to excite vibrations in a molecule, which can then be studied by analyzing the photons emitted from the molecule in turn. When examining the energy of a molecule, it can be written that

Etotal = Eel + Evib + Erot (2.22)

where the three components correspond to the energy associated with the motion of electrons, the vibrations of the atoms, and the rotation of the entire molecule, respectively.124 The molecule will become excited and transition between electronic

states when an incident photon, or for that matter any electromagnetic radiation, is of

such energy that it matches the energy required for the state transition. Considering the

wavenumber of light, ν , is defined as the inverse of the wavelength, and is sometimes known as the frequency shift, this energy requirement can be written as:

c ∆E = hν = h = hcν (2.23) λ

This energy requirement, when compared to the electromagnetic spectrum, results

in the quantized observation regions depicted in Figure 2.24. In general, the energy terms

in Equation 2.22 are orders of magnitude apart; electronic energies lie in the UV-Vis

range of 104 – 106 cm-1, vibrational energies fall within the Raman/IR range of 102 – 104

cm-1, and rotational energies are within the microwave region of 100 – 102 cm-1, as

illustrated by Figure 2.25. Thus, Raman and infrared spectroscopies generally study the

effect of impinging light exciting a vibration within a molecule. Not all vibrations are

possible nor are all vibrations detectable by both techniques; there are rules based on the

symmetry of the molecule that determine whether the transition is possible, and then

whether it is Raman active, IR active, or both.124, 125

Figure 2.24. The electromagnetic spectrum by various units, and the analytical techniques operating in each region.124

Generally speaking, the difference in IR and Raman spectroscopies lies in what

effect the impinging light has upon the molecule’s dipole. In Raman, the molecule must

undergo a distortion of its electrons that momentarily alters the polarization of the

molecule for the vibration to be active. When the molecule relaxes, it scatters photons of a slightly different wavelength than the impinging light, defined by the wavenumber of the emitted light. In IR, vibrations are active when the dipole moment of the molecule is changed by the vibration, resulting in absorption of the impinging light. The difference in detection then lies in that IR spectra examine the reduced intensity at the initial wavelength of incident light, whereas Raman spectra examine the minute shifts in the wavelength of the incident light.124

As the vibrational spectroscopies are so sensitive to the energies required to induce dipole moment shifts and polarization in a molecule, the resultant spectra are extremely specific to exact bonds, including the elements being bonded. Thus, every phase of every material should have a relatively unique fingerprint by examining its vibrational spectra, and there have been attempts at determining the exact vibration of every peak within the spectrum of boron carbide, though some of these are somewhat disputed.126 Still, the differences between crystalline and amorphous phases are generally

readily observed by vibrational spectroscopies, and even the presence of dislocations and

twins, a relatively minor change in the bonding structure of a material, can be detected by

vibrational spectroscopies. The power of the vibrational spectroscopies and the

specificity inherent to them thus make them ideal for examining structural changes

induced by processing, and will therefore be one of the main techniques utilized in this

dissertation.

When applied to boron carbide, the vibrational spectroscopies are tools that have been relied upon for years for detailed bonding information about the material, yet neither

FTIR nor Raman microspectroscopy has been fully detailed as to the origin of each peak.

There is extensive literature on both techniques, often giving contradictory assignments to each characteristic peak in the spectra.10, 26, 36, 69, 86, 99-101, 104, 106, 107, 126-147 The expected spectra for FTIR and Raman microspectroscopy for boron carbide is shown in Figure

2.26.139

Figure 2.26. Expected Raman and IR spectra for boron carbide.139

The FTIR spectrum is of greater ambiguity between the two vibrational spectroscopies employed in this dissertation. The most common assignment of peaks to vibrations is proposed by Werheit et al. and Shirai et al.86, 139, 143, 146 These groups generally agree that the peak between 1560 – 1600 cm-1 is assigned to the

intericosahedral chain of boron carbide, typically for a C-B-C bonded chain. As boron substitutes for carbon on the chain, this peak should shift to a higher wavenumber. The intericosahedral chain has one other peak, at approximately 410 cm-1. However, this

peak is located at the low wavenumber detection limit for most IR spectrometers,

including the Mattson Galaxy 5000 used for this dissertation. As such, the peak at 410

cm-1 cannot be reliably used for analysis of boron carbide in this dissertation. All other

IR peaks are generally assigned to icosahedral vibrations and stretches without specific

assignment, with one notable peak located at approximately 1080 cm-1. The 1080 cm-1

peak is assigned to a general B-C bond and shifts to higher wavenumbers with increasing overall carbon content of the sample. This peak has been suggested to be influenced by vibrations of molecules attached to a host network; that is, it relates more to an impurity carbon source bonding to the boron carbide structure rather than purely to the stoichiometry of the boron carbide.141 Finally, a peak has been observed when examining

-1 amorphized boron carbide at approximately 800 cm . This peak is indicative of B12

icosahedra, and has been postulated to develop a low intensity, broad peak in response to

amorphization.106, 135 The three peaks at 800, 1080, and 1560 cm-1 are then the key

characteristic peaks that can help analyze boron carbide through FTIR.

Raman microspectroscopy as an analytical tool for boron carbide, although often

utilized in peer reviewed literature, has the same major drawback as FTIR in that many of

the characteristic peaks remain unassigned to specific vibrations. The spectrum possesses

many bands stretching from 270 cm-1 to 1080 cm-1. The two lowest wavenumber peaks,

at 270 cm-1 and 320 cm-1 are generally assigned to either rotations of the intericosahedral

chain86, 126, 129, 139, 140, 143, 144 or to a disorder-induced modified selection rule for Raman

activity and relate to acoustic phonons.137, 142 While the acoustic phonon assignment

appears to be inaccurate due to the presence of the peaks in anti-Stokes Raman spectra,147

indicating a true Raman vibration and not a disorder-induced acoustic phonon, the assignment of the peaks to the rotation of the intericosahedral chain also has not been confirmed. Most of the higher wavenumber peaks, between 600 cm-1 and 1200 cm-1,

have also not been definitively assigned a vibration, although most are believed to arise

from vibrations of the icosahedron itself. Generally, the peak at 1080 cm-1 is assigned to

the icosahedral breathing mode,140 but it is possible that several peaks contribute to this

very strong feature.137

There are two peaks in the Raman spectrum of boron carbide between the two

groupings discussed, located at 480 cm-1 and 535 cm-1. These peaks are typically

assigned to the intericosahedral chain,126, 129, 137, 143 although the 535 cm-1 peak has also

been attributed to the librational mode of the icosahedra through calculations.137, 148

Disparity in assignment arises from whether the vibrations responsible for these peaks are

both stretching modes of the chain,129 or whether the 480 cm-1 peak is possibly rotation of

the chain.137 Both assignments have experimental support, and thus either are possible.

Recent work by Domnich et al.133 has shown that while many researchers have noted the

dependence of the intensity of these peaks upon carbon content, the location of the peaks

also strongly correlates with stoichiometry. As shown in Figure 2.27, the peaks at 480

and 535 cm-1 tend to shift towards one another with decreasing stoichiometric carbon content at multiple laser wavelengths. This is significant as peak intensity and location are typically dependent upon laser wavelength, but for the peaks to shift towards each other across a multitude of wavelengths indicates a material phenomenon that can be used to identify stoichiometry of a boron carbide through Raman microspectroscopy.

Figure 2.27. Peak location of the boron carbide Raman active peaks at 480 and 535 cm-1 as a function of carbon content.133

Finally, boron carbide has a series of peaks that respond strongly to the presence

of amorphization: two standard peaks that are present even without amorphization above

1200 cm-1, located at 1330 cm-1 and 1580 cm-1, and a third peak which only appears with

the observance of amorphization at 1800 cm-1. It is tempting to assign the two lower

wavenumber peaks to free carbonaceous materials. It has been shown that two peaks in

that range to are characteristic of carbon, a “D” peak at 1370 cm-1 and a “G” peak at 1590

cm-1.149, 150 These peaks arise from the sp2 bonding of C=C bonds for the “G” band, and

breathing vibrations of aromatic rings within the graphite for the “D” band. The “D”

band is only active when there is structural disorder present in the graphite, allowing for

an electron-hole pair to form that releases a phonon upon reformation. However, these

peaks in boron carbide are not purely “D” and “G” peaks of carbonaceous material. Not

only are the peak locations shifted, indicating an altered bond relative to the typical “D”

and “G” bonds, but the relative intensities of the two peaks is impossible in carbonaceous

materials. The ratio of the “D” peak to the “G” peak can be used to determine crystallite

size of graphite149, 151, yet it never exceeds 2.5.152 In boron carbide spectra, the “D” over

“G” ratio is often between 4 and 5. It seems apparent that the bands located at 1330 cm-1

and 1580 cm-1 in boron carbide are not truly resultant from secondary carbon content

alone, and only the peak at 1800 cm-1 appears to be solely indicative of amorphization at

this time.

2.9.2 X-ray Diffraction

X-ray diffraction has been widely utilized since its first known observation in

1912, mainly in the identification of crystal structures of materials.153 The well-known fundamentals behind x-ray diffraction include application of Bragg’s Law during diffraction of an incident x-ray beam by a crystalline sample:

2d sinθ = nλ (2.24)

where d is the spacing between crystallographic planes in the material, θ is the angle of

diffraction, n is an integer, and λ is the wavelength of incident radiation. Diffraction

peaks only appear when constructive interference occurs at conditions that satisfy

Bragg’s Law, resulting in a finite number of diffraction peaks for every crystal. Thus,

analysis of the location and intensity of diffracted peaks can lead to precise identification

of the crystal structure of the sample, thus inherently lending itself towards the qualitative

identification of phases present and in turn phase transitions in general.125, 153-155

However, boron carbide, as mentioned in Section 2.3, has a very large solid

solubility range, and minor changes in the stoichiometry of the material may not be easily

detectable by typical qualitative analysis x-ray diffraction. For instance, it was suggested

that some stoichiometries are more susceptible to phase transitions than others,104 and the

nature of the Higgins furnace18 inherently lends itself to the possibility of incomplete

reactions, resulting in localized stoichiometric variations. To examine this element of the

nature of phase transitions, it is possible to use precision lattice parameter evaluation to

determine the lattice parameters of the crystal, which has been shown in boron carbide to

have a fairly direct correlation with stoichiometry.10, 36-41, 156 The exact relationship over

the full range of the solid solubility has been difficult to determine, leading to Conde et

al.’s10 assertion of the three-regime relationship depicted in Figure 2.28. Over the more

narrow range of 13 at% to 20 at% carbon, most literature and experimental work tends to

agree with the relationship shown in Figure 2.29, as proposed by Aselage et al.36 and endorsed by Mayo,38 which stresses the higher correlation of the a lattice parameter with

stoichiometry. For this dissertation, the following relationships, derived from work based

on Figure 2.29,38 will be utilized to correlate lattice parameters to stoichiometry:

a = 5.674343 − 0.003801[C] (2.25)

c = 12.357596 − 0.014427 [C] (2.26)

c = 2.177993 − 0.001108[C] (2.27) a

In the above equations, a and c are the respective lattice parameters in Angstroms, and

[C] is the atomic percent carbon in the boron carbide sample.

Figure 2.28. Three-regime correlation between stoichiometry and lattice parameters as proposed by Conde et al.10

Figure 2.29. Correlation between lattice parameters and stoichiometry as observed by Aselage et al., depicting a linear relationship between 13 at% and 20 at% carbon.36

Typical x-ray diffraction scans can be used to obtain rough lattice parameter data, but this technique is subject to extreme variations in calculated values of lattice spacing, d, due to a non-linear dependence of sinθ upon θ, as shown in Figure 2.30. Thus, care must be taken to minimize the effect of this non-linear dependence.

Δsinθ

Δθ

sinθ

Figure 2.30. Generalized dependency of sinθ upon θ, leading to inaccurate calculations of d-spacings in lattice parameter calculations.154

The first step is to use many high-angle peaks, not normally examined under a typical qualitative scan, up to 120° 2θ or more, as shown in Figure 2.31. Second, an extrapolation function can be used to calculate more accurate lattice parameter data.

Cohen’s Method is widely used as it does not assume anything about systematic errors dominating the dependence, and treats every possible cause of error in the scan and calculations.155 Through a series of least squares refinements, a fit is made to the location of the peaks in the diffraction scan, and the lattice parameters are calculated. A least

squares refined fit and lattice parameter calculation using Cohen’s Method can calculate

precision lattice parameters to within 30 ppm on a typical diffractometer. By

comparison, TEM lattice imaging is typically only able to calculate lattice parameters to

within 10,000 ppm of the true value.156 Thus, the highly accurate precision lattice

parameter calculations by x-ray diffraction can be used to obtain stoichiometric

information about the boron carbides examined in this dissertation and correlate the stoichiometry to the phase transitions observed.

Figure 2.31. Typical XRD pattern for boron carbide extended to 120 ° 2θ, with peaks for boron carbide and graphite delineated.133

3. Method of Attack

Boron carbide powder synthesis and processing is complex. This dissertation will focus on one specific area, namely, the effect of contact pressure during comminution from large fused aggregates to fine particle sized powders. To determine the extent of variations and transformations evident in processing of boron carbide, this dissertation will first simulate the powder comminution stages of powder processing on raw, unprocessed boron carbide. Then, sequentially finer grained commercially available powders will also be comminuted to determine whether commercial powders are subject to any variations or transformations in structure induced via processing. These powders will be analyzed to determine first whether any preexisting bulk property-threatening impurities, such as secondary carbon, have been introduced during processing, and second whether the amorphization event so evident in high pressure situations has been induced via processing. Both the evolution of material specific impurities and the rise of amorphization, if either or both are present due to processing, will be attempted to be eliminated or accommodated by further processing.

3.1 Objective 1: Comminution of Boron Carbide

Accomplishing this objective will require first obtaining raw boron carbide from a powder producer. In the Americas, there is only one producer of boron carbide powder-

Washington Mills, Inc. Washington Mills is a powder producer specializing in high temperature and high hardness ceramic powders, ranging from oxides like aluminum oxide and zirconium oxide to carbides such as silicon carbide and boron carbide. With the help of Washington Mills, raw, unprocessed boron carbide will be obtained from

across the ingot product of the carbothermal reduction furnace. Because of lot size for

their materials, Washington Mills utilizes a Higgins Furnace18 for their carbothermic reactions, which is essentially a thin metal pot with constant water cooling that is lined with loose coke insulation and heated by three phase electric arc melting using three graphite electrodes close to the top surface of the melt. Boron carbide materials will be gleaned from locations near the unreacted edges of the furnace; partially-reacted and well-reacted melt areas; and locations of possible over-reaction and graphitization, as close to the graphite electrodes as possible. These materials will represent a fairly wide range of product, inasmuch as some will be material that is preferentially screened the remainder considered scrap. Still, due to the optical nature of the classification process in typical powder factories, this scrap material may be mixed in with sold product to some degree and will be examined also.

Prior to comminution and at each stage of comminution, each material will be analyzed as discussed below. Then large, inch-sized aggregates from the raw ingot, broken off by hand and hammer, will be jaw crushed to obtain an appropriate feed size for the shaker mill. The shaker mill will reduce particles from roughly 5 mm and less to

200 μm and finer. These powders will then be jet milled, first to obtain a uniform fine particle size distribution on the order of a single micron. Then, iterative jet milling will be conducted to examine the capabilities of the jet mill to reduce particle size and induce phase transformations in the powders. Select powders that exhibit comminution induced transformations or variations will be subjected to further processing in an attempt to remove introduced impurities and reverse any induced transformations.

3.2 Objective 2: Characterization of Comminuted Powders

The boron carbide powders will be analyzed for the presence of phase

transformations through the use of vibrational spectroscopic and x-ray diffraction

techniques. The vibrational spectroscopies of FTIR and Raman microspectroscopy, for

their ease of use, quick data collection time, and the capability of Raman

microspectroscopy for examining any size scale product, make them invaluable

characterization tools. The nature of Raman as a surface analytical technique will be

beneficial as the most probable locations for phase transformation will be near the surface

of the particles, near locations of fracture during comminution. For this dissertation, a helium-neon laser of wavelength 633 nm will be utilized. As shown in Figure 3.1, different wavelength lasers will excite varying portions of the boron carbide spectrum preferentially. The 633 nm laser is, of the lasers available at Rutgers, more responsive to the appearance of high wavenumber peaks indicative of amorphization with less lost signal due to over-excitation of low wavenumber peaks. IR spectroscopy will be used to compare and confirm the results given by the Raman spectra, although only on fine powders of less than 100 μm particle size. X-ray diffraction for phase identification and precision lattice parameters will be conducted on any powder of similar particle size

requirements to the IR spectroscopy. In completion of this part of the objective, it will be

necessary to comment upon the current state of information within the international

Powder Diffraction File used for phase analysis, as recent findings have shown a lack of

accuracy for a multitude of boron-containing hard materials. Field-emission scanning

electron microscopy will be utilized to note morphological changes and surface flaws in

the powders.

Figure 3.1. Raman spectra of boron carbide before (a) and after (b) nanoindentation, displaying the propensity of particular wavelength lasers to excite various vibrations.132

3.3 Objective 3: Elimination of Impurities and Reversibility of Transformations

Of profound impact upon the future of boron carbide as a multi-application

structural ceramic will be the ability to effectively eliminate any undesirable variations

and transformations induced by the powder processing. This will be attempted first

through chemical and thermal attack of any noticeable impurities, specifically carbon,

that evolve during the comminution process. Several methods have been discussed for

chemical and thermal attack of carbon.157-162 Out of these methods, this dissertation will

conduct thermal attack at low temperature and high temperature in air and in inert

atmosphere. For chemical attack, the chemical solution suggested by Nazarchuk161, 162

seems to hold the most promise. The success of each method will be gauged by the

reduction and/or the elimination of the carbon presence in the sample without inducing

further transformations. Secondly, any amorphization induced during the extent of the

comminution process will be heat treated in an attempt to evaluate the reversibility of the

transformation. Low temperature heat treatments were shown to have success in

reducing the amorphization signal in Raman spectra previously.107 These low

temperature treatments will be duplicated for the current study on comminuted powders.

If unsuccessful, higher temperature treatments and treatments in differing atmospheres

such as argon, nitrogen, 15% hydrogen, and high vaccum will be attempted. Again, the success of the heat treatments in reversing the amorphization event will be judged by the reduction and elimination of amorphization peaks in the vibrational spectroscopies.

4 Experimental Procedures

4.1 Comminution

Commercially available boron carbide powders were obtained from two

producers, ESK and UK Abrasives, in a range of sizes, from several millimeters to less

than 100 μm. Those powders that were finer than approximately 75 μm and could pass through a 200 mesh sieve were held for direct jet milling. Powders larger than 75 μm were milled in a multi-axis vibratory mill for increments of ten minutes until they passed through the 200 mesh sieve and were then held for jet milling. In contrast, raw unprocessed boron carbide samples were obtained from Washington Mills and were in the form of large quarter meter segments directly taken from the ingot, shown in Figure

4.1a prior to initial breakage and 4.1b as the quarter meter segment. No processing was performed on these samples following removal from the carbothermic melt. The samples were crushed by hand and then placed into a benchtop jaw crusher, shown in Figure 4.2, which reduced the particle size to approximately 200 μm. The jaw crushed powders were then milled in a multi-axis vibratory mill, the Spex SamplePrep 8000m Mixer/Mill shown in Figure 4.3, with aluminum oxide milling media used to assist comminution as it is relatively easy to distinguish from boron carbide through the analytical techniques utilized. This milling served to reduce particle size to less than 75 μm, at which point they passed the 200 mesh sieve. This was considered an optimal starting particle size for the jet mill.

Figure 4.1. a) Raw ingot of boron carbide prior to separation and classification. b) Quarter-meter long segment that has been separated by hand.

Figure 4.2. Benchtop Jaw Crusher.

Figure 4.3. Spex SamplePrep 8000m Mixer/Mill (image courtesy Fisher Scientific).

Jet milling was conducted on a Sturtevant 2 in. Micronizer model laboratory scale table-top jet mill with tungsten carbide lining, shown in Figure 4.4. Initial comminution runs were done at the highest pressures possible, with a feed backing air pressure of 100 psi and a grinding air pressure of 100 psi. Feed rates of the stock boron carbide into the jet mill were 5 grams per minute via vibrating hopper. Resultant average particle sizes were on the scale of 900 nm. By lowering the feed rate and air pressure of both the feed

79 and grind lines, it was possible to reduce the overall energy of impact and comminution inside the mill. Select samples were jet milled iteratively, so as to observe the effect of repeated high-energy impact upon the microstructure of the powders. Further samples were milled at air pressure of 80 psi and 60 psi for both feed air and grind air in an attempt to determine a lower limit for observed microstructural changes.

Figure 4.4 Sturtevant 2 in. Micronizer Laboratory scale jet mill.

4.2 Particle Analysis of Powders

4.2.1 Light Scattering Particle Size Analysis

Particle size analysis was performed using a Malvern Mastersizer 2000 light scattering particle sizer, shown in Figure 4.5, with a Hydro 2000S wet sample analysis accessory. The Mastersizer was able to generate particle size distribution data in the size range of 200 nm through 2000 μm, although light scattering particle size reliability is limited below 500 nm due to the wavelength of light. Powders were dispersed in

80 deionized water at a low 1% solids loading of 0.5 grams of powder per 45.5 mL of water.

The Mastersizer was operated under conditions set to de-air the water, so as to not encounter air bubbles that would skew the particle size distribution. These conditions included ultrasonication of the sample in the system and use of a rotating mixer blade.

The sample chamber had to be examined after each run to determine if any sample material had coated the window of the sample chamber, and if so, taken apart and cleaned.

Figure 4.5. Malvern Mastersizer 2000 Light Scattering Particle Sizer with Scirocco Dry Powder Feeder (front left) and Hydro S Small Volume Wet Dispersion (front right) accessories.

4.2.2 Scanning Electron Microscopy

Particles were also examined on the Field Emission Scanning Electron

Microscope (FESEM) for general morphology and evidence of fracture. Samples were produced by loading a small amount of powder onto an aluminum stud covered by double-sided carbon tape. Canned dry air was then used to blow excess powders off the stud, leaving only small unagglomerated quantities of powder remaining on the stud.

This methodology was determined to be optimal for sample preparation of these powders due to the ability to view small numbers of particles at a time without having to compensate for any agglomeration or drying effects due to sample preparation via liquid mediums. The samples were not coated due to the desire to examine the surfaces in detail. Samples were then examined using a Zeiss Σigma (sigma) FESEM, depicted in

Figure 4.6, with secondary electron imaging. Initial operating conditions were a voltage of 2.5 kV and a working distance of 8.0 mm, adjusted as necessary per individual sample’s charging characteristics. Images were taken via line integration to obtain crisp, detailed images.

Figure 4.6. Zeiss Σigma field emission scanning electron microscope (image courtesy Carl Zeiss NTS, LLC).

4.3 Vibrational Spectroscopy

4.3.1 Raman Microspectroscopy

Raman microspectroscopy was of critical importance in this dissertation due to the aforementioned strong response to amorphization in boron carbide. The system utilized was a Renishaw inVia Confocal microRaman Spectrometer, shown in Figure 4.7,

with a laser of wavelength 633 nm and laser specific penetration depth into boron carbide

of less than 1 μm. This wavelength was chosen for its low influence on the low

wavenumber “disorder” peaks in the spectrum of boron carbide compared to the 785 nm

wavelength laser, as shown in Figure 3.1 and mentioned in Section 3.2. The

microspectrometer was used with a 20x objective lens, for an effective spot size of

approximately 5 μm diameter, to counteract localized inhomogeneities vastly influencing

the acquired spectra. With an exposed area of close to 20 μm2, and average particle size

of the comminuted powders being approximately 900 nm, scans typically examined tens

of particles at a time. This higher number of particles exposed to the beam at one time

led to more statistically significant acquisitions from each point. Twenty acquisitions of

data were accumulated for each spectrum, leading to spectral scan times of half an hour,

and spectra were taken from at least twenty locations on each sample.

Figure 4.7. Renishaw InVia Confocal micro-Raman Spectrometer (image courtesy Renishaw plc).

Powders were loaded onto glass slides for initial examination, with no attempt made to tightly pack the powders. However, the nature of the flat surface of the glass slide may have induced some level of texturing in the thin layer of powders placed onto the slide. To counteract this texturing, a select number of samples were extruded to induce texture and then examined from multiple angles to determine any spectra dependence upon orientation of the particles.

Acquired spectra were analyzed using Renishaw’s incorporated Wire version 2.9 software. Background removal was performed using a cubic spline fit to attain the best fit possible. Following background removal, smoothing at 9 reference points, and location of peak centroids, spectra were examined for the presence of amorphization peaks and the intensity of such, if present. Spectra were then imported into OriginLab’s

OriginPro 8 for plotting and minor mathematical operations as desired, such as average spectra for an entire sample.

4.3.2 Fourier Transform InfraRed Spectroscopy

Fourier Transform InfraRed Spectroscopy (FTIR) was performed on comminuted powders using a Mattson Galaxy 5000 spectrometer, shown in Figure 4.8. Prior to comminution, FTIR could not be performed on the samples due to the large size of the specimens that would not fit into the FTIR sample holder. Spectra were acquired from

400 to 4000 cm-1 so as to be able to determine the validity of each spectra and the amount of distortion due to atmospheric humidity. Initial spectra were acquired in transmission mode, on pelletized samples compacted from 100 mg of potassium bromide (KBr) containing 0.3 mg of comminuted boron carbide powder. These transmission samples contained so little boron carbide due to the extreme scattering power of the material.

More than 0.3 mg of boron carbide in each pellet would result in almost perfectly non- transmitting samples.

Figure 4.8. Mattson Galaxy Series 5000 FTIR Spectrometer.

To counteract this issue, a diamond Attenuated Total Reflection (ATR) cell was

obtained. The ATR cell allowed for powder samples to be examined without sample size

limitations due to lack of transmission of the infrared source. Approximately 10 mg of

powder were loaded onto the diamond surface of the ATR for each spectrum. An anvil was used to apply minor pressure to the powders purely to hold them in place on the diamond surface. Initial FTIR spectra taken using the ATR were unidentifiable, as shown in Figure 4.9, with no peaks at all evident. Following repeated attempts to improve the spectra from the ATR, it was determined that boron carbide could not be examined using the available ATR system due to its extreme scattering nature. As a consequence, no samples with particle size larger than a few hundred micrometers could be examined by FTIR during this dissertation.

Figure 4.9. Sample FTIR spectra taken with ATR setup. No pattern is evident.

Spectra were taken from at least ten sub-samples from each comminuted powder.

At least 30 acquisitions were taken for each initial spectrum to check successful KBr

pellet formation. Full analysis was conducted on spectra composed of 400 acquisitions,

with background spectra acquired prior to each individual sample. Background removal

was performed automatically by the incorporated software Winone. Spectra were

examined in the range 500 – 2000 cm-1 for influence of amorphization on the spectrum boron carbide, specifically interested in the “amorphous boron” peak at 800 cm-1, as

previously shown in Figure 2.12.106 Full mathematical spectral averaging and plotting

were performed using OriginLab’s OriginPro 8.

4.4 X-ray Diffraction

X-ray diffraction (XRD) experiments were conducted with a Panalytical X’Pert

MPD model powder diffractometer, shown in Figure 4.10, with a copper radiation source

operating at 35 kV and 30 mA. As the high angle peaks were of particular interest in

determining lattice constants, diffraction patterns were taken in a continuous scan from

10 – 140° 2θ at a virtual step size of 0.02° and dwell of 15 seconds. Powder samples

were hand mixed with pure silicon powder for use as an internal standard. In cases

where, due to the extreme x-ray transparency of boron carbide, a packed powder bed was

simply not desirable, a Panalytical X’Pert MPD Pro at 45 kV and 40 mA was utilized

with a rotating stage and single crystal silicon zero-background holder. This enabled diffraction patterns to be obtained from a monolayer of comminuted powder dispersed in

methanol deposited on the holder.

Figure 4.10. Panalytical X’Pert MPD X-Ray Diffractometer.

Lattice parameters were determined within MDI Corporation’s Jade version 9.0 analytical software. A hybrid least squares fit/whole pattern refinement approach was utilized to determine lattice parameters, based on the whole diffraction pattern, to within

+/- 0.0003 Angstrom precision. These values for the lattice parameters were used to determine the local stoichiometry of comminuted samples. The diffraction patterns were also examined for the presence of amorphous humps indicating an influence of the induced amorphization on the XRD patterns.

4.5 Excess Carbon Removal

During the course of the dissertation, it became evident that excess carbon would be encountered while attempting to characterize boron carbide powders, much to the detriment of many of the characterization techniques. Attempts were made to determine the optimal method for carbon removal based on available literature and industry

“general knowledge.” The most frequent suggestions were to conduct excess carbon removal via simple froth flotation, wherein a liquid with density between that of the two materials in question, boron carbide and free carbon, is used to disperse the powders.

The liquid is then frothed to promote separation of the two materials, with the denser material sinking and the less dense material floating atop the frothing liquid. However, this was deemed an inappropriate method due to it being a purely mechanical separation technique. Some of the excess carbon observed in the boron carbide samples was chemically bonded to the boron carbide, and thus would require a more direct method of removal.

For this reason, attack of the carbon through heat and chemical oxidation was attempted. For thermal removal of excess carbon, samples were fired to 900°C in air and three flowing atmospheres of argon, nitrogen, and 15% hydrogen in argon.

Thermogravimetric analysis (TGA) was also conducted on samples utilizing a Perkin

Elmer TGA 7, depicted in Figure 4.11, up to 850°C in air and nitrogen atmospheres.

Further high temperature firings were conducted under high vacuum.

Chemical oxidation routes of carbon and boron carbide were detailed in depth by the work of Nazarchuk et al. 161, 162 The concern with chemical oxidation was whether or not routes that were significantly strong enough to oxidize carbon would also oxidize the

boron carbide, destroying the microstructure of interest to this dissertation. The

formulation recommended by Nazarchuk for successful oxidation of carbon with minimal

impact on the boron carbide is:

3 3 3 10 cm HNO3 + 10 cm H2SO4 + 5 cm HClO4 + 20 mg K2Cr2O7 (4.1)

However, as K2Cr2O7 is a known carcinogenic, first attempts were made with mixtures of

nitric acid and peroxide as well as mixtures of potassium manganate and sulfuric acid.

Samples were boiled in the mixtures for 30 minutes at ten minute intervals then analyzed

for weight loss and oxidation of both carbon and the boron carbide.

Figure 4.11. Perkin Elmer TGA 7 apparatus (image courtesy Fleming Polymer Testing & Consultancy).

4.6 Extrusion of Boron Carbide Rods

To investigate the possibility of orientation dependence of induced microstructural variations and transformations, samples were fabricated with induced texture via capillary extrusion. Batches were composed of 77 wt% boron carbide powder, 3 wt% methylcellulose binder (Hercules Culminal Nonionic Cellulose Ether

Type: MC-3000P, Hercules Tianpu Chemical Company Ltd), and 20 wt% DI water.

Batches were plasticized using a Haake Rheocord 9000 at 50 rpm for 30 minutes, to ensure activation of the methylcellulose binder and retention of shape following extrusion. Extrusion was then performed using a Malvern Instruments Rosand RH2000

Capillary Rheometer, as seen in Figure 4.12, fed through an orifice of 1 mm with a die length of 16 mm at a piston speed of 2 mm per minute, with no applied heating.

Figure 4.12. Malvern Rosand RH2000 Capillary Rheometer.

Induced texture was determined via x-ray diffraction. The extruded rods of boron carbide required the use of the Panalytical X’Pert MPD Pro system for its linear detector.

Without the specialized detector, the small sample size of the rod, 1 mm in diameter on the plane of incidence, would have severely limited the intensity of the spectra. Whereas quality data acquisition would have required days per sample spectra on the MPD system, the MPD Pro enabled quality data acquisition in 2-3 hours per sample. The extrudates were mounted in a small amount of backing putty, oriented so the cross section of the extrudate lay on the plane of incidence for the x-ray beam. A glass slide was used to push the extrudates into the putty such that the flat face of the extrudate was level with the sample holder. Diffraction data were collected from 15° to 60° 2θ to enable examination of the intense low index peaks. Texture was then determined from the intensities of the peaks by following the equation for the Texture Coefficient TC given by

Barrett and Massalski155:

I (hkl) I TC = 0(hkl) (4.2) I 1 ∑ (hkl) n I 0(hkl)

Here, n is the number of peaks being examined, I(hkl) is the observed intensity for a given index peak, and I0(hkl) is the reference intensity for that given index peak from the PDF card file. A TC of 1 is considered perfect random orientation, with a value higher than 1 indicating preferred orientation along that plane. To double check the accuracy of the TC

92 calculations, the sum of all the TC values for all peaks examined should be equal to the number of peaks examined, n.

5. Results and Discussion

5.1 Comminution

Within this section, the low energy and high energy comminution of several

boron carbide materials will be discussed. For the purposes of the rest of the discussion,

a naming convention must be used. Powders referred to as “commercial” powders are boron carbides of various average particle size that can be purchased as-is from commercial boron carbide suppliers. “Unprocessed” powders refer to boron carbides that were obtained from Washington Mills directly from the ingot following carbothermic reduction, without any post-melt processing or refinement.

5.1.1 Low Energy Comminution: Vibratory Milling

Large segments of raw boron carbide and coarse commercial boron carbide were separated by hand and jaw crushed to approximately 1 mm in particle size. Initial low energy comminution was conducted on coarse commercial powders from two industrial providers of boron carbide, ESK and UK Abrasives. During vibratory milling for 10, 20, and 30 minutes, no significant reductions in particle size have been observed through optical measurements for ESK powders. For UK Abrasives powders, 10 minutes of vibratory milling led to a size reduction of average particle size being roughly 0.5 mm, with the size of the powders after 20 minutes of milling further reduced to the range of 5-

200 µm, and yet more size reduction to a range of 2-100 µm after 30 minutes of vibratory milling. Both commercial series, when vibratory milled for 60 minutes, had a reduction of average particle size to the range of 1 μm. The resultant powder shards from 30 minutes of vibratory milling, shown in Figure 5.1, display a great deal of debris, as well

as two distinct surface “types”: flat and highly optically reflective, or dark and non-

reflective of visible light. From optical observations, milling led to more fractured

surfaces with higher amount of dark and less reflective areas.

The raw, unprocessed boron carbide was obtained from Washington Mills. The

0.25 m segment portrayed in Figure 5.2a is an example of material that was entirely classified as “good” boron carbide that would then be further processed and sold as technical grade boron carbide. The segment was sorted into the four sections portrayed in Figure 5.2b-e based on visual appearance- a relatively solid and well reacted section,

“S;” a section of large, 1 cm or more in diameter, porosity, “P;” a section that appeared to resemble a more carbonaceous material, “C;” and a center band that resembled the S

95 section, “CB.” Each section was subjected to 30 minutes of vibratory milling and screened with a 325 mesh sieve. The initial SEM micrographs after comminution, shown in Figure 5.3, present the powders as having an average particle size of approximately 10

μm, with apparent coatings of fine, nano-scaled cluster-like particles. When examined under Energy Dispersive Spectroscopy (EDS) in the FESEM, the clusters appeared to be comprised mostly of carbon. All four sectors possess large, chip-like particles.

However, the “C” and “S” sectors appear to possess a large relative number of aggregates instead of single grain particles.

CB C P

(a)

(b) (c)

(d) (e)

(a) (b)

5.1.2 High Energy Comminution: Jet Milling

Following appropriate vibratory milling and screening, powders were jet milled in the Sturtevant 2 in. Micronizer at 100 psi grinding pressure at a feed rate of 5 grams per minute. As mentioned in Chapter 2.8, contact pressures for the boron carbide powders could vary between 1 – 800 GPa depending upon effective geometry of impact. This range of contact pressures is mostly far above the pressure for inducement of amorphization, determined to be between 15 – 40 GPa.99, 107 Initial tests on grinding power of the jet mill were performed with a commercially available 100 μm particle size

98 boron carbide powder. The resulting particle size distributions, shown in Figure 5.4, clearly show the efficacy of the jet mill in comminuting even hard ceramics. One round of jet milling resulted in the initial 100 μm average particle size powder being reduced to an average particle size of less than 1 μm.

Figure 5.4. Particle size distributions of test boron carbide powders before (top) and after (bottom) jet milling at 100 psi and 5 grams per minute feed rate.

For the unprocessed boron carbide, more of the visually classified “good” unprocessed segments were obtained from Washington Mills and subjected to jet milling, as the bulk of the C, CB, P, and S powders from vibratory milling had been consumed by various analytical testing methods. These segments were subjected to the same vibratory milling and screening conditions as the previous powders. The resultant -45 μm powders were then jet milled at 100 psi feed air pressure and grinding air pressure. Interestingly, a large amount of powder was lost to the fines filter bag on the Micronizer. Inspection of the difference in particle size between the fines filter bag and the product collection bin,

shown in Figure 5.5, found the powders in each chamber to be roughly of the same

particle size, with a D90 less than 900 nm. However, the light obscuration factors of

samples from either chamber were drastically different, implying a fundamental

difference in the elemental nature of the powders, and suggesting that the less dense

carbonaceous materials have been somewhat classified in-line and separated from the

denser boron carbides.

Following initial jet milling, samples were taken from the powders for x-ray

diffraction, FTIR, and Raman microspectroscopy analysis. The jet milled powders were

then processed through the jet mill again a second and third time, in an attempt to

increase the observed frequency of any induced transformations. Due to the nature of the

small particle size of the powders, during the second and third rounds of jet milling

100

nearly all the powder was classified into the fines filter bag. As such, it became

necessary to collect samples from the fines filter bag to analyze, as no sample set could

be gleaned from the product collection bin. This also required the second and third

rounds of jet milling to be conducted on powder combined from that collected from the

fines filter bag and the product collection bin, rather than merely on the powder in the

product collection bin. SEM micrographs of the powders after 1, 2, and 3 rounds of jet

milling are shown in Figure 5.6a-c. The powders did not exhibit drastic particle size

reduction following iterative jet milling. However, the powders did appear to become

more and more concentrated with small clusters of nanoscale aggregates, with the larger

~ 900 nm particles nearly obscured by the aggregates after two and three rounds of jet milling. In order to assess the continued feasibility of inducing transformations in boron carbide powders, commercial 10 μm particle size powders were subjected to one round of jet milling and analyzed, shown in Figure 5.6d. Average particle size of the commercial

powders reduced to 1.6 μm following jet milling, although there are obvious large 10 μm

particles remaining.

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(a) (b)

(c) (d) Figure 5.6. SEM micrographs of unprocessed boron carbide following (a) one, (b) two, and (c) three rounds of jet milling, and (d) commercial boron carbide after one round of jet milling. All powders display marked presence of nanoscale material coating the larger micrometer scale particles.

5.2 Analysis

5.2.1 X-ray Diffraction

X-ray Diffraction (XRD) patterns were taken for all unprocessed boron carbide

powders following comminution. Due to the rough surface of the large unprocessed

segments, it was impossible to obtain a flat surface to satisfy Bragg conditions for

diffraction prior to comminution. Also, x-ray diffraction of boron carbide is intrinsically

difficult due to the x-ray transparency of boron carbide, and several issues with the

structure database, which will all be discussed in the Appendix.

102

Typical x-ray diffraction patterns are shown in Figures 5.7-10 for the vibratory milled boron carbide powders. Table 5.1 lists the four vibratory milled samples, C, CB,

P, and S, with the weight percent graphite in each sample, lattice parameters, and approximate carbon content in the boron carbide by lattice parameters, as discussed in

Chapter 2.9. Interestingly, the weight percent graphite in the four samples is drastically different. Although the C sample appeared visually to be the most resembling an unreacted carbonaceous precursor material, it had by far the lowest weight percent graphite of all the samples, at 10 wt% graphite. Sample S, which visually appeared to be a well reacted and desirable boron carbide, had 65 wt% graphite. The weight percent graphite gradually increases along the length of the segment from the C sample to the S sample, although sample P was variable between multiple diffraction scans, with the average being 20 wt% graphite. As the S sample was located on the edge of the segment

that would have been in contact with the graphite electrode in the arc reactor, it is likely

that some graphite from the electrode may have broken off and been included in this

section. Also, as the area closest to the electrode, sample S most likely was the most well

heated section of the segment. If high heat was maintained for long enough, the boron

carbide could have begun to leach carbon from the electrode as boron diffused into the

electrode and the melt achieved a higher resultant concentration of graphite. The

relatively low graphite weight percent in the C sample may be explained as the precursor

material might not have been graphite but rather another form of carbon, and will be

discussed further in the vibrational spectroscopy results.

103

Table 5.1. Vibratory milled boron carbides and select data obtained from analysis of the x-ray diffraction patterns. Wt % %C Average Sample Graphite a (Å) c (Å) c/a (c/a) %C (a) %C C 10 5.6045 12.0855 2.1564 19.6 20.7 19.7 CB 15 5.6054 12.0884 2.1566 19.5 20.4 19.5 P 20 5.6098 12.1027 2.1574 18.7 19.1 18.5 S 65 5.5995 12.0691 2.1554 20.6 22.1 20.9

Figure 5.7. Typical XRD pattern for vibratory milled sample “C.”

104

Figure 5.8. Typical XRD pattern for vibratory milled sample “CB.”

105

106

Figure 5.10. Typical XRD pattern of vibratory milled sample “S.”

The impressive difference in weight percent graphite in each sample was not the only intriguing result gleaned from XRD analysis. Lattice parameter refinement from the diffraction patterns enabled calculation of stoichiometry of each sample. The entire segment at large was considered “good” boron carbide and processed to meet commercial specifications, and as such it would have been assumed to have a certain level of consistency in stoichiometry across the length. However, the lattice parameter data clearly shows a trend of decreasing stoichiometric carbon content from the far edge of the segment, the C section, through towards the near edge. The S section, with its probable presence of overreacted material from being too near the electrode and excess graphite

107

from the electrode itself, has a much higher stoichiometric carbon content than the other

sections. Although the calculated carbon content is higher than the stoichiometry range

of boron carbide when examining the a lattice parameter, the trend is still valid. This

trend implies that stoichiometry of boron carbide will vary along the length of an ingot, with maxima at the nearest and farthest edges of the ingot in relation to the electrode.

Table 5.2 lists the lattice parameters and stoichiometric carbon content in the

iteratively jet milled boron carbides. As the multiple milled powders were of the same

origin as the single milled powder, it was expected that the stoichiometries of these

samples would be constant throughout the experiments. Indeed, the calculated

stoichiometric carbon contents are within experimental deviation of one another, with the

Milled 3x sample having slightly higher calculated carbon content. This higher carbon

content is likely experimental error arising from the presence of secondary phases within

the diffraction pattern, as it is extremely unlikely stoichiometric carbon would increase

with comminution. It can be concluded that comminution will not induce a wide-spread stoichiometry change, whether through evolution of material out of the boron carbide or through deformation mechanisms such as slip planes and dislocations inducing a fundamental structural change.

108

Table 5.2. Lattice parameters and resultant stoichiometric carbon content calculated from x-ray diffraction patterns for iteratively jet milled boron carbide. %C %C Average Sample a (Å) c (Å) c/a (c/a) (a) %C Mill 1x 5.6025 12.0745 2.1552 20.7 21.3 20.5 Mill 2x 5.6009 12.0721 2.1554 20.5 21.7 20.7 Mill 3x 5.6009 12.0671 2.1545 21.4 21.7 21.1

It was hoped that XRD patterns of these samples, shown in Figures 5.11-13,

would be useful in determining overall amorphous content of the jet milled powders, if any amorphization did indeed occur during jet milling. However, analysis of the diffraction patterns was hindered by the presence of broad peaks indicative of nanoscale tungsten carbide. This material lines the Micronizer, while the SamplePrep 8000m vibratory mill is made of aluminum oxide and was used with aluminum oxide media.

Despite the nature of the jet mill preventing impact of feed material with the walls and

liner of the mill, it is apparent that a small degree of wearing of the liner of the jet mill

did indeed occur. The worn liner material is nanoscale, and thus nearly impossible to

segregate from the product. Unfortunately, the presence of this nanoscale material makes

calculation of the amorphous content of the powders from the diffraction patterns

unreliable. A full pattern fitting refinement would analyze the shapes of the peaks and

background of the diffraction pattern for consistent peak broadening and low, nearly

undetectable amorphous humps in the background of the pattern. The nanoscale

powders, while having clearly identifiable peaks, interfere with the amorphous humps

and thus render calculation of amorphous content impossible.

109

Figure 5.11. Typical XRD pattern of boron carbide after 1 round of jet milling.

110

Figure 5.12. Typical XRD pattern of boron carbide after 2 rounds of jet milling.

111

Figure 5.13. Typical XRD pattern of boron carbide after 3 rounds of jet milling.

5.2.2 Fourier Transform InfraRed Spectroscopy

5.2.2.1 FTIR for Carbon Content and Location

The FTIR data of vibratory milled unprocessed boron carbide, shown in Figure

5.14, is useful for helping to determine the nature of the carbonaceous variations first illuminated by x-ray diffraction. To reiterate the discussion in Chapter 2.9, the spectra depict the presence of carbon in boron carbide through the intensity and shift of absorption bands at 1080 cm-1 and 1570 cm-1.137, 139, 142, 143 The 1080 cm-1 band is

indicative of the B-C bond, and therefore enables relative measuring of overall carbon

content. A shift to higher wavenumbers and increasing relative intensity compared to the

112

rest of the spectrum implies higher overall carbon content in the sample. The 1560 cm-1

band, however, is characteristic of the intericosahedral chain, and will shift to higher

frequencies with decreasing stoichiometric carbon content as boron atoms substitute onto

the intericosahedral chain. The overall intensity of FTIR spectra will also decrease with

increasing free carbon content due to the high scattering and absorption of carbon.

Intensity will also decrease with increasing weight percent boron carbide in the sample

pellet, requiring careful sample formation. From the spectra, it can be observed that sample C is composed of a high number of B-C bonds through the intensity of the 1080 cm-1 band and high stoichiometric carbon content through the lack of shift of the 1560

cm-1 band. Sample CB is possessed of a moderate amount of B-C bonds yet lower

stoichiometric carbon than sample C. The most intriguing spectrum is for sample P,

which indicates relatively low stoichiometric carbon through the large shift of the 1560

cm-1 band to 1570 cm-1, yet has a high degree of B-C bonds indicated by the strength of

the 1080 cm-1 band. These factors would imply that sample P has low stoichiometry, yet

a relatively large amount of impurity carbon content, most likely either as an isolated

substitution or thin film. Through the low intensity of the S sample spectra, it can be

inferred that S has a very high presence of free carbonaceous material, with the lack of

shift in the 1560 cm-1 peak and low intensity of the 1080 cm-1 band indicating the

stoichiometry of sample S is carbon rich. Similar to the x-ray diffraction data, the FTIR

spectra of vibratory milled boron carbide then indicate a range of stoichiometries across

the length of the unprocessed boron carbide segment, with the most carbon rich

stoichiometries on either end of the segment. There also appears to be a great deal of

113 carbon present in the powders, either as a large free carbon, as seen in sample S, or as possibly a thin film, as seen in sample P.

5.2.2.2 FTIR Analysis for Presence of Amorphization

The FTIR spectra of jet milled powders looking specifically for amorphization was more productive than the lack of peer reviewed literature on the subject would lead one to believe. Spectra of several unprocessed boron carbides following jet milling are shown in Figure 5.15a with an expanded view of the low wavenumber region in Figure

5.15b. Of the jet milled powders, approximately 1 in 2 FTIR samples possessed a spectrum similar to those in Figure 5.15. These spectra are distinguished by the evidence of a low broad band at or just below 800 cm-1. The peak at 800 cm-1 has already been

114

discussed as evidence of amorphized B12 icosahedra. However, the low amorphous

bands in many samples were shifted below 800 cm-1 to between 780 – 790 cm-1, and

several possessed a double peak within that region. This shift is due mainly to the

substitution of carbon atoms into the icosahedron. As a heavier atom substitutes into a

bond, the energy required to excite that vibration will decrease, and the IR band will shift

to a lower wavenumber.124 In the case of boron carbide specifically, Lazzari et al. have

calculated the approximate locations of the IR bands for B11C icosahedra, portrayed in

Figure 5.16.137 Thus, it would seem that the jet milled boron carbide powders have at

least two differing stoichiometries, resulting in a double peak for the amorphous band

near 800 cm-1.

115

(a)

(b)

116

The iteratively jet milled boron carbide powders were expected to show

increasing intensity of the characteristic amorphization band. Spectra taken on

unprocessed boron carbide powders after being jet milled once, twice, and three times are

shown in Figure 5.17. As expected, the powders do not experience any shift in the bands

at 1560 cm-1 or 1080 cm-1, indicating there is no large scale stoichiometry change

117 induced merely by jet milling the powders. However, the double peak below 800 cm-1 indicative of amorphization by the icosahedra does become slightly more relatively intense by the end of three rounds of jet milling. The intensity difference is not statistically significant, but it may be a sign of minor compounding of amorphization through iterative jet milling. The FTIR spectrum of a commercial powder before and after jet milling is shown in Figure 5.18. This powder also showed a propensity for amorphization following jet milling, as indicated by the appearance of the double peak below 800 cm-1.

118

(a)

(b)

119

(a)

(b)

Figure 5.18. FTIR spectra of commercial boron carbide powders before and after jet milling, (a) full view and (b) expanded view. After jet milling broad, low intensity peaks appear below 800 cm-1.

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5.2.3 Raman Microspectroscopy

5.2.3.1 Raman for Carbon Content

The first samples to be vibratory milled, the commercially available boron carbide powders, were subjected to Raman microspectroscopy as the quickest and easiest method of determining if any structural variations were being induced by comminution. The results of 10 minute integrals of vibratory milling up to 30 minutes for two commercial powders, from ESK and UK Abrasives, are shown in Figure 5.19. Both powders displayed a drastic increase in “D” and “G” peak intensity as comminution proceeded, indicating a rise in the surface presence of carbonaceous materials. This increase in carbon signal led to the in depth study on the unprocessed boron carbide materials.

G D G D

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Unprocessed boron carbides from the four sections described in Chapter 5.2.1.1

were examined by Raman microspectroscopy before and after 30 minutes of vibratory

milling. The resulting Raman spectra are shown in Figure 5.20. The visual sorting had

suggested section C would contain the most carbonaceous material, with a decreasing

trend across the segment towards section S. However, the Raman spectra before milling

clearly show section P as exhibiting purely carbonaceous materials, with section S

showing significantly higher relative intensity of “D” and “G” peaks than sections C and

CB. Following vibratory milling, however, sample P began to exhibit characteristic spectra of boron carbide with vastly reduced relative intensity of the “D” and “G” peaks, while section S exhibited no relative change in “D” and “G” peak intensities and sections

C and CB both exhibited marked increase in relative intensity of the “D” and “G” peaks.

This phenomenon can be explained by the nature of Raman microspectroscopy as

a surface characterization technique. The penetration depth of the laser utilized in Raman

microspectroscopy, and thus the extent of material characterized, is greatly dependent

upon the scattering power of the material being examined. Carbon is a strong scatterer,

with penetration depth varying proportionally with relative amount of sp3 bonds in the

carbon.149, 163, 164 For diamond, the relative content of sp3 bonds is high; for other forms

of carbon, such as graphite and amorphous carbon, the sp3 bond content is low, and thus

the laser will only have a penetration depth of 10 nm.164 Thus, a thin film of carbon upon

the surface of the P section, as suggested by the FTIR band shifts, would result in the

carbon completely obscuring the boron carbide beneath it via Raman. When examined

following vibratory milling, new surfaces beneath the thin film are exposed and the boron

carbide spectrum is once more exhibited.

122

D G

(a)

G D

(b)

G D G D

(c)

123

D G D G

(d)

The vibratory milled samples were also examined for stoichiometric carbon content via Raman. Table 5.3 lists the average locations of the 480 cm-1 and 535 cm-1 peaks, and the resultant wavenumber spacing between them. Compared with the data from Domnich et al.,133 these results support the XRD and FTIR conclusions. Although the difference in wavenumber separation between the two peaks is somewhat minimal in the range of 18 – 20 % carbon, it is still evident that the separation will decrease with decreasing carbon content. For the vibratory milled boron carbides, the separation is at a maximum in the C and S sections, with the P section having the smallest separation.

Thus, the results suggest a trend of decreasing stoichiometric carbon from the C section across the CB section into the P section, with an increase in the S section again. All the sections are in the range of 18 – 20 % carbon.

124

Table 5.3. Locations of the 480 cm-1 and 535 cm-1 peaks in samples with varying stoichiometric carbon133 and for the vibratory milled unprocessed boron carbide. Separation between the peaks suggests a decreasing stoichiometric carbon content from section C to section P, with section S having a much higher stoichiometric carbon content again. All values in cm-1. Sample 20 % C 18 % C 13 % C 11 % C C CB P S 480 cm-1 480 481 484 486 477 479 477 479 535 cm-1 532 532 530 529 530 531 528 532 Separation 52 51 46 43 53 52 51 53

5.2.3.2 Raman Analysis for Presence of Amorphization

Raman microspectroscopy on the jet milled unprocessed boron carbide powders

proved useful in determining whether amorphization had occurred during jet milling.

The Raman spectra for the jet milled powders, shown in Figure 5.21, clearly displays a

low broad peak at 1800 cm-1, indicative of the amorphization previously observed.

Unlike FTIR, the frequency of this feature was only approximately 1 out of every 7 scans

collected per sample, and it was often difficult to predict visually if an area would exhibit

the amorphization peak. Areas that would appear to be well comminuted would not

necessarily always possess the amorphization peak, and areas that appeared to not have

been reduced in particle size would occasionally have a clearly defined amorphization

peak. It seemed that visual classification was not sufficient to predict amorphization.

125

Figure 5.21. Typical Raman spectrum of unprocessed boron carbide following jet milling, clearly showing the characteristic amorphization peak at 1800 cm-1.

The occurrence of the amorphization peak in 1 out of 7 scans in seemingly random locations could have been a result of a concentration effect within the powders.

The 20x objective lens used on the Raman microscope, with its 5 μm diameter spot size, would allow analysis for tens of particles at a time with average particle size of less than

900 nm. If all the boron carbide particles possessed a very minor peak at 1800 cm-1, then

if enough of those particles were examined at one time the cumulative effect may have

given rise to the overall peak seen with the 20x objective lens. In areas where fewer

particles were within the laser spot, and the resultant cumulative relative intensity of the

1800 cm-1 peak was low, then the 20x objective lens scan would not depict an overall

amorphization peak. To determine whether this concentration effect was responsible for

the frequency of observation of the amorphization peak, an area was found on a jet milled

unprocessed boron carbide that exhibited the amorphization peak, as shown in Figure

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5.22. The magnification was increased to a 50x objective lens with a spot size approximately 2 μm in diameter, allowing for analysis of many less boron carbide particles at once. A 70 μm by 3 μm raster was performed over the area at 1 μm step size.

Figure 5.23 depicts the typical Raman spectra zoomed in on the amorphization peak from the raster, where a spot with strong intensity for the characteristic 1800 cm-1 peak is surrounded by spots with decreasing intensity of the peak and then no peak at all. This leads to the conclusion that the frequency of observation is not a concentration effect, but rather a true frequency wherein only 1 in 7 particle sets analyzed at 20x objective lens magnification experienced amorphization.

127

(a) (b)

(e) (f)

128

(g) (h)

(i) (j)

(k) (l) Figure 5.23. Raman spectra taken at 1 μm intervals across the sample from Figure 5.22, proceeding from left to right (a) – (l). Only select spectra exhibit the characteristic amorphization peak at 1800 cm-1, indicating the frequency of observation in jet milled unprocessed boron carbides is not a concentration effect.

129

Raman spectra for unprocessed boron carbide powders iteratively jet milled are shown in Figure 5.24. Although XRD and FTIR combined to depict no structural change and no evidence of increasing frequency of amorphization, it was predicted that Raman would be more sensitive to any increase as a function of iterative milling. However, it was impossible to truly determine the sensitivity of the Raman to iterative milling due to the overwhelming carbon signal present upon further rounds of jet milling which dampens out the boron carbide pattern, including the amorphization peak. The nature of sample collection, as described in Chapter 5.1.2, required that the fines filter bag material be collected for analysis. The amount of fine carbon particles in this material results in the Raman laser being unable to penetrate the coating of carbonaceous particles on nearly every boron carbide particle. Unfortunately, no conclusions as to the effect of iterative milling upon the amorphization of boron carbide can be reached via Raman microspectroscopy at this time. However, the locations of the 480 cm-1 and 535 cm-1 peaks for each round of jet milling are shown in Table 5.4. As expected, no significant shift in peak location occurred with iterative jet milling, indicating no stoichiometry change with iterative comminution, supporting results from XRD and FTIR. Thus, it would appear that the carbonaceous materials that are obscuring the patterns and preventing analysis of amorphization are indeed secondary intragranular phases and are not evolving from the boron carbide structure itself.

130

D G

Table 5.4. Peak locations for iteratively jet milled unprocessed boron carbide powders. No peak shift, and thus no stoichiometry change, is apparent as a function of iterative comminution. All values in cm-1. Milled Milled Milled Sample 20 % C 18 % C 13 % C 11 % C 1x 2x 3x 480 cm-1 480 481 484 486 475 471 479 535 cm-1 532 532 530 529 528 524 532 Separation 52 51 46 43 53 53 53

Commercially available boron carbide powders also exhibited Raman spectra that

depicted a strong amorphization peak following jet milling, as shown in Figure 5.25. The

appearance of the characteristic 1800 cm-1 peak is evident when comparing the spectra for the commercial powder before and after comminution. Even commercial powders, processed and refined for sale, appear to be subject to amorphization under comminution.

131

D G

Figure 5.25. Raman spectra of commercial boron carbide powders before and after jet milling. The characteristic amorphization peak at 1800 cm-1 is very prominent after comminution.

5.2.4 Summation of Induced Variations

The previous sections have addressed the results gleaned from the various analytical techniques separately, and begun to draw a picture of induced variations in boron carbide during comminution. This section will tie those results together into three cohesive variations observed as a result of comminution.

5.2.4.1 Emergence of Pre-existing Compositional Variations

The stoichiometry of boron carbide has nearly always been assumed to be fairly constant across the bulk of a powder. However, calculations of stoichiometries on vibratory milled boron carbide powders from XRD data displayed a shocking disparity in stoichiometry. From near the graphite electrode across the ingot to locations near the edge of the ingot, even across a segment that would have been entirely classified as good

132

boron carbide and processed for sale, there was a stoichiometry difference of up to 2 at%

carbon in the powders. FTIR data supports this conclusion by showing a difference in

network carbon across the unprocessed boron carbide segment. The change in network

carbon could be explained by varying isomorphs dependent upon the stoichiometry across the segment. Raman data also supports the XRD and FTIR results, portraying a change in stoichiometric carbon content across the segment. The variation indicated by the separation between the 480 cm-1 and 535 cm-1 peaks is on the order of 2 % carbon from the far edge of the segment to the middle, rising back up in stoichiometry in samples closer to the electrode. While this difference in stoichiometry and isomorph

likely would not make a significant enough difference in properties for the realms of

abrasives or nuclear shielding materials, for high load structural ceramics applications the

present compositional variations can have drastic consequences.

5.2.4.2 Evolution of Carbon

From optical microscopy observations, electron microscopy observations, and

Raman microspectroscopy spectra, it has become evident that carbonaceous materials are

being released from large particles of boron carbide upon comminution and proceed to

coat the boron carbide particles. The origin of this carbonaceous material can be debated

to an extent. In the vibratory milled samples, the graphite weight percent of the samples,

combined with the vibrational spectroscopy data, can lead to a conclusion of several

possible origins. The first possible carbon source is unreacted raw materials, like those

seen in section “C” of the unprocessed segment. These materials are characterized by

low graphite weight percent, as they are typically a coke, coal, or reactive carbon

133

precursor, and not graphite. Secondly, highly reacted, graphitized material from near the

electrode, characterized by high graphite weight percent, may be present. A third

possible source is the presence of carbonaceous films on the inside of pores in the

segment, created as the melt cooled and gases were trapped inside the viscous melt. As

indicated by the Ellingham Diagram in Figure 5.26a and expanded view in Figure 5.26b, carbon monoxide will preferentially transform to carbon dioxide when cooled below

700°C. These carbon films then likely arise from the deposition of carbon during the cooling of carbon monoxide gases due to the lack of oxygen to react with.

134

(a)

135

(b)

Figure 5.26. a) Ellingham Diagram portraying the free energies of several oxides,117 expanded in b) to portray the lines for reactions of carbon with oxygen as calculated from 165 equations for Gibbs free energy. Of note is the intersection of the 2C + O2 = 2CO and 2CO + O2 = 2CO2 lines near 700°C.

However, not all carbonaceous materials may be the result of preexisting intergranular carbonaceous secondary phases in the boron carbide powders. When examining the typical graphite inclusions in carbide materials, shown in Figure 5.27, it becomes readily evident that typical secondary carbon in boron carbide is of an elongated, platey morphology and oftentimes several micrometers in length. However,

magnified images of the carbonaceous materials coating the boron carbide powders

comminuted in this dissertation, shown in Figure 5.28, are possessed of a very different

morphology. These carbonaceous materials are clusters of nanoscaled material, and

136 have a relatively equi-axed morphology that would not be obtained by simple comminution of the larger platey inclusions. It may be assumed that carbonaceous materials are evolving from the structure of boron carbide during comminution, but the lack of stoichiometry change before and after comminution indicates these carbon materials were secondary phases originally. It can be inferred that some of the carbonaceous materials becoming exposed during comminution are in fact arising from intragranular inclusions, possibly of the sort noticed in Figure 2.6.

Figure 5.27. Graphitic inclusions in carbide armor materials, appearing as plate-like structures in both silicon carbide (left)166 and boron carbide (right).167

137

(a) (b)

(e) (f)

Figure 5.28. SEM micrographs, in increasing magnification from (a) to (f), portraying clusters of nanoscale, equi-axed carbonaceous material coating the larger shard like boron carbide particles.

138

5.2.4.3 Amorphization of Boron Carbide

The combination of Raman microspectroscopy and FTIR data show an effect that

has been previously labeled as pressure induced amorphization in high load situations has been induced by high energy comminution. It is currently unclear whether the material

exposed by comminution is newly transformed material, or a previously existing

intragranular amorphous phase inherent to the boron carbide structure that has never

before been identified due to the lack of high resolution analysis beneath the particle

surface. As the observed amorphous boron carbide is assumed to be highly reactive to

applied pressure, a focused ion beam (FIB) would be required to expose material inside

the boron carbide grains. Supporting the view of amorphization being a preexisting sub- surface condition in boron carbide are several works that emphasize the common knowledge that under pressure, most materials (excluding H2O) experience a phase

transition to a more highly ordered structure, not an amorphous structure.115, 116, 168-170

Also, in a material that exhibits a compositional gradient such as that shown in the ingot

segment of unprocessed boron carbide, it is common to find periodic deformations,

resembling the amorphization bands observed in boron carbide, in the crystal lattice due

to the thermal gradient across the segment during cooling.171 As there has been no

detailed atomic structural analysis of boron carbide in the unprocessed condition to this

point in time, the possibility of preexisting amorphization in boron carbide cannot be

discounted.

There is an abundance of previous studies that support the theory of

amorphization being a pressure induced transformation, in a multitude of materials.99, 110,

113, 131, 172-174 If the amorphization is viewed as a pressure induced amorphization, the

139 contact pressures involved in comminution, depending on effective geometry of impact, are well above those required for amorphization, and the vibrational spectroscopy data supports the theory of amorphization due to comminution well. The key characteristic peak of amorphized boron carbide, the peak at 1800 cm-1 in the Raman spectra, was observed in powders following jet milling, but not prior to jet milling. Although the occurrence is slightly low, being observed in roughly 15% of random sampling areas per powder, the observations are high enough to definitively conclude that there is a new peak being observed following jet milling. Similarly, FTIR data allowed a more frequent observation of the amorphous boron peak at 800 cm-1 which should be coincidental with amorphization. The higher frequency of observations, approximately 50% of samples per powder, can be explained as the FTIR samples a much larger amount of powder than the

Raman. Although very small, the 0.3 mg of powder examined in the FTIR is nearly five orders of magnitude more than the mass of the powder examined in the Raman. A spot size of approximately 5 μm in diameter would mean only roughly a nanogram of powder will be examined under Raman microspectroscopy, thus making it no surprise that the observations of amorphization are less frequent in the smaller sampling area of the

Raman. Both techniques are reliable in their demonstration of the amorphization peaks following high energy comminution.

5.3 Elimination of Induced Variation

This section will discuss the three main variations in boron carbide powders exposed by this dissertation, in order of decreasing difficulty of accommodation.

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5.3.1 Compositional Variations

The first powder variation readily exposed by this dissertation was the extreme

range of compositional, stoichiometric variation in boron carbide powders.

Unfortunately, this is also the most difficult variation to address in powder processing.

With the exception of alloying or high temperature diffusion reactions, there is virtually

no way to adjust the stoichiometry of a carbide once it has crystallized into a solid form.

While there have been efforts to examine the feasibility of mechanical alloying for

production of boron carbide from precursor materials,11, 27, 28 there has been no evidence

to suggest mechanical alloying can alter the stoichiometry of a boron carbide already in

solid form. Diffusion of elements into the boron carbide is also difficult, and often

requires temperatures in excess of 2000°C, even when utilizing substitutional metal

species for accelerated diffusion.64, 66, 175, 176

Adding to the difficulty of addressing the stoichiometric variations in boron

carbide powders is the fact that many bulk properties may not be greatly influenced by

the variation. The discussions in Chapter 2.4 lead to these conclusions: It has been

shown that stoichiometric B4C is desirable for its maximum hardness of all boron carbide

stoichiometries,53 as shown in Figure 5.29, and the elastic properties also vary with stoichiometry. However, there is little that changes with stoichiometry that would enable easy segregation of the various stoichiometries. There is no noticeable visual change in

powders with a 2-3 at% of carbon difference. Electrical properties do not alter enough to

enable a magnetic separation. Density does not vary to a degree at which Stokes settling

would be effective at eliminating undesired stoichiometries. Calculating from general

Stokes Flow equations117 for terminal velocity of a particle in a liquid medium, under

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gravity settling the difference in velocity between an 18 at% carbon and 20 at% carbon boron carbide would be only on the order of 100 nm/s. The low difference in velocity makes it practically impossible to differentiate between the two stoichiometries. The bonding in the boron carbides are very minimally changed, meaning a preferential thermal or chemical attack would likely not succeed in segregating between varied stoichiometry boron carbides. In the end, it seems that if the stoichiometry of the boron carbide is indeed an issue, it must be addressed prior to achieving a solid.

Figure 5.29. Dependency of Vickers Hardness upon stoichiometry for boron carbide. 53 Maxima is reached at stoichiometric B4C.

5.3.2 Removal of Excess Carbon

The excess carbonaceous materials evolved through comminution of boron carbide are considered to be detrimental to boron carbide’s desirable mechanical properties. This conclusion arises mainly from prior work, highlighted by McCauley et

al. and Bakas et al., 166, 167, 177 wherein large carbonaceous inclusions, such as those

illustrated in Figure 5.27, were found to be stress concentrators, and lead to premature

anomalous failure of the hard ceramics that contained them. Thus, there is strong

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motivation for removal of excess carbonaceous materials from a bulk hard ceramic. In

the case of carbides, it is difficult to predict a removal method that will not adversely

affect the bulk carbide as well. In an attempt to achieve such selective material removal,

however, this dissertation approached thermal attack of carbon in several atmospheres

and chemical oxidation of the carbonaceous materials.

5.3.2.1 Thermal Attack of Carbon

As a first attempt to remove carbon, firings in various atmospheres were attempted in order to decompose the carbon preferentially over the boron carbide. Since

carbon is extremely stable, this is extremely difficult. To determine the effect of firings

upon the boron carbide itself, thermogravimetric analysis (TGA) was performed in

various atmospheres, and shown in Figure 5.30.

Figure 5.30. TGA curves of raw and commercial boron carbides, fired to 900°C in air and nitrogen atmospheres. All samples showed significant weight gain during firing, rather than weight loss.

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This weight gain implies the boron carbide is reacting with both air and nitrogen

to form a secondary compound that is retained even after firing. The reaction initiates

near 600°C, creating difficulties for thermal oxidation of carbon when considering the

general rule of thumb is carbon will oxidize around 500°C. SEM micrographs of samples

fired in air are shown in Figure 5.31, and depict the growth of whiskers of what is most

likely boron oxide. Raman spectra taken before and after firing in air are shown in

Figure 5.32, which, compared with the Raman spectrum of boron acid in Figure 5.33,

indicates a growth of surface oxide on the particles as the new major peaks after firing.

The peaks located at 500 cm-1 and 880 cm-1 are the characteristic peaks of boron acid, indicating a growth of oxide that then hydrates to the acid form. The secondary peak at

-1 178 800 cm supports this conclusion as it is characteristic of boron oxide of B2O3.

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Figure 5.31. SEM micrographs of boron carbide powders before (top) and after (bottom) firing in air to 900°C. The powders have become completely enveloped by the growth of oxide whiskers.

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Figure 5.33. Characteristic Raman spectrum of boric acid.179

In an attempt to remove this oxide layer, the fired in air samples were subjected to vibratory milling to break down aggregates and washing in methanol and hydrochloric acid. The resulting Raman spectra from each of these tests, shown alongside the original stock powder and fired in air powder spectra in Figure 5.32, indicate with very strong

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“D” and “G” peaks that carbonaceous material has not been removed effectively from the

powders. Work done by Nazarchuk et al.162 suggests that the oxidation of the boron

carbide has created a passivation layer on the surface of the material, preventing further

oxidation of any boron carbide or carbon still remaining. Sogabe et al. and Steinbruck et

al. concurred with this conclusion, demonstrating the boron oxide layer would indeed

form a protective coating around residual surface carbon.43, 180 Thus, it can be recognized

that attempts to thermally attack carbon in air will be ineffective, although perhaps

different atmospheres would be more successful.

Thermal attack of carbon was then attempted in a variety of inert firing

atmospheres to prevent growth of the oxide on the boron carbide powders. Initial firings

were to 900°C in atmospheres of argon, nitrogen, and a mixture of 15% hydrogen in argon. The hydrogen/argon atmosphere was thought to have the most significant capacity for reacting with the carbon to form methane under heating. Powders were also fired to 1700°C in a high vacuum of 10-4 Torr in a loose powder bed in a graphite

crucible. It was hoped that the carbon present, as it was not a graphitic carbon, would be

more reactive and sublime under heating. Heating to this temperature initiated pre- sintering, however, and some powder close to the crucible was aggregated into compacted segments. The Raman spectra of powders taken from all atmosphere firings are shown in Figure 5.34. As feared, the inert atmospheres, even the hydrogen/argon gas atmosphere, achieved no significant carbon removal after firing, as indicated by the high intensities of the “D” and “G” peaks following firing. These results clearly show that thermal attack of carbon is not an effective means of carbon removal from boron carbide.

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Figure 5.34. Raman spectra of boron carbide powders subjected to various firing atmospheres. All atmospheres resulted in no significant selective oxidation of carbon for carbon removal.

5.3.2.2 Chemical Oxidation of Carbon

Chemical oxidation of carbon proved to be even more difficult than thermal attack of the carbon. Although the literature had presented several chemical cocktails purported to preferentially oxidize carbon over the host boron carbide,158-162 there were doubts as to whether any chemical would truly preferentially oxidize to such a degree so as to not break down the boron carbide structure. The most promising chemical system based on the literature was that suggested by Nazarchuk et al.,161 consisting of a mixture of potassium dichromate, perchloric acid, sulfuric acid, and nitric acid. Care had to be taken due to the risk of explosion with this mixture, and the nature of chromates as severe carcinogens.181 Systems investigated included mixtures of nitric acid and peroxide;

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potassium manganate and sulfuric acid; and the Nazarchuk mixture of potassium

bichromate, perchloric acid, sulfuric acid, and nitric acid.

To test the efficacy of the chemical systems, tests were first run separately on half

a gram of graphite powder and carbon black in each of the systems. The reactivity

difference between graphite and carbon black was thought to impact the efficacy of the

oxidizing chemical systems. However, after boiling in the chemical systems for half an

hour then filtering, washing, and drying the residual powders, it was found that

difficulties were present with each system. The sodium hydroxide and hydrogen

peroxide mixture achieved no statistically significant weight reduction. The mixture of

sulfuric acid and potassium manganate resulted in formation of manganate oxide

precipitates that were impossible to wash away from the carbon powders. The presence

of such precipitates required the careful use of the original Nazarchuk solution, which

also achieved no statistically significant weight loss. Still, samples were taken from the

Nazarchuk solution for further vibrational spectroscopy analysis.

FTIR spectra taken of powders before and after washing in the Nazarchuk

solution are shown in Figure 5.35. Of note here is the significant increase in intensity of

the absorption bands near 1030 cm-1, indicative of the growth of boron oxide,182 and the

bands ranging from 2850-2950 cm-1 which are characteristic of carbon-hydrogen bonds.

The Raman spectra of several samples of boron carbide powder following half an hour boil in the Nazarchuk mixture, filtering, washing, and drying, are shown in Figure 5.36.

As can be clearly seen by the continued presence of the “D” and “G” peaks in the spectra, carbon has not been efficiently removed via chemical oxidation. Still, there is a

noticeable trend of reduced “D” and “G” band intensity. This trend is offset, however, by

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the appearance of a small peak at 880 cm-1, indicating oxidation of the boron carbide has been initiated as well. Both vibrational spectroscopies therefore indicate that the

Nazarchuk solution has not preferentially oxidized the carbonaceous material, but rather

has begun oxidation of the boron carbide as well and provided for adsorption of hydrogen

to the carbon. As mentioned in Chapter 5.3.2.1, oxidation of the boron carbide will

compete with oxidation of the carbon, and provide a passivation layer that will limit

further oxidation of carbonaceous materials. Thus, even chemical oxidation for removal

of carbonaceous materials is not very effective. It would seem that any attempt at

eliminating excess carbonaceous materials must be addressed prior to solidification of the

boron carbide.

Figure 5.35. FTIR spectra of boron carbide before and after chemical oxidation in Nazarchuk’s solution.161 Increase of intensity of bands near 1030 cm-1 and 2800-2900 cm-1 are indicative of growth of boron and carbon oxides, respectively.

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5.3.3 Addressing the Amorphization

As the amorphization of boron carbide is considered by industry to be the limiting factor in several applications, limiting and eliminating it are of key importance in discussing future applicability of boron carbide. This section will address two possible methods examined for limiting amorphization, and then discuss the reversibility of amorphization.

5.3.3.1 Orientation Dependence of Amorphization

As mentioned in Chapter 2.4.3, boron carbide has a rather pronounced anisotropy of elastic properties. As such, it can be presumed that orientation of grains during

151 pressure loading should have a marked effect upon activation of deformation mechanisms, including amorphization. Also, previous work mentioned in Chapter 2.6 depicted a strong orientation dependence of the amorphization in both ballistic impacted and nanoindented boron carbide.99-101, 103, 106, 107, 131, 132, 135 These factors combine to give rise to the theory that the frequency of observation of amorphization in jet milled boron carbide may be affected by orientation of the particles being examined.

To investigate this possibility of orientation dependence of amorphization, samples were fabricated with induced texture via capillary extrusion. Commercially available micron-scale boron carbide with an aspect ratio of approximately 10 to 1 was used to provide a baseline; samples were also produced from the unprocessed boron carbide following one round of jet milling to obtain textured amorphized material.

Following extrusion, extrudates were subjected to XRD for determination of the texture coefficient, TC.

For the extruded boron carbide powders, it was found that a preferred orientation did indeed exist. Table 5.5 depicts several planes and their TC values, with the (101) showing by far the highest TC for both samples. Unsurprisingly, the higher aspect ratio commercial powder (~ 10:1) was more highly oriented than the lower aspect ratio comminuted powder (~ 4:1). Still, the level of induced texture in the extrudates implies that if there is an orientation dependence of amorphization, then these samples should be able to illuminate such a dependence.

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Table 5.5. Expected and observed intensity values and TC values for several planes of boron carbide extrudates. Both samples show relatively strong orientation of the (101) plane parallel to the cross section of the rod. hkl 101 003 012 110 104 021 113

I0 7.5 24.6 57.0 11.1 60.9 100.0 5.0 Commercial I 13721 21132 45070 7951 41634 65120 3818 TC 2.03 0.96 0.88 0.80 0.76 0.72 0.85 Comminuted I 13430 23540 45979 12001 62646 68906 6515 TC 1.64 0.87 0.74 0.99 0.94 0.63 1.19

Due to the shape of the rods, FTIR could not be performed on the extruded

samples. Raman microspectroscopy was possible by wrapping the end of an extrudate in

scotch tape, then folding the scotch tape to make a flat blade. This flat blade was

mounted in a razor blade handle whose base had 8 flat edges around the handle, enabling

rotation of the sample at a consistent 45°. For comparison with the spectra to be taken on

the extrudates, the Raman spectrum of a typical methylcellulose, taken with a laser at

wavelength 785 nm, is shown in Figure 5.37b.183 The peak at 890 cm-1 is of note as it will be evident in the Raman spectra of the extruded boron carbides, albeit shifted to slightly lower wavenumber due to the shorter wavelength of the experimental laser, and should not be confused with a characteristic boron carbide peak.

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Representative Raman spectra from 0° through 180° are shown in Figure 5.38 for the commercial platey powder and the comminuted powder. Interestingly, the characteristic amorphization peak at 1800 cm-1 was observed minimally at 90° and more intensely at 180° rotation, yet not at any other rotation angle. Once again, the observation was approximately one in seven scans at these angles. This angular dependence depicts a clear orientation dependence of the amorphization, as expected. While an extremely highly textured sample may display more dependency of the amorphization, even this level of induced texture is enough to influence the observations of amorphization.

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(a)

(b)

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5.3.3.2 Low Pressure Jet Milling

As mentioned in Chapter 2.8, the impact pressure of jet milling is on the order of

several GPa or higher, depending upon geometry of impact. Indeed, it is likely that the

amorphization seen, having shown no obvious orientation dependence, thus is most likely

induced through a combination effect of achieving the optimal geometries for maximum

impact pressure and possibly multiple impacts per particle during milling. As such, it

appears at present that the “standard” milling pressure, 100 psi air pressure in the case of

this dissertation and the 2 in. Sturtevant Micronizer in question, is sufficient to readily

achieve localized pressures during impact of such magnitude that amorphization is

induced. In order to determine a lower limit to this amorphization in hopes of enabling

jet milling for comminution without inducing amorphization, two further series of

samples were run at lower grinding air pressures.

Grinding and feed air pressures used in the low pressure jet milling were held

equal to each other, lowered for the first samples to 80 psi and 60 psi for the second

series. Following the calculations of Chapter 2.8, approximate contact pressures are

shown in Table 5.6 for multiple contact geometries. Assuming both particles under

impact to be of equal particle size of 45 μm, D1 is taken to be the effective particle size

of a protrusion on the surface of a particle leading to a smaller contact area upon impact.

Of note is the fact that for submicron effective particle size, contact pressures are still

well into the 300+ GPa range, well above the observed 15-40 GPa onset of

amorphization.99, 107 60 psi is approximately the low limit on grinding pressure for

successful operation of the Micronizer. SEM of the two low pressure grindings are shown in Figure 5.39. The powders clearly have not reduced particle size nearly to the

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level of the higher pressure jet milling, yet still have clear indications of brittle fracture in

the form of steps along the face of the powders. Also, the finer ground carbonaceous

material continues to coat the powders even at reduced grinding pressure. Average

particle size of the materials appears to be approximately 30-100 μm at 60 psi grinding

pressure and 10 μm at 80 psi grinding pressure. Of note is the rounded features present in

both micrographs, indicative of abrasive wearing of the boron carbide particles as

opposed to brittle fracture.

Table 5.6. Calculated contact pressures for several effective particle sizes at differing grinding air pressures. D1 is the effective particle size for an impacting protrusion on the surface of a larger host particle. Host particle diameter and D2, particle size of the impacted particle, are assumed to be 45 μm. D1 D2 Air Pressure Contact Area Contact Pressure (μm) (μm) (psi) (m2) (GPa) 45 45 100 7.11E-13 1.54 0.05 45 100 1.38E-15 791.97 45 45 80 6.13E-13 1.43 0.05 45 80 1.75E-15 501.27 45 45 60 5.06E-13 1.30 0.05 45 60 1.94E-15 338.31

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Raman microspectroscopy of the low grinding pressure samples was hoped to

show little to no amorphization. However, typical spectra for both samples, shown in

Figure 5.40, have clearly demonstrated the emergence of the 1800 cm-1 peak indicative of

amorphization. These samples exhibited the 1800 cm-1 peak at a rate consistent with

those exhibited in the regular, higher grinding pressure samples discussed previously.

Approximately one in seven scans, randomly located on the sample, displayed the amorphization peak. Similarly, FTIR spectra of the two sample series, shown in Figure

5.41, also still exhibit a small peak at 800 cm-1, supporting the conclusion that

amorphization has indeed occurred in these materials. The emergence of the peak at

1800 cm-1 in Raman spectra and the low peak at 800 cm-1 in FTIR is proof that even at lower contact pressures, comminution via jet milling inherently possesses the capability to induce amorphization in boron carbide. The lack of significant change in frequency of observation of amorphization is compelling evidence that the pressure induced transformation is triggered by a pressure not greatly affected by the grinding air pressure,

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and truly is likely an effect of geometry of impact creating extremely large localized applied pressures. Interestingly, the samples ground at 80 psi grinding pressure had a slightly lower relative intensity of amorphization peaks in both Raman and FTIR spectra.

This may be yet more indication of the occurrence of amorphization being more

dependent upon a relatively random variable, such as geometry of impact, and the change

in grinding pressure does not significantly affect amorphization. The higher presence of

rounded particles in SEM images of the 80 psi grinding pressure samples supports a

scenario wherein high aspect ratio impacts are less common than a sphere on sphere

geometry, as well.

159

(a)

(b)

160

(a)

(b)

Figure 5.41. FTIR spectra taken on samples jet milled at 60 psi (a) and 80 psi (b) grinding pressure. Both samples exhibit evidence of a low peak at 800 cm-1, indicative of amorphization.

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5.3.3.3 Reversibility of Amorphization

In any induced transformation, one of the first questions asked about it is, “Is it

reversible?” A transformation can be limited in its impact upon functionality of the

material due to its ease of reversibility. In the case of amorphization in boron carbide,

present thinking is that the pressure induced transformation possesses a negative impact upon high strain rate applications. Thus, full and easy reversibility is desired. Prior work showed reversibility upon annealing of boron carbide, so powders that had demonstrated amorphization were subjected to annealing for one hour at 600°C. Powders were then examined using Raman microspectroscopy, with before and after annealing spectra portrayed in Figure 5.42. As expected from the literature,100, 101 the spectra observed

exhibited no peak at 1800 cm-1. Similarly, the “G” peak at 1570 cm-1 grew extremely

pronounced, and the “D” peak at 1330 cm-1 also increased in intensity, albeit to a lesser degree than the “G” peak.

While initially this damping, and indeed apparent total elimination, of the characteristic peak at 1800 cm-1 can lead to the assumption that the amorphization has been effectively reversed via annealing, one must be careful in leaping to that conclusion.

It is highly unlikely that carbonaceous materials have recrystallized at as low a

temperature as 600°C, with work in the literature expecting recrystallization only at

temperatures exceeding 1700°C for amorphous boron carbide.184 One possible

explanation would follow Yan’s model of amorphization, discussed in Chapter 2.6.3,

wherein bending of the intericosahedral chain results in aromatic carbonaceous rings forming from the chain upon unloading.107 As such, the increase in the “G” peak, and to

an extent the “D” peak, and the elimination of the characteristic peak at 1800 cm-1, would

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result from coarsening of the aromatic carbonaceous rings. As the low annealing

temperature drives coarsening, the aromatic rings would coarsen to a degree where the

effective hybrid sp2 bonding would no longer be the dominant bonding mechanism, and

sp3 bonding would become dominant. However, this explanation is dependent upon the

assumption that the “D” and “G” peaks in pure boron carbide are exactly those of carbon.

This is not the case, as mentioned in Chapter 2.9, but rather the new peaks after low

temperature annealing appear to be the higher wavenumber carbon “D” and “G” peaks, as

opposed to the slightly lower wavenumber boron carbide peaks. This assumption’s

invalidity does not necessarily also invalidate the explanation for reversal of the

characteristic 1800 cm-1 peak, but it does imply that a simple coarsening mechanism is

not the only mechanism at play. Without fully understanding what causes the

amorphization to begin with, it is difficult to comprehend the quasi-reversibility observed via Raman microspectroscopy. What is understood is that at relatively low temperatures, the single most characteristic analytical feature of amorphized boron carbide has been eliminated.

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(a)

(b) Figure 5.42. Raman spectra of jet milled boron carbide powders before (a) and after (b) annealing at 600°C. The characteristic peak at 1800 cm-1 has been effectively eliminated from the spectra, and the relative intensity of the “D” and “G” peaks are markedly increased.

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6. Conclusions

Boron carbides of various origins were subjected to several energies of comminution in an attempt to determine the extent of microstructural variations such as carbon inclusions and amorphous regions that are induced or exposed by powder processing alone. Powders investigated included commercially available boron carbide and unprocessed boron carbide, gleaned directly from the ingot following carbothermic reduction and representative of a range of boron carbides. The combined analyses from

XRD, FTIR, and Raman microspectroscopy provided for a number of conclusions:

1. Boron carbide, as produced via carbothermic reduction, exhibits a wide range of stoichiometries. This range exists entirely within material that, when gleaned from the unprocessed ingot, are all classified as “good” boron carbide, yet might not be optimal materials depending on desired application.

1a. These stoichiometric variations are relatively wide, with a difference in excess of

2 at% carbon between extrema determined by XRD lattice constant calculations and

Raman microspectroscopy peak position determination. Shifts in the FTIR peak locations support the stoichiometric variations concluded from XRD and Raman. Prior work has shown a strong dependence of bulk properties upon stoichiometry, so powders with a variety of stoichiometries will then lead to variable bulk properties in any body produced from those powders.

1b. The presence of stoichiometric variations is difficult, if not impossible, to correct once the boron carbide is in solid form. There is a lack of readily available literature on controlling stoichiometry after solidification, and general knowledge states such an

165 endeavor is impractical. Stoichiometric control of boron carbide must be addressed prior to solidification.

2. Carbonaceous materials of various origins on and within the boron carbide particles are exposed by comminution. Upon exposure to the surface, the carbonaceous materials proceed to coat the surface of boron carbide particles, and are extremely difficult to remove by thermal attack in various atmospheres and by chemical attack.

2a. Analyses led to the assignment of at least three preexisting carbonaceous materials that are exposed upon comminution: unreacted precursor carbon; carbon thin films deposited during cooling of CO gases trapped within the highly viscous carbothermic reduction melt; and graphitic material either spalled from the arc furnace electrodes or produced by overheating of the melt close to the electrode. There is also the possibility of nanoscale intragranular carbonaceous material that is exposed upon comminution. Intragranular carbon has been observed for years in bulk boron carbide bodies, and can be assumed to be present in the powders.

2b. Thermal attack of carbonaceous materials is ineffective at best. TGA curves displayed boron carbide has a tendency to oxidize at temperatures near those of carbon, and firings in air only grew a passivation layer of boron oxide over the carbonaceous materials. Firings in various atmospheres achieved no significant removal of the carbonaceous materials.

2c. Chemical attack of carbon is equally ineffective. Put generically, any chemical strong enough to oxidize carbon is also strong enough to oxidize boron carbide. Nitric acid mixed with peroxide and potassium manganate mixed with sulfuric acid both achieved no significant carbon loss. Even the Nazarchuk solution resulted in oxidation of

166 the boron carbide and a significant carbon presence was remaining following chemical attack. Thus, it appears full removal of excess carbon may only be achievable when addressed before solidification of the boron carbide.

3. Amorphized boron carbide is readily exposed by high energy comminution.

Unfortunately, iterative comminution proved incapable of examining compounded exposure of amorphized material due to overwhelming presence of carbonaceous materials.

3a. The origin of the amorphized boron carbide cannot be conclusively determined by this study. It is possible that it either was a preexisting intragranular feature that is only exposed to the surface upon particle fracture, or it may be a pressure induced transformation.

3b. FTIR results portrayed amorphous boron carbide existing in roughly half of the materials scanned following jet milling. The double peak observed where one peak representing a B12 icosahedron is indicative of carbon atoms substituting into the icosahedron.

3c. Raman microspectroscopy depicted amorphous boron carbide in only 1 in 7 materials scanned following jet milling. The reduced sample size compared to FTIR accounts for this change in observation frequency.

3d. Raman mapping of an area of boron carbide that included a spot observation of amorphized boron carbide eliminated the possibility of the frequency of observation being influenced by a concentration effect. This concentration effect would involve low quantities of amorphization existing in all materials, yet at the magnification used, only strong concentrations would result in an area averaged scan that exhibits the

167 characteristic amorphization peak. Points separated by 1 μm, when scanned at a higher magnification and smaller spot size than initial scans, displayed markedly different presence of the characteristic amorphization peak, with many points exhibiting no amorphization peak at all. Amorphization is not present in all boron carbide materials that have been comminuted.

3e. There is an orientation dependence of amorphized boron carbide. Extrudates formed from comminuted boron carbide exhibited a slight texture. Raman microspectroscopy of these extrudates only observed the characteristic amorphization peak at 90° and 180° about the rotation axis from a random point. The orientation results need to be further explored, but there is definitive orientation dependence observed.

3f. High energy comminution through jet milling, between grinding air pressures of

60 and 100 psi, will always enable exposure of amorphized boron carbide. If amorphization is preexisting, then jet milling will fracture the particle through the carbon inclusions, exposing surfaces exhibiting the amorphization. If amorphization is a pressure-induced transformation, then depending upon the geometry of impact, jet milling at low grinding pressures still possesses enough contact pressure to induce amorphization.

3g. The characteristic amorphization peak in Raman spectra of boron carbides can be eliminated by low temperature annealing. While it seems unlikely that this is a recrystallization of the amorphized material, it may be a result of coarsening of the amorphized bands. True reversibility of amorphization is likely not possible unless annealed at high temperatures in excess of 1700°C.

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7. Future Work

1. Carbon/structural analysis of commercial boron carbide: Vast amounts of free carbon have been shown to be released during powder processing by this dissertation.

However, there is still a large amount of uncertainty as to the true origin of this free carbon, and its effect upon the atomic scale structure of boron carbide. It is recommended that a study be undertaken to conduct high resolution TEM with EELS upon commercially available boron carbide. This study would focus upon the determining the extent of carbon inclusions in boron carbide and what effect, if any, these

inclusions have on the stoichiometry and structure of boron carbide in contact with an

inclusion. Kinetic and thermodynamic theories state that there should be a certain

amount of diffusion from a concentrated source of one atom (free carbon inclusions) into

a surrounding matrix with a relatively lower concentration of that atom (boron carbide).

This is countermanded by the observation of thermodynamic stability of a boron-rich

boron carbide phase preferentially to a carbon-rich boron carbide. If diffusion is a

stronger motivational force in boron carbide than the thermodynamic stability of boron-

rich boron carbide, EELS should be able to depict a diffusional transition between the

carbon inclusion and the bulk boron carbide. It has been proposed that this diffusional boundary is in fact the amorphous material seen after ballistic testing. In this proposal, the inclusions would serve as fracture origins and the surfaces exposed after fracture

would then be a diffusional boundary, not a material that has undergone a structural

transition. A proper TEM study should be able to determine whether or not this is true,

and the diffusional boundary can be mistaken for amorphous boron carbide. As an

additional benefit, an in-depth TEM study such as this can also be utilized to examine the

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structure of boron carbide, and attempt to determine the nature and extent of isomorphs

present in the bulk boron carbide crystal. This study would confirm whether the

observed amorphization was a pressure-induced event as previously assumed or present

as a result of the chemical synthesis of the grain. As such, the study is critical to the

future of boron carbide and needs to be addressed.

2. Development of carbon-free boron carbide and proof thereof; effect upon bulk

properties: Even after looking at the structural effect of having carbon inclusions within

boron carbide, the true impact of said inclusions upon bulk properties will still have not

been fully examined. At present, it is assumed, through experience with similar

materials, that carbon inclusions are detrimental to boron carbide’s truly desirable

properties, most specifically to mechanical properties. To examine this in full, “pure”

boron carbide must be produced without any free carbon inclusions. At present, this may

best be approached through examination of the rapid carbothermal reduction method first

utilized at Dow in the 1980’s and currently under construction at Rutgers University.

Rapid carbothermal reduction possesses great promise for at the very least laboratory

scale production of “clean” boron carbide, stoichiometrically pure at B4C with no free carbon inclusions. The powders produced through the rapid carbothermal reduction experiments must first be examined in-depth through TEM to ensure that they truly are inclusion-free. Following verification of purity, bulk samples will have to be made with the powders to enable mechanical testing. From a standpoint of comparing a pure boron carbide to one with a significant secondary phase present in carbon inclusions, mechanical testing should show a significant difference in properties between commercial boron carbide and “pure” boron carbide.

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3. TEM of comminuted materials to confirm amorphization and reversibility:

This dissertation has asserted through use of proven corollary methods (vibrational

spectroscopy) that comminution of boron carbide results in amorphization. However,

there is certainly room for visual proof of the amorphization. A TEM study looking for

evidence of amorphization in comminuted powders would be relatively simple in scope,

although time consuming. The final result would be full, definitive, incontrovertible

proof of amorphization in boron carbide as a result of comminution. Also, TEM would

enable examination of annealed samples to determine what process is truly occurring to

dampen the 1800 cm-1 peak in Raman spectra of amorphized boron carbide during low temperature annealing. High temperature annealing leading to full recrystallization of the amorphous material can also be examined through detailed TEM studies.

4. Comminution of carbon-free boron carbide to show lack of carbon evolution and amorphization: Combining elements of Future Work Items 1, 2, and 3, calls to mind two intriguing factors- first, carbon evolution is only assumed to be from carbon inclusions in boron carbide. Comminution of the pure boron carbide powders created in

Future Work Item 2 will quickly show whether or not carbon evolution is still present with pure powders. It is expected that there will be no carbon evolution, as there are no inclusions to break down; however, if there is, then it implies that the physical structure of boron carbide is breaking down during comminution and carbon is being released from the material itself, opening up fundamental questions about the true stability of boron carbide. The second intriguing factor would be what exactly is the cause of amorphization, the “weak link,” so to speak, that breaks down during comminution?

Prior theoretical work of Fanchini et al. 104 and Yan et al. 107 has suggested that the CCC

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chain in the B12(CCC) isomorph is what breaks down during high pressure applications,

resulting in an amorphous band. There has also been a suggestion that the concentration

of stress around a carbon inclusion assists this transformation. If B12(CCC) is also more

prevalent near a carbon inclusion, then a theoretical picture can be painted of carbon

inclusions and non-stoichiometrically pure boron carbide as part of the origination of amorphization. If such is the case, then a pure boron carbide with no carbon inclusions, such as those created in Future Work Item 2, would possibly experience little to no amorphization under high load applications. If nanoindentation and comminution of these powders shows no evidence of amorphization through vibrational spectroscopy and

TEM, the production and application of boron carbide may experience a significant impetus for change.

5. Temperature-controlled comminution of boron carbide: This dissertation

showed a marked correlation between comminution and appearance of structural changes

in boron carbide. However, there may be room for examination of comminution of boron

carbide under non-standard conditions. Experiments were conducted at room

temperature for this dissertation, imparting a certain level of thermal energy to the

comminution. It would be interesting to observe whether altered temperatures would

affect the observed structural changes. Cryogenic milling, often used to comminute

plastic materials that would decompose through room temperature milling, might be able

to “lock in” the boron carbide structure and prevent structural changes during milling.

High temperature steam-powered milling may provide enough excess thermal energy to

enable recovery from the structural changes. One, both, or neither of these techniques

may prove effective at limiting the occurrence of structural changes in boron carbide due

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to comminution, and it would be very interesting to observe their true impact upon the

structure.

6. Full modeling of the impact energies and pressures of comminution: This

dissertation has made several assumptions to enable inclusion of impact pressures into the

discussion of the effect of comminution on particle structure. However, it is fully

recognized that these assumptions do not resemble a real-world material or situation. As such, a detailed effort should be made to model the kinetics of a multi-planar, cyclonic jet milling apparatus. Factors taken into account would include the “true” morphology of particles, as the assumed spherical shapes are much closer to a shard-like morphology with an aspect ratio in the range of 3:1. The opposing air pressures of grinding air and feed air must be accounted for instead of assuming acceleration by only feed air pressure, as well as the rotational velocity of the particles as opposed to a linear velocity. Air pressure differentials across the particles and non-laminar air flow would further complicate the model. Finally, the true nature of impact for two shards must be accounted for by examining edge on plate and point on plate impacts as the most likely impact orientations. This model would then give a more realistic, complete determination of the kinetics and resultant applied impact pressures in comminution.

7. Detailed modeling and description of the Raman spectrum of boron carbide:

As has been mentioned throughout this dissertation, the Raman spectrum of boron

carbide still is quite the mystery. The originations of many features in the spectrum are

not well understood, and often times theories pertaining to the origin of features are

contradictory and even mutually exclusive. While there have been many attempts to

systematically identify every feature of the spectrum, much of the confusion arises from

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the nature of the not well understood crystal structure of boron carbide. The substitution

of boron and carbon atoms interchangeably in the end site of the three atom

intericosahedral chain or within a polar site of the icosahedron itself can result in

modified selection rules differing from what has been expected and calculated for

stoichiometric B4C. Thus, features at various wavenumbers can arise from any of the multitude of boron carbide isomorphs, and each isomorph can have its own set of selection rules. A highly detailed, multiple excitation wavelength, multiple applied pressure and temperature study should be undertaken to attempt to specify what bonds give rise to each feature of the Raman spectrum. Of great boon to this endeavor would be the use of isomorphically pure single crystals; however, as of this time, it is unknown if such materials are even able to be fabricated.

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Appendix: X-ray Diffraction Analysis of Boron Carbide

When conducting XRD on powders during this dissertation, a number of

challenges were encountered to effectively analyzing boron carbide. This appendix will

address several of these in a brief overview.

Boron carbide is one of several materials that are extremely x-ray transparent. As

such, use of a packed boron carbide sample in a cup style diffraction sample holder will

result in scans with significant peak shift and create difficulty in accurately analyzing the

sample. As such, a monolayer of powder should be placed on a zero-background sample holder, such as a single crystal of silicon as available from a number of commercial x-ray

diffraction supply companies. A monolayer of powder will necessarily limit the

crystallographic planes under incident x-rays, so a sample spinning stage should be used to rotate the sample during the experiment and increase the number of planes brought into diffraction position. These simple steps will help ensure an accurate diffraction scan is obtained.154

Analysis of boron carbide can also depend greatly upon the powder diffraction

file (PDF) card used to compare the experimental scan. At Rutgers, there is access only

to the ICSD database of PDF cards. While overall a useful database, the ICSD does not

have well-maintained cards for boron carbide. The result is a phase identification that

can result as shown in Figure A1: the peak locations and intensities appear to mark the

experimental scan as B13C2 and not B4C. An examination of the full card files with notes,

as shown in Figures A2-A5, shows that the B4C PDF cards contained within the ICSD are

from the 1943 work on single crystals of boron carbide. A knowledge of the experiments

shows that these samples were not spun, resulting in preferential plane orientation in

175 diffraction position with greatly skewed peak intensities and minorly offset peak locations. The B13C2 cards are then similar enough to theoretical B4C patterns that a computerized analysis will ignore the ICSD B4C cards and label the material as B13C2.

This problem is most easily avoided by using a PDF database other than the ICSD; if such is not possible, the user must manually determine an appropriate B4C card to utilize in analysis and attempt to correct for the skewed card during refinement.

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Figure A1. Typical boron carbide diffraction pattern (top) with several relevant ICSD

PDF cards showing the propensity for patterns to resemble B13C2 over B4C.

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Figure A2. PDF card notes for B4C card file 97-065-4971.

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Figure A3. PDF card notes for B4C card file 97-002-9093.

179

Figure A4. PDF card notes for B13C2 card file 97-000-0446.

180

Figure A5. PDF card notes for B13C2 card file 97-061-2568.

181

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Curriculum Vita

Daniel W. Maiorano

Education:

2011 Ph.D., Materials Science and Engineering, Rutgers University

2011 M.S., Materials Science and Engineering, Rutgers University

2005 B.S., Ceramic and Materials Engineering, Rutgers University

Work Experience:

2005 – 2011 Graduate Research Assistant, Department of Materials Science and Engineering, Rutgers University

Publications:

D. W. Maiorano, D. E. Niesz, R. A. Haber, “Analysis of Texture in Controlled Shear Processed Boron Carbide.” Proceedings of the 30th International Conference on Advanced Ceramics and Composites. (2006)

D. W. Maiorano, N. Venugopal, R. A. Haber, “The Effect of Soluble Sulfate Concentration on the Rheological Behavior of Nanoparticulate Titania Suspensions.” Journal of Ceramic Processing Research. 4 [8] 266-270 (2007).

N. Venugopal, D. W. Maiorano, R. A. Haber, “Effects of Starting Powder Characteristics on Bulk Assembly of Titania.” International Journal of Applied Ceramic Technology. 4 [6] 541-548 (2007).

V. Domnich, D. W. Maiorano, R. A. Haber. “Compositional Variations of B4C Structural Parameters by Raman Spectroscopy and X-Ray Diffraction.” Physical Review B. Submitted. (2011)