Experimental observations of amorphization in stoichiometric and boron-rich boron carbide

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Experimental observations of amorphization in stoichiometric and boron-rich boron carbide Ankur Chauhan , Mark C. Schaefer , Richard A. Haber , Kevin J. Hemker PII: DOI: Reference:

S1359-6454(19)30642-1 https://doi.org/10.1016/j.actamat.2019.09.052 AM 15561

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Acta Materialia

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27 June 2019 19 September 2019 27 September 2019

Please cite this article as: Ankur Chauhan , Mark C. Schaefer , Richard A. Haber , Kevin J. Hemker , Experimental observations of amorphization in stoichiometric and boron-rich boron carbide, Acta Materialia (2019), doi: https://doi.org/10.1016/j.actamat.2019.09.052

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Experimental observations of amorphization in stoichiometric and boron-rich boron carbide Ankur Chauhana,b, Mark C. Schaeferc, Richard A. Haberc, Kevin J. Hemkera,b* a

Department of Mechanical Engineering, The Johns Hopkins University, Baltimore, MD, 21218,

USA b

Hopkins Extreme Materials Institute, The Johns Hopkins University, Baltimore, MD 21218,

USA c

Department of Materials Science and Engineering, Rutgers, The State University of New Jersey,

Piscataway, NJ 08854, USA Corresponding author: *[email protected] Abstract Boron carbide is extremely hard but has been shown to undergo stress-induced amorphization when subjected to large nonhydrostatic stresses. This localized amorphization has been associated with the sudden loss of shear strength and poor ballistic performance. Recent quantum mechanics predictions suggest that boron-enrichment may be used to mitigate amorphization in boron carbide. As a means to test this hypothesis, stoichiometric boron carbide (nominally B4C) and a novel composition of B-rich boron carbide (nominally B6.3C) were investigated. Nanoindentation followed by Raman spectroscopy revealed an obvious reduction in the Raman peaks associated with amorphization in the B-rich material. Transmission electron microscopy observations of the region below the nanoindents facilitated direct observation of amorphization, confirmed the Raman finding that amorphization is reduced in the B-rich specimens, and provided additional insight into deformation mechanisms. It is surmised that boron-rich alloys offer a path to reducing local amorphization in boron carbide.

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Keywords Boron carbide; Amorphization; Nanoindentation; TEM; Raman spectroscopy 1. Introduction and background Boron carbide is an ideal candidate material for protective body armor application due to the combination of its unique properties such as superior hardness, high hugonoit elastic limit and low density. This combination of properties is a result of its unique lattice structure with high inter-atomic electron density and strong covalent bonds. Belonging to a R3m space group, the boron carbide structure is comprised of 12-atom (B12 or B11C) icosahedra located at the vertices of a rhombohedral unit cell and a 3-atom linear chain connecting icosahedra along the longest body diagonal [1]. Although boron carbide has several polymorphs, the energetically favorable B4C structure is generally considered to have a C-B-C chain and B11C icosahedra with the C atom at a polar site [1–3]. In spite of the above-mentioned advantages, there are two challenges associated with using boron carbide for ballistic applications. First, it is very brittle and lacks local plastic deformation mechanisms like those observed in metals and even silicon carbide. Second, during high-speed impacts boron carbide undergoes localized solid-state amorphization with a significant increase in fragmentation [4]. High strain rate/high stress amorphization in boron carbide has commonly been observed as nano-scale bands with large aspect ratios that are separated by crystalline material [4–13]. This phenomenon has been identified under various testing conditions such as indentation [5,6], pressure shear plate impact [7,8], diamond anvil cell [9], mechanical scratching [5,10], laser shock [11], radiation [12], and electric fields [13]. Early density functional theory (DFT)-based simulations focused on hydrostatic pressure induced transformation to alternating layers of graphite and all-boron icosahedra [3], chain bending [3,14–17], and ejection of the B atom from the middle of the -C-B-C- chain [18,19]. Based on Yan et al.’s [20] diamond anvil experiments that highlighted the critical role of non2

hydrostatic stresses over hydrostatic pressure, An and Goddard performed quantum mechanics (QM) simulations to study incremental shear deformations for a variety of crystallographic slip systems [21]. Global shear deformation on the

̅

slip system was shown to

result in twinning, the formation of amorphous bands and eventual cracking within the amorphous band. Upon closer inspection, An and Goddard suggested that amorphization is initiated by strong interactions between B chain atoms and neighboring icosahedra C atoms and that these interactions distort and rupture icosahedra and trigger amorphization [21,22]. The theoretic debate and predictions continue, as an alternative explanation for amorphization was recently been put forth by Awasthi and Subhash, who interpreted their QM and molecular dynamics (MD) simulations to suggest that defects and inhomogeneities present in the material act as stress-concentrators, causing a spike in local pressure and temperature, which leads to highly localized melting followed by rapid quenching and amorphization [23]. Altering the chemistry of both the 12-atom icosahedra and the 3-atom linear chain has been suggested as a way to avoid amorphization [14,15,24,25]. QM simulations suggest that substituting carbon atoms with boron distorts the rhombohedral unit cell, which may improve toughness by favoring local plasticity [14,24,25]. Moreover, replacing carbon with boron in the icosahedra transforms it from B11C to B12 thus modifying twin formation [26] and chainicosahedra interactions in a way that is predicted to result in higher shear strength and decreased amorphization [14,15]. Stress-induced amorphization may also be suppressed by replacing the C-B-C- chain with silicon atoms [27,28], but the current study is focused on boron-rich boron carbide. Experimental validation is needed to confirm the simulation results. Experimental measures of the effect of boron enrichment on mechanical properties, such as hardness and fracture

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toughness, have been mixed. Niihara et al. [29] fabricated chemically vapor-deposited boron carbide with various B/C ratios and found decreases in both hardness and fracture toughness with increasing boron content. On the other hand, Champagne et al. [30] and Larsson et al. [31] reported that boron additions did not have significant effects on hardness or fracture toughness of hot-pressed and hot-isostatic pressed boron carbide. Recently, Xie et al. [32] and Cheng et al. [33] investigated hot-pressed B-rich boron carbides and found reduction in nanoindentation hardness by boron substitution, however, no apparent effect on the fracture toughness was observed. The current study was undertaken to test the hypothesis that boron-enrichment alleviates amorphization. For this, two polycrystalline samples with different B/C ratios (stoichiometric B4C and B-rich B6.3C) were fabricated via spark plasma sintering. Nanoindentation was used to measure hardness and as a means of triggering stress-induced amorphization in both types of boron carbide. Raman spectroscopy investigations were performed before and after indentation experiments and peak intensities were used to determine and contrast the relative amounts of amorphization in the samples. The Raman results were then compared with quantitative transmission electron microscope (TEM) examination of focused ion beam (FIB) prepared foils that were extracted from the region beneath the indents. Supplementary results of the observed relation between amorphization and cracking from similar TEM observations of a single-crystal regular B4C sample are also documented. 2. Materials and methods 2.1.Materials and processing In this study, three different samples, i.e. a boron carbide single-crystal and polycrystalline B4C and B-rich samples, were characterized. The single-crystal was fabricated using the float zone

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method in the PARADIM Bulk Crystal Facility [34]. The growth direction of the crystal is close to the <111> body diagonal of the rhombohedral unit cell. This single crystal sample is a part of an ongoing study and detailed characterization of this sample will be reported separately [34]. Fully dense stoichiometric and B-rich polycrystalline samples were consolidated via spark plasma sintering. To ensure minimal impurities, powders were washed prior to sintering via hydrochloric acid, then decanted with deionized water, and finally rinsed and dried with ethanol. A LECO TC600 analyzer revealed a very small amount (< 0.2 at. %) of impurities present in the initial boron carbide powder batch. To obtain the target boron-contents and to eliminate excess free carbon, the stoichiometric boron carbide powders were blended and sintered with amorphous boron powder. The sintering was carried out under 50 MPa pressure, pulsed DC current and with a holding time of 5 minutes at temperatures up to 1900°C. Cross-sectional slices were extracted from the center of the sintered disks, using a diamond wire saw, embedded in a conductive epoxy, and mechanically polished to a mirror-like surface finish using diamond lapping films (with the final grit size of 1 μm). 2.2.Chemical and phase analysis To determine precise stoichiometries of the polycrystalline samples, two separate standard chemical analysis methods were employed. Boron content was determined via boron mannitol titrations following ASTM Standard C791-04 [35], while a LECO C/S 230 chemical analyzer was used to determine carbon content by oxidation of the sample, followed by infrared detection. X-ray diffraction (XRD) was employed to investigate the phase compositions of the prepared samples. The XRD scans were collected using a Panalytical X'Pert diffractometer with a Cu-Xray source at 45 kV and 40 mA over a continuous scan range of 15°–90° 2; at a virtual step size of 0.0131° and counting time of 200 s. 5

2.3.Nanoindentation To interrogate hardness and to activate stress-induced amorphization, nanoindentation experiments were carried out on polished boron carbide samples. A Micro Materials NanoTest Vantage nanoindentation system equipped with a Berkovich indenter was used at a constant load of 500 mN, loading/unloading rate of 2.5 mN/s and with a hold time of 2 s at maximum load. 2.4.Raman investigations Raman spectroscopy was performed to obtain additional insights into the compositional variation within a phase of the prepared samples. Raman data was collected using a Renishaw inVia Reflex system. Amorphous zone distribution within the nanoindent impressions were evaluated using a 633 nm laser with spatial resolution of 0.5 m (100x objective). The Raman spectra were deconvoluted using WIRE 4.2 software. 2.5.Transmission electron microscope (TEM) investigations Cross-sectional TEM specimens from beneath nanoindent impressions were prepared by a liftout technique using a Thermo Fischer Helios G4 Dual Beam FIB system. First, a protective carbon layer was deposited on selected nanoindents using a 20kV electron beam. The nanoindented samples were then tilted to an angle of 52°, and an additional protective Pt layer was deposited over previously sputtered carbon layer using low Ga+ beam current (0.79 nA) with the accelerated voltage of 30kV. A large trench was excavated on either side of this Pt layer using a high Ga+ beam current (21nA). After removing the remaining residual portion (including typical J-cut at -10º angle), the specimen was tilted back to its original position and a tungsten microprobe was attached to the top edge of the Pt layer by Pt deposition. The TEM lamella was

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then in-situ lifted out and attached onto a copper TEM half-grid using Pt deposition, followed by removal of the microprobe tip from the lamella via Ga+ beam cutting. Finally, the lamella was thinned down to a thickness of 50 nm with multiple passes of a low-energy ion beam. The prepared TEM samples were carefully examined in a Tecnai TF30 microscope at 300KV. Conventional TEM bright-field (BF) imaging was carried out to visualize the undeformed and deformed microstructure beneath the nanoindent. High-resolution TEM (HRTEM) investigations were performed to determine the existence of nano-scale mechanisms such as localized amorphization. In order to investigate local orientation changes and related phenomena, selected area electron diffraction (SAED) patterns and automated crystal orientation mapping (ACOM) were utilized. For ACOM data acquisition, the TEM was equipped with a NanoMegas ASTAR system. Using a spot size of 9, gun lens of 3, extraction voltage of 4.5 kV and a 30m C2 aperture, a probe with ~ 1 nm diameter and a convergence semi-angle of 1.4 mrad was generated. A camera length of 135 mm was used to acquire diffraction patterns with a step size of 5 nm. These patterns were primarily indexed/analyzed using commercial ASTAR software to obtain digital crystallographic datasets. The datasets were then further analyzed using standard electron backscattered diffraction (EBSD) software (OIM version 8 developed by EDAX, Inc). 3. Results and discussion 3.1.Stoichiometries and phase composition of the samples To obtain precise stoichiometries of the prepared samples, both carbon and boron content was analyzed using standard chemical methods. The carbon content of the polycrystalline stoichiometric and B-rich boron carbide samples was measured to be 19.9 0.1 at. %, and 13.7  0.3 at. %, respectively. The corresponding boron content was measured to be 79.8  0.1 at. %

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and 86.2  0.1 at. % for stoichiometric and B-rich boron carbide samples, respectively. Using these values, the final stoichiometries were estimated to be B4C and B6.3C based on their postprocessing constituents, rather than their pre-sintered estimated stoichiometries. XRD investigations were performed to determine phase composition of the prepared samples. Fig. 1a presents XRD patterns acquired from the polycrystalline stoichiometric B4C and B-rich B6.3C samples. All diffraction peaks from both samples correspond to the primary boron carbide phase with the rhombohedral Bravais lattice of R3m space group (Fig. 1a). No free carbon was detected and the additional boron atoms are incorporated into the boron carbide crystal structure, creating a B-enriched solid solution. All B-rich peaks were found to be shifted towards lower angles in comparison to that of the stoichiometric sample, indicating an increase in unit cell volume (Fig. 1b illustrates clear examples of (104) and (021) peak shifts). Such lattice expansion has also been predicted by atomistic simulations and observed experimentally [14,24,32,33].

Fig. 1: (a) XRD spectra of the polycrystalline stoichiometric B4C and B-rich B6.3C samples. (b) An expanded view of the (104) and (021) peaks, showing a clear peak shift towards lower angles (lager unit cells) for the B6.3C sample.

3.2.Hardness measurements 8

Berkovich nanoindentation was used to interrogate hardness and to activate stress-induced amorphization in all samples. The measured nanoindentation hardness of the stoichiometric B4C samples range from 36 to 43 GPa (with an average of 40 2 GPa), while that of the B-rich B6.3C sample from 34 to 38 GPa (with an average of 37 1 GPa). Note that the average hardness and the highest hardness value of the stoichiometric sample are both statistically larger than that of the B-rich sample, indicating that the boron addition causes a reduction in boron carbide hardness. This finding is in agreement with those recently reported by Xie et al. [32] and Cheng et al. [33]. This decrease in hardness may be related to the lattice expansion induced by boron substitution. Furthermore, it is know that the boron carbide receives its extraordinary hardness from its strong covalent bonds [1] and the B-C bonds have been reported to be stronger than the B-B bonds [36]. Substituting boron for carbon reduces the overall proportion of B-C bonds, which is consistent with the measured lower intrinsic hardness [32]; however, the effects of quasi-plasticity cannot be discounted and may have a strong effect on the hardness of these samples. 3.3.Raman investigations Raman spectroscopy provided additional insight into the effect of excess boron on the crystalline nature of boron carbide. Fig. 2 presents typical Raman spectra measured from the pristine surface of the two different types of polycrystalline boron carbide samples. The resultant spectra are characterized by a series of peaks extending from 200 to 1200 cm−1 assigned to the crystalline boron carbide [37,38]. Typically, the peaks in the high frequency regime (530 to 1200 cm−1) are associated with the vibration modes originating from the icosahedra, while the narrow peak at 485 cm−1 is attributed to stretching of the three-atom linear chain [28,37,39,40]. The low frequency Raman peaks, such as the 270 and 320 cm−1 doublets, are generally attributed to

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disorder-activated acoustic phonons [1,28,40]. From Fig. 2, it can be seen that simply changing the stoichiometry of boron carbide has a few significant effects on its Raman spectrum. The peaks at 1088, 1000, 830 728, and 530 cm−1 are present in both samples but broaden and shift towards lower frequencies in the boron-rich sample. Consistent with the literature, this observation can be linked to the volume expansion due to B substitution in the icosahedra [28,32,33,41]. Furthermore, two additional new peaks appear in the B-rich spectrum. A broad peak is evident at 375 cm−1 and a shoulder peak at 1170 cm−1. Similar observations have also been made previously [32,33,37,41]. Xie et al. [32] proposed that the origin of 375 cm−1 peak is related to structural changes in B-rich boron carbide as boron atoms substitute for carbon in the chains. The 1170 cm−1 peak is close to the vibrating mode of the icosahedra and may signify enhanced B-B bonding in the icosahedra as a result of the substitution of carbon by boron [33]. Therefore, for present B-rich composition (B6.3C) boron appears to have substituted carbon in both icosahedra and chain units.

Fig. 2: Typical Raman spectra of the as-prepared stoichiometric and B-rich samples acquired at an excitation energy of 1.95 eV (633 nm). In contrast to the stoichiometric spectrum, boron enrichment results in shifting some characteristic Raman peaks towards lower frequency and appearance of two additional new peaks at 375 cm−1 and 1170 cm−1.

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The creation of boron-rich boron carbide has been suggested as a way to mitigate stress-induced amorphization [14,15]. The idea is to substitute C atoms in the icosahedra with B atoms [14,24,25,42,43], which eliminates the strong C-B interaction that is predicted to initiate amorphization. To test this claim, Raman spectra and maps of the area around multiple indents in both stoichiometric and B-rich samples were collected. Each indent and Raman map spans over multiple grains. Fig. 3 presents representative comparisons of the Raman spectra acquired from pristine and indented regions of the two types of samples. In contrast to the pristine region spectra, three additional peaks are visible in the indented region spectra (Fig. 3a and b). A strong peak is apparent at 1330 cm−1, followed by a shoulder peak at 1480 cm−1 and a broad peak at 1810 cm−1. The precise origin for the formation of these peaks is still under debate, but the prevailing opinion is that they are associated with the formation of amorphous boron carbide as a result of indentation [5,6,23,28,44]. In addition to these dominant characteristics, a slight rightshift of the crystalline peaks above 400 cm−1 is also observed in the indented spectra. These shifts have been observed before [16, 37–39]. Awasthi and Subhash recently conducted first principles and MD simulations of the indentation and amorphization processes, and they have associated the Raman peaks to residual compressive stresses formed as a result of the amorphization process [23,44]. Comparing the relative intensities of the crystalline and amorphous peaks in Fig. 3, it is evident that the intensity of the amorphous peak at 1330 cm-1 is 3.5 times greater than the crystalline peak at 1088 cm-1 in the indented stoichiometric sample and only 1.5 times greater in the indented B-rich sample. All of the amorphous peaks were found to be smaller in the B-rich sample than in the stoichiometric sample, and the overall intensity of the Raman maps that were collected was significantly higher for the stoichiometric samples than it was for the B-rich samples. Taken as a whole, these Raman measurements provide indirect

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evidence that introducing excess boron reduces the degree of amorphization in boron carbide. In order to further validate these Raman results, a direct observation using TEM was undertaken, the results of which can be described as follows.

Fig. 3: Representative Raman spectra acquired at an excitation energy of 1.95 eV (633 nm) from the pristine and residual indented regions of polycrystalline (a) stoichiometric B4C and (b) B-rich B6.3C. In contrast to the pristine region spectra, three additional broad peaks are visible at 1330, 1480, and 1810 cm−1 in the spectra obtained from the indented regions, signifying formation of amorphous phase. A slight right-shift of the crystalline peaks above 400 cm−1 in the indented spectra could be associated with residual stresses. The prominence of the amorphous peaks is greater in the stoichiometric sample, suggesting that there is less amorphization in the B-rich sample.

3.4.TEM investigations to understand deformation characteristics In order to characterize the deformation mechanisms associated with nanoindentation, FIB liftout samples were prepared to enable TEM investigations. Similar to the location of the Raman spectra, all TEM samples were also lifted-out close to the center of the nanoindents. Fig. 4 presents a representative BF-micrograph of the region under an indent in a B-rich sample. The position of the indent is highlighted on the top of the micrograph. The cracks visible at the bottom of the micrograph grew during the thinning process, releasing the tensile stress that formed as a result of the plastic zone below the indent. The undeformed grains are a few micrometers in diameter and were observed to contain a high density of planar defects, see for

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example Fig. 5). Xie et al. [32] also provided a detailed high-resolution characterization of these planar defects in both stoichiometric and B-rich boron carbide. The density of planar defects was observed to be significantly greater in the B-rich samples, and they were shown to be growth twins and stacking faults that appear to be associated with either a decrease of stacking fault energy upon boron-enrichment [32] or local chemical accommodation in the form of Wadsley defects [46,47].

Fig. 4: BF-micrograph presents a cross-sectional view of the B-rich B4C lift-out sample taken close to the center of 500 mN nanoindent using standard FIB techniques. The marked microcracks grew during sample preparation owning to the presence of the tensile residual stresses. The undeformed grains are predominantly in the range of few micrometers and contain a high density of planar defects (also see Fig. 5). The indent induced a highly deformed region.

Fig. 5: (a) BF-micrographs (like the one shown here) and (b) HRTEM images (like this) both attest to the presence of a high-density of planar defects (twins and stacking faults) in the B-rich boron carbide samples. The FFT pattern in (b) shows well-defined twin and matrix reflections.

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The indents created plastically deformed regions that spread across a handful of grains, and magnified cross-sectional views of the deformed regions were also obtained, see for example Fig. 6. The strain contrast visible in this image is a clear indication that the grains have undergone considerable permanent deformation. Moreover, transgranular nanoscale bands were observed across the whole deformed region, see the sub-micron features marked with arrows in Fig. 6. These bands are nominally aligned with the maximum shear directions for the indent, although deviations of ~ 20 degrees have been observed. It appears that both the macroscopic stress state and the crystallographic orientation of individual grains both influence the orientation of these nanoscale shear bands, as previously reported by Zhao et al. [11]. Based on their contrast, thickness and position from the indent’s impression, these shear bands are found to be of two types. Thinner bands (less than 5 nm wide) with alternating dark-bright contrast were observed and found to be located relatively far away from the indent (Fig. 7a). Closer to the indent where the stresses were higher, thicker bright contrast bands (with widths of 5-10 nm) were observed (Fig. 7b). Interestingly, there exists additional strain contrast around the thinner bands that appears to be absent around the thicker bright bands. This loss of strain contrast signals a relaxation of residual elastic strain.

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Fig. 6: BF-micrograph presenting a magnified cross-sectional view of the area beneath a nanoindention in B-rich boron carbide. The extended strain contrast indicates that the material is plastically deformed. The arrows mark transgranular nanoscale bands that were formed during loading and are loosely oriented along the maximum shear direction, although deviations of ~ 20 degrees are observed.

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Fig. 7: BF-micrographs shows two types of nanoscale transgranular shear bands formed in a close proximity to a nanoindentation apex in B-rich boron carbide. (a) Thinner bands (< 5 nm wide) with alternating dark-bright contrast and (b) thicker bright contrast bands ( 5 to 10 nm in width).

In order to characterize these nanoscale shear bands further, HRTEM investigations were carried out. The HRTEM images in Fig. 8 show crystalline lattices on either side of the bands, and a loss of lattice fringes inside of both bands, indicating localized amorphization. Fast Fourier transformations (FFT) further confirmed the atomic disorder of the amorphous bands, as shown in the insets in Fig. 8. The FFT patterns from area “A” inside the nanoscale shear bands show “a diffuse halo” indicative of the amorphous phase. On the other hand, the spot patterns obtained from regions marked “C” are indicative of the rhombohedral crystal structure. It has been suggested that the highly anisotropic nature of boron carbide will lead to amorphization along preferred crystallographic directions [4]. The FFT patterns suggest that these nano-amorphous

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bands are roughly, but not exclusively, aligned along low indices planes e.g. (011) in Fig. 8a. This indicates that both the stress state and the crystallographic orientation have a role to play in determining the orientation of the amorphous bands. The strain contrast highlighted in Fig. 6 and 7 is also evident in Fig. 8a and the observation that the contrast fringes are nearly perpendicular to the thin bands in this image may be taken as evidence of a shear displacement. The contrast fringes were not observed around the thicker bright bands (e.g. Fig. 8b). These observations suggest that formation of the thin amorphous bands creates a residual strain field and that this strain is then relaxed as the bands get wider and trigger intragranular cracking.

Fig. 8: HRTEM micrographs taken from a B-rich sample showing inhomogeneous loss of crystalline order inside both types of shear bands, confirming them to be amorphous. The FFT patterns of the area inside the shear bands present the “diffuse halo” typical of amorphous regions, while the spot patterns from the surrounding matrix index to the rhombohedral crystal structure. The bands are loosely aligned along low indices planes, e.g. (011) in (a). Contrast fringes associated with the thin band in (a) is absent around the thicker band in (b).

Indentation investigations were also carried out on stoichiometric boron carbide samples in both polycrystalline and single-crystalline form (Fig. 9 and Fig. 10). The grains size of the stoichiometric polycrystalline boron carbide was similar to the B-rich material. The stoichiometric polycrystalline sample exhibited similar transgranular nanoscale amorphous band

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formation beneath the indents (Fig. 9), although the number density of amorphous bands was clearly higher for the stoichiometric samples than the B-rich samples (discussed below). The amorphous shear bands were also observed in the single-crystalline sample (Fig. 10), although they are significantly longer, possibly due to the absence of the grain boundaries. This suggests reduction in grain size as an another possible approach to mitigate amorphization in boron carbide, a finding that appears to be consistent with the recently reported Raman results [27,48]. The amorphous shear bands observed in this study are in close agreement with past observations made on indented [5,6] and ballistically impacted [4] boron carbide, where mode II and mode III shear instabilities are considered responsible for the formation of such narrow amorphous bands. Past studies have hypothesized that amorphization, cracking and comminution are all interrelated, and it is reassuring to note that the current study provides a clear connection between amorphization and cracking. Fig. 10c is unique in that both crack tips were captured within the image. A magnified view of this crack and its relation to the amorphous band is given in Fig. 10e. The presence of this crack with both crack tips provides clear evidence that cracks can and do nucleate and propagate along amorphous shear bands. Moreover, the observed change in the propagation direction of a crack, when it intersected an amorphous shear band, in Fig. 10d provides direct visual confirmation that resistance to crack propagation is lower in amorphous boron carbide than it is for crystalline boron carbide. These observations provide strong experimental evidence to conclude that amorphization occurs first and then provides a trigger and crack path for intragranular cracking. As such, these findings would be consistent with QM simulations that predict cavity formation and cracking within nanoscale amorphous bands [22] and mesoscale models that assume cracking after a critical

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shear threshold [49]. However, in contrast to the QM simulations [22], no deformation twinning has been observed in the samples examined for this study.

Fig. 9: (a) BF-micrograph presenting a cross-sectional view of a stoichiometric polycrystalline lift-out taken close to the center of a 500 mN nanoindent. (b) A magnified view of the plastic zone under the indent shows numerous nanoscale shear bands, with associated residual strain contrast. (c) The HRTEM micrograph confirms the nanoscale band to be amorphous.

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Fig. 10: BF-micrograph (a) presents a cross-sectional view of the region under a center of 500 mN indent in a single-crystal of boron carbide (indented along its <111> growth direction). The magnified micrographs show: (b) microns long amorphous shear bands (highlighted with white arrows, also marked in (c) and (d)) with residual strain contrast and bend contours, (c) evidence of a crack that nucleated along an amorphous shear band (magnified rotated view in (e)) and (d) evidence of a crack that intersected an amorphous shear band, kinked, and propagated along the band.

In addition to the formation of nanoscale amorphous shear bands, the formation of randomly oriented nanocrystalline grains was also observed in the deformation zone beneath the indents in all of the samples investigated. Fig. 11 presents an example of a B-rich B6.3C sample showing clearly defined deformed and undeformed regions beneath a nanoindent. A SAED pattern obtained from the deformed region in Fig. 11a is shown in Fig. 11b; the organization of spots

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into rings is an indication of randomly oriented nanocrystalline grains. This finding was further supported by the formation of orientation maps of the deformation zone with the TEM-based ACOM technique. For example, local orientation changes and randomly oriented nanocrystalline grains that are much smaller than the original grains, can be seen in Fig. 12. Similar observations have also been reported by Ge et al. [5], and possible mechanisms that could be used to explain these observations are: (i) boron carbide undergoing successive solid-state phase transformations under high contact pressures with the resultant high-pressure phase nucleating randomly oriented nanocrystalline grains as it relaxes back to the rhombohedral structure after load release, (ii) the extensive local deformation of the original boron carbide grains that results in highly localized shear, rotation and the formation of high-angle grain boundaries, or (iii) localized melting and nucleation of nanograins upon quenching [11,23].

Fig. 11: (a) Evidence of clearly defined deformed and undeformed zones beneath an indent in a B-rich B6.3C sample. (b) A SAED pattern taken from the deformed region in (a) exhibits a ring pattern indicative of the formation of randomly oriented nanocrystalline grains.

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Fig. 12: Orientation map acquired beneath an indent in stoichiometric B4C with the ACOM technique. A color-orientation key is provided. The deformed region exhibits randomly oriented nanocrystalline grains that are much smaller than the original grain size. The indent appears to have hit at a grain boundary between two parent grains (blue and green). The blue grain shows the clear formation of nanocrystalline grains but both grains show evidence of local orientation changes.

3.5.TEM agreement with the Raman results Raman investigations revealed that the intensity of the peaks associated with amorphization in polycrystalline B-rich B6.3C was lower than it was for stoichiometric B4C. In order to validate these Raman results, i.e. to directly visualize amorphization and quantify amorphous shear band densities, images were acquired at the same magnification and proximity to intents in both B4C and B6.3C. One example of the comparison of the microstructures beneath indents for two polycrystalline materials is shown in Fig. 13. The BF-micrographs of similar areas beneath nanoindents in stoichiometric B4C and B-rich B6.3C show comparatively lower amorphous shear band density in the latter sample. To estimate this reduction quantitatively, a line intercept method was employed and the density of amorphous bands defined as: (

∑ ∑

∑ ∑

)

In this method, a grid consisting of 9 vertical and 9 horizontal lines is superimposed on the BFTEM micrographs. The number of intersections of bands with the vertical (

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) and horizontal

(

) grid lines are counted and ∑

and ∑

are the total lengths of the vertical and horizontal

test lines. The TEM lift-out sample thickness was measured to be ~ 50 nm for both samples, and the band density for the stoichiometric and B-rich samples were found to be 5.5  0.4 1014 m-2 and 3.8  0.4 1014 m-2, respectively. This analysis was repeated for two samples from each of the alloys and the difference between the stoichiometric and B-rich samples was observed to be greater than the variation within either type of sample. These results further suggest that the observed difference in the amount of amorphization is not directly attributable to the random orientation of the grains beneath the indent. Taken as a whole, the TEM observations indicate that amorphization is reduced by 30% with the boron-enrichment, which is fully consistent with our Raman results as well as previously reported simulations [14,15]. This observed reduction in amorphization may be expected to have a significant influence on ballistic performance because micro-cracks are shown to nucleate and propagate along the nanoscale amorphous shear bands. And, the results of this study suggest that changes in stoichiometry offer one potential path to improve ballistic properties of boron carbide.

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Fig. 13: Representative micrographs from similar areas beneath nanoindents show a comparatively lower number density of nanoscale amorphous shear bands in polycrystalline B-rich B6.3C sample than in stochiometric polycrystalline B4C sample. Using line intercept method this reduction was found to be  30 %.

4. Summary and Conclusions Boron-enrichment has been suggested as a possible way to mitigate stress-induced amorphization in boron carbide, and the experimental study presented here was undertaken to test this hypothesis. Stoichiometric B4C and B-rich B6.3C samples were produced and examined with nanoindentation, X-ray diffraction, Raman spectroscopy and electron microscopy. Nanoindentation hardness measurements indicate that substituting carbon atoms with boron softens boron carbide, possibly due to the lattice expansion and replacement of the stronger B-C bonds with weaker B-B bonds. The XRD and Raman spectroscopy results revealed an increase in unit cell volume with B substitution into the icosahedra as well as chain units. Comparison of the Raman spectra from pristine regions with no indents and from beneath nanohardness indents

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provided indirect evidence of amorphization during indentation and the intensity of the amorphous peaks was significantly lower for the B-rich samples. TEM and HRTEM observations of foils lifted out from the region directly beneath the indents provided direct evidence of: nanoscale amorphous shear bands in a strained crystalline lattice, localized orientaion changes accompanied by the formation of randomly oriented nanocrystalline grains, and intragranular microcracking. The TEM results point to a 30% reduction in amorphization with boron-enrichment, which is consistent with the Raman measurements. Moreover, observations of the interplay between cracks and amorphous bands provide clear evidence to suggest that amorphization occurs first and provides a trigger and easy crack path for intragranular cracking. Taken all together, these results provide a mechanistic road map for modeling the dynamic amorphization and comminution of boron carbide and suggest that variations in stoichiometry offer a potential path to improving its ballistic performance. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

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Acknowledgements The authors would like to thank Prof. Michael G. Spencer from the State Morgan University, USA for providing single-crystalline boron carbide sample. This research was sponsored by the Army Research Laboratory and was accomplished under Cooperative Agreement Number W911NF-12-2-0022. The views and conclusions contained in this document are those of the authors and should not be interpreted as representing the official policies, either expressed or implied, of the Army Research Laboratory or the U.S. Government. The U.S. Government is authorized to reproduce and distribute reprints for Government purposes notwithstanding any copyright notation herein.

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Graphical abstract

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