Dynamic Indentation Response of Fine-Grained Boron Carbide

Dipankar Ghosh

Department ofMaterials Science and Engineering, Michigan Technological University, Houghton, Michigan 49931-1295

Ghatu Subhash*,w

Department of Mechanical and Aerospace Engineering, University of Florida, Gainesville, Florida 32611

Tirumalai S. Sudarshan and Ramachandran Radhakrishnan*

Materials Modifications Inc., Fairfax, Virginia 22031

Xin-Lin Gao

Department of Mechanical Engineering, Texas A&M University, College Station, Texas 77843-3123

Boron carbide disks with three different grain sizes were con-solidated from submicrometer-sized boron carbide powder usingthe plasma pressure compaction technique. Static and dynamicindentations were performed to determine their loading-ratedependence on mechanical properties. Dynamic indentationsresulted in a decrease in hardness and fracture toughness, andinduced more severe damage compared with static indentations.Using Raman spectroscopy, the mechanism responsible for lossof strength under dynamic loads was identified as the solid-statestructural phase transformation in the dynamically loaded re-gions. The influence of processing conditions and the resultingmicrostructure on the observed rate dependency of mechanicalproperties are discussed.

I. Introduction

BORON carbide (B4C) ceramic is an excellent candidate mate-rial for structural applications1–7 at room and high temper-

atures because of its high melting point (24501C), high elasticmodulus (450 GPa), high hardness (HV 25–35 GPa, next only todiamond and cubic boron nitride), high flexural strength (350–500 MPa), low density (2.52 g/cm3), and excellent wear resis-tance. It is used as a grinding medium for hard materials, as alightweight ceramic armor, as wear-resistant sandblasting nozzlematerial, and as a neutron absorber in nuclear reactors.1,8 How-ever, the processing of boron carbide is challenging due to thedifficulty associated with sintering of the starting boron carbidepowder. Traditionally, boron carbide has been consolidated us-ing (i) hot pressing with and without sintering additives,1,9,10 (ii)hot isostatic pressing (HIP),1,11 (iii) pressureless sintering withsintering additives,1,2,4–6,12,13 (iv) pressureless sintering in agaseous atmosphere of hydrogen and helium,14 and (v) micro-wave sintering.15 Among the above processing techniques, hot-pressing and pressureless sintering are the most commonly usedmethods to produce boron carbide ceramics with 95%–99%of the theoretical density and with grain sizes in the rangebetween 1.5 and 60 mm. However, in these processing methods,

the sintering temperatures were relatively high (420001C) and thesintering times were on the order of hours. In addition, sinteringaids were found to reduce fracture strength moderately,13 and theresulting sintered ceramics were not suitable for nuclearapplications1 where high-purity boron carbide is required forneutron absorption. A dense boron carbide ceramic can be pro-duced at lower temperatures using HIP but this method is notsuitable for bulk processing.1,11 Boron carbide powder heat treat-ed in a gaseous mixture of hydrogen and helium and then sinteredin the presence of pure helium also requires a sintering temper-ature above 22001C.14 Microwave sintering has also been used toconsolidate boron carbide (95% density) in a short duration oftime (12 min) but it also requires high temperatures around20001C.15

Recently, Klotz et al.16 applied a novel non-conventionalmethod known as plasma pressure compaction (P2Cs) for bo-ron carbide consolidation. Sintering was performed at 16501Cwithin a short consolidation time of 5 min, but the resultingdensity was only around 91% of the theoretical value. Severalsintering additives such as graphite, alumina, and titanium dibo-ride were used to achieve up to 97% of theoretical density. In thecurrent work, boron carbide was consolidated using the aboveP2Cs method with an intent to produce theoretically densecompacts with various grain sizes but without the addition ofsintering aids.

Investigations on the evaluation of the mechanical propertiesof boron carbide have been mainly focused on determining elas-tic modulus, flexural strength, fracture toughness,1–7 and nano-indentation response.17,18 Owing to its promise as an effectivearmor material against low-velocity ballistic threats,1,8 highstrain rate experiments were conducted on boron carbide. How-ever, these experiments are expensive, time consuming, and re-quire large-size specimens.8,19–21 Therefore, in the current work,dynamic indentation experiments at high loading rates wereconducted to capture the strain rate sensitivity of indentationresponse in boron carbide ceramics of different grain sizes. Anovel dynamic indentation tester22–24 was utilized to determinethe dynamic hardness and then to compare these results with thestatic hardness measurements. To study the damage accumulat-ed beneath the indentation, subsurface studies using the bonded-interface technique was conducted.25 Raman spectroscopy is auseful nondestructive tool for analyzing structural changes aswell as phase transformations in crystalline materials.26 Thistechnique has been used to investigate the contact damage in-duced by nanoindentation in single crystal and polycrystallineboron carbide ceramics.17,18 Therefore, in the current work,Raman spectroscopy was used to analyze the damage beneaththe static and dynamic indentations.

G. Pharr—contributing editor

This work was funded by a grant from the US NSF (grant# CMS-0324461) withDr. Ken Chong as the program manager.

*Member, American Ceramic Society.wAuthor to whom correspondence should be addressed. e-mail: subhash@ufl.edu

Manuscript No. 22412. Received October 28, 2006; approved February 20, 2007.

Journal

J. Am. Ceram. Soc., 90 [6] 1850–1857 (2007)

DOI: 10.1111/j.1551-2916.2007.01652.x

r 2007 The American Ceramic Society

1850

II. Experimental Procedure

Commercially available boron carbide powder (Grade HS, H.C.Starck, Newton, MA), with particles of sizes around 800 nm,was used as the starting material. The powder was consolidatedusing the P2Cs method16,27–29 under high vacuum (200–300mTorr) without the use of any sintering aids. In the P2Cs meth-od, the powder is subjected to a constant low-voltage (5 V) di-rect current, which causes ‘‘Joule’’ heating along theinterparticle contact areas. As a result, the temperature of thepowder compact increases with increasing current. Simulta-neously, an external pressure (88 MPa) is applied to aid pow-der consolidation. In the current work, a maximum currentdensity of around 4400 A/cm2 was used. In each batch, 33 g ofpowder was used to produce a sintered compact in the form of acylindrical disk of 51 mm diameter and 6.4 mm thickness. Thesedisks were sintered either at a temperature of 16501 or 17501C,and the consolidation time was varied between 2 and 30 min.Table I shows the processing conditions for all the disks. SampleIDs used in Table I reflect the processing conditions, i.e., the firstset of numbers indicate the sintering temperature (1C) and thenext set of numbers indicate the processing time (min) at thatsintering temperature. More details of the P2Cs method areavailable elsewhere.16,27–29

The densities of the sintered disks were measured using theArchimedes method.Microstructural analysis was performed onboth fractured and polished surfaces. Polished specimens wereprepared by using standard metallographic principles. The sur-faces were ground successively with 120, 240, 320, 400, and 600grit silicon carbide papers for 10 min each and then polished forat least 5 min each using 6 and 1 mm diamond paste. Then thepolished surfaces were etched electrolytically using 1% KOHsolution at a current density of 0.03 A/cm2 for 30 s. From theoptical micrographs of the polished and the etched specimens,the average grain size was determined (by the line interceptmethod) as per ASTM standard E 112-96.

Two types of specimens were used: (i) rectangular specimensfor determination of indentation hardness and for estimation ofcrack lengths on the surface for fracture toughness measure-ments (KC) and (ii) bonded split specimens25 for determinationof subsurface damage beneath the indentation and for Ramanspectroscopy analysis of the damaged regions. Rectangular spec-imens of dimensions 4 mm! 3 mm! 6 mm were cut from thesintered disks and 4 mm! 3 mm surfaces of each specimen werepolished (using the same polishing technique as discussed be-fore) for static and dynamic Vickers indentation experiments.

Around 15 static indentations were performed at each load of2.94 N (300 g), 4.9 N (500 g), and 9.8 N (1000 g) for 15 s on eachgrain size specimen. The dynamic indentation experiments were

performed at comparable load levels using a dynamic indenta-tion hardness tester, shown schematically in Fig. 1.22–24 In thistechnique, elastic stress wave propagation in a slender rod isutilized to deliver the desired load within a 150-ms duration. Thistechnique is parallel to the static indentation technique andtherefore, a direct comparison between static and dynamic in-dentations can be made. The dynamic hardness tester consists ofa slender rod with a Vickers indenter mounted at one end and amomentum trap (MT) assembly at the other end. A high-fre-quency load cell Kistler mounted on a rigid base measures theload. The validity of the indentation load measurements in thecurrent work using the above load cell can be found else-where.22–24 The specimen is sandwiched between the diamondindenter and a load cell. A short striker bar is launched from agas gun toward theMT end of the incident bar, thus generating acompressive stress pulse, followed by a tensile pulse (due to MT)of known duration and amplitude in the incident bar. The MTassembly ensures that only a single compressive pulse reaches theindenter, thus causing the indentation, and then the tensile pulseretracts the indenter.22,30 Furthermore, all the successive wavereflections will be tensile while traveling toward the indenter end,thus causing the bar/indenter assembly to retract further awayfrom the specimen. Therefore, single dynamic indentation on thespecimen is ensured. The dynamic hardness is calculated basedon indentation diagonal size and load. Unlike the static inden-tation tests at fixed loads, the dynamic indentations were con-ducted at loads between 2.94 and 14.7 N by increasing thevelocity of the striker bar in the dynamic indentation tester.Around 30 tests per specimen type were conducted at load rangessimilar to those under static indentation. The indentation-in-duced strain rate was estimated to be around 1000 s"1.22

Cracks extending from the corners of the static and dynamicVickers indentations were used for fracture toughness (KC) mea-surements. Numerous empirical fracture toughness relations areavailable in the literature depending on the nature of the cracksystems.31 For half-penny cracks,32 the relationship between thehalf-crack length c and load P is expected to be of type c5AP2/3

whereas for Palmqvist cracks,33 the difference between the half-crack length c and the half-diagonal, a, i.e., l5 c " a, is expectedto be linear with respect to load, i.e., l5BP. In the current work,the above quantities were plotted and fitted to an equation oftype y5 bxn to determine the nature of the crack system. The cversus P plots yielded n values in the range of 0.62–0.69 andtherefore, matched the equation y5 bxn closely. On the otherhand, the plots of l versus P did not result in a linear relationshipbetween l and P as the n values were found to be in the range of0.68–0.92. Therefore, it was concluded that in the current work,static and dynamic indentations resulted in a half-penny cracksystem beneath the indentations. Accordingly, the Evans andCharles31 fracture toughness equation, KC 5 0.0824P/c1.5, basedon a half-penny crack system has been used. Recently, Lee andSpeyer7 used a similar fracture toughness equation based on ahalf-penny crack system to calculate KC for boron carbide with96%–97% of the theoretical density.

For subsurface studies, the rectangular specimens were cutinto two halves along the length, and these cut surfaces werepolished using the steps described previously. They were thenbonded with a high-strength adhesive and kept clamped for 3–4h. The top surface containing the bonded interface was polishedflat for static and dynamic indentations along the interface. Theindentation was performed with one of the diagonals of the

Flange Rigid mass

Sleeve

Digital Oscilloscope

Charge Amplifier

Momentum Trap

Striker Bar

Incident Bar

Indenter Load CellSpecimen

Fig. 1. Schematic of the experimental setup for dynamic indentationhardness measurements.

Table I. Processing Conditions in the P2Cs Method for Sintering of Boron Carbide Compacts

Sample ID Temperature (1C) Time (min) Pressure (MPa) Density (% theoretical) (g/cc) Grain size (mm)

1650_5 1650 5 88 2.3370.002 (92.46%) Not measured1650_30 1650 30 88 2.3570.003 (93.25%) Not measured1750_2 1750 2 88 2.4270.002 (96%) 1.61750_5 1750 5 88 2.5070.004 (99.2%) 2.01750_30 1750 30 88 2.5070.004 (99.2%) 2.7

June 2007 Dynamic Indentation Response of B4C 1851

indenter aligned parallel to the interface as shown in Fig. 2. Forthese subsurface damage studies, indentations were intentionallyperformed at slightly higher loads (up to 21 N) so as to cause asufficiently large damage zone beneath the indentation. Afterindentation, the bonded surfaces were separated and scanningelectron microscopy (SEM) was performed to observe the sub-surface damage. The subsurface-damaged zones were then an-alyzed, for any possible phase change, using a Ramanspectrometer equipped with an optical microscope for focusingan incident He–Ne laser beam (632.4 nm) of 1 mm spot size.

III. Results

(1) Density and Microstructure

Table I provides the density measurements on specimens cutfrom boron carbide compacts. Boron carbide sintered at 16501Cattained only 92%–93% of the theoretical density. On increas-ing the sintering temperature to 17501C, a significant improve-ment in density was observed. Compacts consolidated for 2 and5 min at 17501C attained densities around 96% and 99% of thetheoretical value, respectively. When the consolidation time wasincreased to 30 min at this temperature, no further increase indensity was noticed. The improvement in density (or decrease inporosity level) with an increase in temperature is clearly evidentfrom the SEM micrographs of the fractured surfaces shown inFig. 3. Owing to the low density (high porosity), samplessintered at 16501C were not considered for further study. The

microstructures of the polished and electrolytically etched spec-imens from the three boron carbide disks produced at 17501Care given in Fig. 4. Optical micrographs revealed nearly equi-axed fine-grained microstructures with average grain sizesaround 1.6, 2, and 2.7 mm for the samples sintered for 2, 5,and 30 min, respectively. Darker regions in the micrographsrepresent mostly grain pull-outs during polishing. From thedensity and average grain size measurements, it is clear thatthe processing parameters (temperature, pressure, and consoli-dation time) adopted in the current work were suitable for theproduction of dense and fine-grained boron carbide ceramics.

(2) Comparison of Static and Dynamic Indentation Hardness

A plot of static and dynamic hardnesses versus indentation loadfor the three grain sizes is presented in Fig. 5. The static hardness(HVs) values appear in groups because the indentations wereperformed at fixed loads available in the hardness tester. A widescatter in the static hardness values was observed for all thegrain sizes. These values are summarized in Table II. The aver-age static hardness values for 2 and 2.7 mm grain size specimensappeared in the same range (27.4572 and 27.4572.5 GPa) forthe loads considered in this investigation. On the other hand, the1.6 mm grain size sample, having the smallest grain size, showeda relatively lower average static hardness value of 25.4171.0GPa. The dynamic hardness (HVd) was lower than the statichardness for all the grain sizes in this load range. Unlike thestatic hardness values, the dynamic hardness values exhibited a

50 µm

(a)

50 µm

(b) Static, 20 N Dynamic, 21 N

Fig. 2. Static and dynamic indentation-induced damages along the bonded interface of 1.6 mm grain size boron carbide specimen.

(a) 1650_5

5 µm 5 µm

1750_5 (b)

2.33 g/cc 2.50 g/cc

Fig. 3. Scanning electron microscopy (SEM) micrographs of fractured surfaces of sintered boron carbide samples at two different processingtemperatures. Note the difference in densities mentioned on the SEM micrographs.

1852 Journal of the American Ceramic Society—Ghosh et al. Vol. 90, No. 6

greater scatter for all the grain sizes as can be seen from Fig. 5and Table II. The 1.6 mm grain size boron carbide showed asignificant decrease in dynamic hardness compared with theother two grain sizes. This decrease is probably due to the lowerdensity (see Table I) or the higher level of porosity in thesespecimens. For the other two grain sizes, the trends betweenstatic and dynamic hardness values were difficult to concludebecause of a large scatter in the values and lack of sufficientnumber of data points. More in-depth discussions on the influ-ence of porosity will be presented in Section IV.

(3) Indentation Crack Lengths and Fracture ToughnessA comparison of the optical micrographs of the top surface ofthe indented regions of a 2.7 mm grain size specimen at an in-dentation load of 2.94 N is shown in Fig. 6. It can be seen thatthe dynamic indentations caused more severe damage comparedwith the static indentations. Measurements of crack lengths re-vealed that, in general, the half-median crack length (c) in-creased linearly with indentation load under both low and highstrain rate loads. The average crack lengths were slightly longerunder dynamic loads compared with static loads for all the grainsizes. Using the Evans and Charles equation,30 static and dy-namic fracture toughness (K s

C and KdC, respectively) values were

calculated and plotted in Fig. 7. A wide scatter in the fracturetoughness values was observed under both conditions. In gen-eral, fracture toughness was lower under dynamic loads com-pared with static loads as summarized in Table II.

Clearly, boron carbide revealed lower hardness (Fig. 5) andlower fracture toughness (Fig. 7) under dynamic indentation thanthose under static indentation. But this trend in the loss of me-chanical properties of boron carbide at high strain ratescontradicts the established trend for many other engineering ma-terials where an increase in hardness,22–24 yield strength,34–37 frac-ture toughness,37,38 and fracture strength39,40 has been observedunder a higher strain rate loading. To investigate the underlyingcause for this anomalous behavior in boron carbide, further stud-ies using Raman spectroscopy were conducted in the indentedregions beneath the surface as discussed in the following.

(4) Subsurface Damage and Raman Spectroscopy

Using a bonded specimen, the subsurface damage was analyzed.Scanning electron micrographs of the half indents (shown inFig. 8) revealed that the subsurface damage for both typesof indentations at similar load levels is significantly different.

20 µm

1.6 µm 1750_2

20 µm

2 µm 1750_5

20 µm

2.7 µm 1750_30

Fig. 4. Optical micrographs of the etched sintered boron carbide samples sintered at various times at 17501C. The average grain sizes are indicated onthe optical micrographs.

16

18

20

22

24

26

28

30

32

2 4 6 8 10 12 14 16Load (N)

Har

dnes

s (G

Pa)

Fig. 5. Comparison of static and dynamic hardness values for threegrain sizes of boron carbide.

June 2007 Dynamic Indentation Response of B4C 1853

Under dynamic indentation, the extent of damage in the lateraldirection was significantly larger than static indentation andseveral cracks were observed to emanate from the boundary ofthe damage region as indicated in Fig. 8(b). Such cracks are notevident under static indentations. These subsurface studies fur-ther confirm that boron carbide is more prone to damage underdynamic loading compared with static loading at similar loadlevels. Moreover, it can be clearly seen that the damage zonesappear to be half-penny shaped beneath the indentation as as-sumed previously for calculation of (KC) in Section (2).

Figure 9 shows Raman spectra obtained from the undamagedsurface as well as from the damaged regions of boron carbideshown in Fig. 8. Raman spectrum from the undamaged polishedsurface in the neighborhood of the indents is consistent with thatreported in the literature for single crystal and polycrystallineboron carbide ceramics.17,18 The various peaks in the spectrumcan be related to the crystal structure of boron carbide41–43 asfollows: boron carbide has a complex rhombohedral crystalstructure containing eight icosahedrons and one linear chain ofthree atoms. Each icosahedron (B11C) consists of 11 boron (B)

atoms and one carbon (C) atom residing in one of the polarsites. The linear chain consists of CBC atoms. The icosahedronsare located at the corners of the unit cell and one of the longestdiagonals along the /111S direction contains the linear chain.The two broad peaks in the lower frequency range (at 275 and325 cm"1) could be related to the chain–icosahedral linkages.The appearance of the next two narrow peaks (at 478 and 532cm"1) has been assigned to the rotation of the CBC chain aboutan axis perpendicular to the [111] direction and the liberationalmode of the B11C icosahedron, respectively. Broad peaks in thehigh-frequency range (between 600 and 1200 cm"1) are associ-ated with the B11C icosahedrons. Apart from these characteristicpeaks, a small peak near 1580 cm"1 was observed in the Ramanspectrum of the undamaged surface. This peak is attributed tothe so-called G peak (graphite peak)18,44 due to the presence offree carbon in the boron carbide specimens.

Raman spectra obtained from the indentation-induced dam-aged zones showed characteristics different from those of theundamaged surface. The intensities of all the peaks decreasedcompared with that of the undamaged surface. But the mostsignificant change was in the high-frequency range where abroad peak around 1340 cm"1 was observed. This particularpeak was noted to be the most intense peak in the Raman spec-tra collected from the dynamically damaged regions. This peakis identified as the so-called D peak (disordered graphite peak)characteristic of amorphous carbon.18 Clearly, these changes inthe Raman spectra indicate a structural change induced by theindentations in boron carbide crystal structure. Also, the extentof structural change or disorder was observed to be significantlyhigher under dynamic indentation than that under static inden-tation. In the following section, we will summarize all the aboveresults and provide a rationale for the strain rate dependence ofindentation-induced structural changes and grain size effects inboron carbide.

IV. Discussion

In general, sintering of boron carbide by conventional sinteringmethods at temperatures in the range of 17001–18001C produceshighly porous and coarsened microstructures,9,11,12 because inthis temperature range grain coarsening dominates over particle

20 µm

(a)

20 µm

(b)

Fig. 6. Optical micrographs of (a) static and (b) dynamic indents for 2.7 mm grain size boron carbide at 2.94 N. Note longer cracks and more severedamage in dynamic indent.

2.2

2.7

3.2

3.7

4.2

2 4 6 8 10 12 14 16

Load (N)

Kc

(MP

a.m

0.5 )

Fig. 7. Comparison of static and dynamic indentation fracture tough-ness for three different grain sizes of boron carbide.

Table II. Comparison of Average Dynamic and Static Hardness and Fracture Toughness Values

Grain size (mm) HVs (GPa) HVd (GPa) % change #HVs"HVd$HVs

KsC (MPa %m0.5) Kd

C (MPa %m0.5) % Change#K s

C"Kd

C$

K sC

1.6 25.4171.00 18.1772.00 28.49 3.3970.25 2.6970.27 20.652.0 27.4572.00 26.1572.50 4.74 3.6170.30 3.0070.29 16.892.7 27.4572.50 23.3773.00 14.86 3.2270.16 2.8170.36 12.73

1854 Journal of the American Ceramic Society—Ghosh et al. Vol. 90, No. 6

densification. This is also aided by long consolidation times thattypically extend into hours. On the contrary, the P2Cs methodused here produced fine-grained dense (above 99% of the the-oretical) boron carbide at 17501C within a short consolidationtime of 5 min. Thus, in the P2Cs method, densification of par-ticles dominated over coarsening under appropriate processingconditions. This can be seen from Fig. 3 where the residual po-rosity decreased with consolidation temperature. Therefore, itcan be rationalized that rapid localized heating, particle rear-rangement, and deformation at the interparticle contact areas(both of which are characteristic of the P2Csmethod) facilitatedthe low-temperature sintering of boron carbide. Also, the fineparticle size of the starting powder may have aided in loweringthe sintering temperature, thereby resulting in minimal graingrowth (see Fig. 4). Clearly, the temperature and the consolida-tion times used were much lower compared with those used inthe conventional sintering methods. Thus, P2Cs provides anattractive non-conventional sintering method to fabricate fine-grained boron carbide ceramics at lower temperatures within ashort consolidation time and without the use of any sinteringadditive.

It is well-known that hardness increases with a decrease ingrain size (the Hall-Petch relationship).45,46 Similarly, a decreasein residual porosity also improves the mechanical properties.But in the current work the 1.6 mm grain size boron carbide,

having the smallest grain size, revealed the lowest static hard-ness, which does not follow the well-known Hall–Petch rela-tionship. No significant variation in hardness was observed

23.8 µm

23 µm

10 µm

68 µm

24.7 µm

17.3 µm

10 µm

37.5 µm

(a)

(b)

Fig. 8. Subsurface-damaged region beneath (a) a static indentation and (b) a dynamic indentation at a load of 19.6 N for 1.6 mm grain sizeboron carbide.

100 300 500 700 900 1100 1300 1500 1700

Raman shift (cm!1)

Inte

nsity

(a.u

)

325

532 478

275

725

832

1090

1000

1580

1340

1340

Static Indentation

Undamaged

Dynamic Indentation

D-peak

G-peak

Fig. 9. Raman spectra from undamaged region and damaged regionsof static and dynamic indentations. Note the evolution of a strong Dpeak in dynamic indentation.

June 2007 Dynamic Indentation Response of B4C 1855

between 2 and 2.7 mm grain size boron carbide specimens. Thepossible reason for such behavior may be the presence of a higherlevel of residual porosity in small grain size boron carbide thanthat in the other two boron carbide specimens, which have morethan 99% theoretical density (see Table I). This increase in po-rosity may offset the effect of smaller grain size on hardness. Asthe larger grain size specimens did not show any variation instatic hardness and they contain negligible porosity, it can besurmised that hardness is less influenced by the grain size varia-tion (in the current grain size range) than the porosity level at alow strain rate loading. Under dynamic loading, hardness wasobserved to change significantly with grain size variation (Fig. 5and Table II). But as mentioned earlier, due to statistically insuf-ficient amount of data, it is difficult to make any definitiveconclusion, except that the dynamic hardness for 1.6 mm grainsize was significantly lower than static hardness. Again, the in-creased level of porosity might have offset the effect of smallergrain size, which resulted in the observed decline in dynamichardness in 1.6 mm grain size specimens. Therefore, it can be in-ferred that the hardness of boron carbide is not only influencedby grain size and porosity level but also by strain rate. A signif-icant lowering of dynamic hardness compared with static hard-ness in 1.6 mm grain size boron carbide indicates that residualporosity could be more detrimental to hardness under dynamicloading than under static loading. A similar conclusion can bemade from the fracture toughness measurements (Fig. 7 and Ta-ble II) where a large decline in fracture toughness was observedfor 1.6 mm grain size compared with other grain sizes underdynamic loading.

The observed decrease in hardness and fracture toughness inboron carbide could be a consequence of some form of weak-ening mechanism that offsets the effect of strengthening, gener-ally observed in engineering materials at high strain rates. Theballistic resistance of boron carbide was found to decreasedramatically at projectile velocities above 850 m/s.19 A similardecline in shear strength was also observed when subjected toshock above the Hugoniot elastic limit (HEL).20,21 It has beenfound that boron carbide exhibits a localized solid-state am-orphization when subjected to high-velocity impacts.19 Recently,Fanchini et al.47 have predicted the origin of solid-state am-orphization in boron carbide from the collapse of B12(CCC)polytype based on Gibbs free-energy calculations using the den-sity functional theory. In our study, Raman spectroscopy clearlyrevealed that the D peak was more pronounced under dynamicloading than under static loading. Such spectra for damagedregions of static and dynamic indentations for boron carbidehave not been reported in literature. Domnich et al.17 reportedRaman spectra from the damaged regions of nanoindentationon boron carbide single crystal and noted a similar D peakaround 1340 cm"1 from which they predicted a structural dis-order or localized amorphization. Ge et al.18 carried out similarnanoindentations on polycrystalline boron carbide and con-firmed from transmission electron microscopy the localized am-orphization. Therefore, the appearance of the D peak, in ourstudy, clearly indicates that evolution of a D peak is associatedwith an indentation-induced deformation (or damage) in theform of some structural change. As the D peak corresponds to adisordered graphite peak,48 the appearance of a D peak in thedeformed boron carbide definitely indicates a disordered graph-ite state that may result from the local collapse in crystal struc-ture. Therefore, based on the results obtained in the currentwork and from other researchers,17,18 we also suggest a similarlocalized amorphization triggered by the static and dynamic in-dentations. The appearance of a stronger D peak in dynamicindentation may also indicate a higher level of structural trans-formation compared with the static indentation.

Therefore, dynamic indentation initiates a form of weakeningmechanism in boron carbide through the formation of a local-ized amorphous phase. The observed decrease in dynamic hard-ness and fracture toughness could be a direct result of thegreater extent of structural damage under dynamic loadingthan that under static loading. As the amorphous phase is

more brittle and relatively weaker than the crystalline phase,19

the higher the amount of amorphous phase, the greater the ex-tent of weakening in the material. This result is verified from thesubsurface damage studies (Fig. 8), extent of crack lengths(Fig. 6), and fracture toughness measurements (Fig. 7). There-fore, in this range of loading rate and grain sizes, pressure-as-sisted solid-state amorphization offsets the inertial effects ondislocation motion and crack propagation under dynamic load-ing and lowers the strength of boron carbide ceramics.

V. Conclusions

(1) Boron carbide powder sintered using the P2Cs tech-nique resulted in near theoretically dense and fine-grained boroncarbide ceramic disks at significantly lower processing time andtemperatures than in conventional methods.

(2) The boron carbide ceramics showed a consistent de-crease in hardness and fracture toughness as well as a greaterextent of damage under dynamic indentations than under staticindentations. The presence of residual porosity was observedto lower dynamic hardness significantly compared with statichardness.

(3) Raman spectroscopic analysis of the indented regionssuggested a crystalline to amorphous structural change in boroncarbide. The dynamic indentations resulted in a greater level ofstructural change compared with the static indentation. Thelower hardness, fracture toughness, and a greater level of dam-age under dynamic loading compared with static loading couldbe attributed to the formation of a greater amount of localizedamorphous phase that is weaker than the original crystallinephase in boron carbide.

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