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Materials Science and Engineering A 531 (2012) 61– 69

Contents lists available at SciVerse ScienceDirect

Materials Science and Engineering A

jo ur n al hom epage: www.elsev ier .com/ locate /msea

emperature and load dependent mechanical properties of pressureless sinteredarbon nanotube/alumina nanocomposites

oumya Sarkar, Probal Kr. Das ∗

on-oxide Ceramics and Composites Division, CSIR-Central Glass and Ceramic Research Institute (India), Kolkata 700032, India

r t i c l e i n f o

rticle history:eceived 18 February 2011eceived in revised form 5 September 2011ccepted 3 October 2011vailable online 19 October 2011

eywords:

a b s t r a c t

Multiwalled carbon nanotube (MWCNT)/alumina (Al2O3) nanocomposites having different CNT contentshave been fabricated by wet-mixing and pressureless sintering in Argon at 1700 ◦C. Depending on inden-tation load, the highest improvement of 15–34% in hardness was achieved in 0.3 vol.% MWCNT/Al2O3

nanocomposite. Indentation size effect (Mayer’s exponent = 1.753) on hardness was the most prominentin 1.2 vol.% MWCNT/Al2O3 nanocomposite due to presence of clustered CNTs, non-uniform interface andporous microstructure. ∼34% increase in fracture toughness was achieved in 0.3 vol.% MWCNT/Al2O3

0.5

ompositesechanical characterization

lectron microscopy

nanocomposite than pure Al2O3 (∼3.84 MPa m ). High MWCNT loaded nanocomposites had reducedR-curve sensitivity because of increased matrix grain refining effect and reduced matrix compressiveresidual stress. High temperature flexural strength results indicated that strength retention of nanocom-posites up to 1100 ◦C in ambient was much better compared to pure Al2O3. To predict structure–propertyrelationship in nanocomposites, detailed microstructural and fractographic studies were also performed.

. Introduction

CNTs being the excellent allotrope of carbon received immensenterest as efficient reinforcing phase in advanced structural andrmor ceramic matrix nanocomposites [1–5]. From mechanicalspect, reinforcement of ceramic matrices with CNT can produceougher, stronger and harder nanocomposites than unreinforcedrittle matrix by effective load sharing, crack deflection, crackridging and CNT pull-out [5,6]. However, the primary limitationf using CNT in ceramic matrix, especially in Al2O3, is poor dis-ersion of CNT in matrix due to chemical incompatibility betweenonstituents, hydrophobicity of CNT and strong Van der Waalsttractive forces among tubes that lead to presence of clusteredNTs with no load carrying capacity and poor densification and

nterface performance [7]. In addition, high temperature (>1250 ◦C)xposure during sintering can also destroy structural integrity ofNT that failed to offer any reinforcing effect in final nanocomposite8]. Attempts have already been made to effectively disperse CNTsn Al2O3 and to maintain structural reliability of CNTs by using com-aratively low-temperature sintering methods, e.g. hot-pressingnd spark plasma sintering [5–7,9–15,17–30]. However, even after

mploying such advanced sophisticated techniques to achievefficient CNT/Al2O3 nanocomposites, consistency in mechanicalroperty data is still lacking. For better realization, a brief review

∗ Corresponding author. Tel.: +91 33 2473 3469/76/77/96; fax: +91 33 2473 0957.E-mail address: probal@cgcri.res.in (P.Kr. Das).

921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2011.10.024

© 2011 Elsevier B.V. All rights reserved.

on mechanical properties of Al2O3 matrix nanocomposites rein-forced with either MWCNT or singlewalled CNT (SWCNT) is given inTable 1 [5,7,10–30]. Hence, further research work on this particularceramic matrix nanocomposite is still required to get consis-tent mechanical performance suitable for real-life applications.The present work aimed at mechanical property evaluation ofMWCNT/Al2O3 nanocomposites containing 0.15–2.4 vol.% MWCNTfabricated by wet-mixing of as-received precursors followed bypressureless sintering in Argon. Indentation size effect (ISE) onhardness and crack growth resistance (R-curve) behavior wereevaluated to assess load dependence of mechanical properties ofnanocomposites. High temperature flexural strength (�FS) tests (upto 1100 ◦C) were also performed to determine extent of strengthretention of nanocomposites with respect to their room temper-ature strength. Mechanical properties of nanocomposites werecompared with those obtained for pure Al2O3 to get an ideaon effectiveness of MWCNT reinforcement in Al2O3 for struc-tural applications. Microstructural and fractographic analysis wereperformed to explore structure–property relationship in nanocom-posites.

2. Experimental

2.1. Raw materials and sample fabrication

MWCNT (>95 wt.% pure, Shenzhen Nano Port Co., China) andAl2O3 powder (A-16-SG, 99.8 wt.% pure, Almatis, ACC Ltd., India)

62 S. Sarkar, P.Kr. Das / Materials Science and Engineering A 531 (2012) 61– 69

Table 1Mechanical properties of CNT/Al2O3 nanocomposites available in literature.

CNT type Method Content % change in KIC % change in HV % change in �FS Ref.

MWCNT

SPSa 0.9/1.9/3.7 vol.% 36 (↑)/30 (↑)/4 (↓) 3/8/24 (↓) 40 (↑)/30 (↑)/10 (↓) [5]PSb 1/3/5 vol.% 24/6/9 (↑) 4/23/46 (↑) 35 (↑)/10 (↓)/30 (↓) [16]SPSa 1 vol.% – 25 (↑) – [29]SPSa 1.5/3.3 vol.% 10/5 (↑) 5/7 (↑) – [17]SPSa 3.5 vol.% 67 (↑) – – [7]HPc 3/6 vol.% 79/56 (↑) – 13 (↑)/9 (↓) [18]HPc 12 vol.% 6 (↑) – 32 (↓) [11]HPc 12 vol.% 80 (↑) – 4 (↓) [11]SPSa 0.1 wt.% 32 (↑) 4 (↑) – [19]SPSa 0.1 wt.% 32 (↑) – – [14]SPSa 0.18/1.07/2.48 wt.% 33/65/152 (↑) 3/5/13 (↓) – [27]HPc 1 wt.% 2.5 (↑) – 12 (↑) [20]SPSa 1 wt.% – 0.8 (↑) – [28]HPc 2.5 wt.% 94/65 (↑) 13 (↑)/7 (↓) 6 (↑)/22 (↓) [21]SPSa 3.19/7.39/8.25/19.1 wt.% 1.3 (↓)/21 (↑)/36 (↓)/66 (↓) 67 (↓)/8.4 (↑)/3.4 (↑)/47 (↓) – [15]HPc 5 wt.% 80 (↑) – 25 (↑) [30]SPSa 0.5–10 wt.% 27–61 (↓) – 32–79 (↓) [12]HPc 2.7/4.1/10.5/12.5 wt.% – 14 (↑)/29 (↑)/16 (↓)/42 (↓) – [22]

SWCNT

SPSa 5.7/10/15 vol.% 139/194/– (↑) 1.5/21/– (↓) – [23]SPSa 5.7/10 vol.% 139/194 (↑) 1.5/21 (↓) – [24]SPSa 10 vol.% ∼200 (↑) – – [25]HPc 1 wt.% 103 (↑) – 19 (↑) [26]

(↑) increase; (↓) decrease. Some numerical data have been approximated from graphs/plots of corresponding report.

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a Spark plasma sintering.b Pressureless sintering.c Hot pressing.

ere used as starting precursors. Diameter and length of MWC-Ts were between 60–100 nm and 5–15 �m, respectively. FiveWCNT/Al2O3 nanocomposites were fabricated by dispersing as-

eceived MWCNTs in isopropyl alcohol by ultrasonic agitation for h in an ultrasonic bath (Oscar Ultrasonic Pvt. Ltd., India) followedy mixing of dispersed MWCNT slurry with aqueous suspensionf Al2O3 by magnetic stirring for further 1–2 h. The MWCNT/Al2O3lurry was then dried to remove volatiles. Green billets were pre-ared by cold isostatic pressing at 150 MPa and sintered at 1700 ◦Cor 2 h in a graphite resistance heating furnace (1000-4560-FP20;hermal Technology Inc., USA) in static Argon at 10 ◦C/min heatingate. Pure Al2O3 were also prepared by the same method. Nomen-lature, theoretical and sintered density and apparent porosity ofll nanocomposites and pure Al2O3 are given in Table 2.

.2. Mechanical properties evaluation

.2.1. Vickers hardness (HV)Hardness of sintered specimens was evaluated using two Vick-

rs hardness testers capable of operating from 10 gf to 2.0 kgf

402 MVD, Wolpert-Wilson, Germany) and 1 kgf to 100 kgf (Wolpert,ermany) at different loads from 4.9 N to 98 N with 10 s dwell atach test load. HV values were calculated using standard formula

able 2asic physical properties of samples sintered at 1700 ◦C for 2 h in Argon.

Sample MWNT loading (vol.%) TDa (g/cc) RDb (%) APc (%)

A0 0.00 3.970 99.50 0.132A1 0.15 3.967 99.47 0.000A2 0.30 3.963 98.92 0.140A3 0.60 3.957 98.69 0.000A4 1.20 3.944 91.27 7.790A5 2.40 3.917 88.49 9.318

a TD, theoretical density calculated using “Rule of mixture” and taking �Al2O3=

.970 g/c.c; �MWCNT = 1.775 g/c.c.b RD, relative density measured by Archimedes water immersion technique.c AP, apparent porosity measured by Archimedes water immersion technique.

[31]. 10–15 indents were made on each sample to verify consis-tency of data.

2.2.2. Fracture toughnessIndentation fracture (IF) technique as proposed by Lawn et al.

[32] was used to measure fracture resistance (KR) of sintered spec-imens at each indentation load by direct measurement of cracklengths produced from corners of Vickers impression. KR valueswere utilized to predict R-curve behavior of present nanocom-posites. Fracture toughness (KIC) values of specimens were alsodetermined by single edge notched beam (SENB) method usingASTM C-1421-09. All tests were carried out in a four-point flexuraltester (422, Netzsch-Geratebau GmbH, Germany) at constant stressrate of 0.03 N/mm2 s. Outer and inner span of test fixture were40 mm and 20 mm, respectively. Average KIC values were evalu-ated from 5 valid tests in which specimens were failed exactly fromnotched region.

2.2.3. Ambient and elevated temperature flexural strengthRoom temperature (∼40 ◦C) and high temperature (viz. 600 ◦C,

900 ◦C and 1100 ◦C) 4-point flexural strength test of sintered spec-imens were carried out according to ASTM-standard 1161-02Cand C1211-02, respectively. The same instrument and test fixtureas mentioned in Section 2.2.2 were used. For high temperaturestrength measurement, sample positioned inside test fixture wasplaced within the furnace attached with the instrument and heatedup to desired temperature. After reaching the target temperature,whole assembly was allowed to soak for 5 min. Followed by this,load was applied at specified cross-head speed.

2.3. Microstructure characterizations

For microstructure and fractographic analyses, polished andthermally etched sintered samples and indented and flexure testedsamples, respectively, were viewed using field emission scanningelectron microscope (FESEM, Supra-35, VP-Carl Zeiss, Germany).

S. Sarkar, P.Kr. Das / Materials Science and Engineering A 531 (2012) 61– 69 63

3

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Fig. 1. Presence of clustered CNTs in sintered A5 nanocomposite; scale 200 nm.

. Results and discussion

.1. Physical properties and microstructure of sintered specimens

Table 2 shows physical properties of sintered nanocompositesnd pure Al2O3. It may be seen from Table 2 that after sinteringt 1700 ◦C, although, nanocomposites containing up to 0.6 vol.%WCNTs were ∼99% dense having very low apparent porosity,

eyond that concentration, samples (i.e. A4 and A5) were porousecause of significant inhibition of consolidation primarily dueo presence of non-uniformly distributed and clustered CNTs inl2O3 matrix (Fig. 1) that acted as pores of equal dimension andestricted material transport through Al2O3 grain boundaries dur-ng sintering. On contrary, well dispersed MWCNTs were found touccessfully bridge Al2O3 grains in nanocomposites containing upo 0.6 vol.% MWCNT (Fig. 2). FESEM image of pure Al2O3 sinteredt 1700 ◦C (Fig. 3) indicates dense microstructure having mostlyquiaxed Al2O3 grains of 6 �m to 10 �m size. Formation of smallerl2O3 grains with increasing MWCNT content was observed dur-

ng microstructural analysis of present nanocomposites. Such grain

efining by CNT incorporation in Al2O3 has also been reported bythers [5,16,24,30]. Fig. 4 shows FESEM image of sintered A3 spec-men having smaller grains (6.06 ± 2.43 �m) than pure Al2O3. In

ig. 2. Isolated MWCNTs bridging Al2O3 grains in sintered A3 nanocomposite; scale00 nm.

Fig. 3. Grain morphology of sintered pure Al2O3; scale 10 �m.

both figures a few porous regions were also detected which weremost possibly produced from closed pores present in samples andcame out during surface preparation for microstructure analysis.

3.2. Hardness

Average HV values of all specimens along with related standarddeviations are plotted in Fig. 5 from which it is clear that exceptA4 specimen, all other nanocomposites had superior hardness thanpure Al2O3 irrespective of indentation load. The highest increase inHV was obtained for A2 nanocomposite which was 15–34% higherthan A0 depending on applied load. On the other hand, at the high-est load, i.e. 98 N, nanocomposite containing 1.2 vol.% MWCNT hadhardness ∼1.5% below A0 (13.81 ± 0.45 GPa). Changes in HV valuesof present nanocomposites matched closely with available litera-ture data on CNT/Al2O3 nanocomposites presented in Table 1. FromFig. 5 it is also evident that all specimens exhibited clear ISE withinpresent range of indentation loads and depending on composition,extent of ISE differed from one to another. As applied indenta-tion load (P) crossed ‘low-load hardness’ region (200 gf < P < 2 kgf)and entered ‘normal hardness’ region (P > 2 kgf) [33], reduction in

hardness was more dramatic irrespective of sample composition.However, A4 specimen encountered the most distinct effect ofapplied load on hardness at ‘normal hardness’ region possibly due to

Fig. 4. Grain morphology of sintered A3 nanocomposite; scale 10 �m.

64 S. Sarkar, P.Kr. Das / Materials Science and Engineering A 531 (2012) 61– 69

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Table 3Regression analysis result of experimental data of all samples using Meyer’s law.

Sample log(A) n R2

A0 3.707 1.820 0.999A1 3.914 1.922 0.999A2 3.906 1.902 0.999

ig. 5. Variation in Vickers hardness of tested materials with applied test load.

resence of clustered CNTs around Vickers impression which acteds defects of similar dimension and poor interfacial performancehat failed to sustain such higher loads and produced higher plasti-ally deformed zone around the impression by lowering extent oflastic recovery on unloading.

To analyze extent of variation in HV at different indentationoads, three most widely accepted empirical law/model namely

eyer’s law [34–36], proportional specimen resistance (PSR) model34] and modified PSR (M-PSR) model [36] were utilized andompared. As per classical Meyer’s law, ‘P’ is related to Vickersmpression diagonal (d) as [34–36]:

= Adn (1)

here A is constant and n is Meyer’s exponent. According to this law,hen hardness of a material is independent of ‘P’, Meyer’s exponent

hould be 2. However, load dependent hardness should result in ‘n’alues less than 2 and larger deviation from 2 signifies higher ISEn test material. ‘A’ and ‘n’ can be evaluated from linear regressionnalysis of log(P) versus log(d) data plot using following logarithmicorm of Eq. (1):

og(P) = log(A) + {n ∗ log(d)} (2)

Results obtained from regression analysis (Fig. 6) of Eq. (2) ofll samples are given in Table 3. log(A) and ‘n’ values obtained for

ig. 6. Plot of log(P) versus log(d) according to Meyer’s law for tested materials.

A3 3.855 1.877 0.999A4 3.647 1.753 0.998

present specimens matched well with reported literature data forhard and brittle materials having no ability of work-hardening dur-ing indentation [35,36]. Since, ‘n’ values were less than 2 for allsamples, therefore, all of them experienced ISE. It may be seen fromTable 3 that except A4 (n = 1.753), ‘n’ of A0 (1.820) was lower thanthose (1.877–1.922) obtained for nanocomposites containing up to0.6 vol.% MWCNT. Therefore, while the lowest ISE was offered by A1nanocomposite, the highest ISE was encountered by A4 followed bypure Al2O3.

PSR model was used to quantitatively describe ISE of presentspecimens. According to Li and Bradt [34], resistance offered bya material during indentation is not constant and increases withincreasing indentation impression and ‘P’ and ‘d’ should be relatedas:

P = a1d + a2d2 (3)

where a1 and a2 are constants and related to the elastic and plasticdeformation properties of test material, respectively [34,36]. Theterm a2 in Eq. (3) is related to load-independent true hardness (HVT)of material according to the following equation:

HVT = k ∗ a2 (4)

where k = 1.8544 for a Vickers indent. (P/d) versus d plots of spec-imens are shown in Fig. 7. Results showed that only A4 specimentraced two distinct region of apparent linearity (R2 > 0.999) thatfall into ‘low-load hardness’ and ‘normal hardness’ region. Thus,it was inappropriate to use PSR model for describing ISE of thisnanocomposite because it might produce several inconsistencies.However, other nanocomposites and pure Al2O3 showed betterlinearity (R2 = 0.996–0.999) throughout applied load range. Regres-sion analysis results are given in Table 4 along with calculatedHVT values which were lower than experimental HV values. Differ-

ence between HVT and HV data possibly caused by the presence ofthermal and/or machining induced residual stresses in specimens.Although, HVT of Al2O3 obtained in present study matched closelywith previous reports [36,37], since no literature data are available

Fig. 7. Plot of (P/d) versus d according to PSR model for tested materials.

S. Sarkar, P.Kr. Das / Materials Science and Engineering A 531 (2012) 61– 69 65

Table 4Regression analysis result of experimental data of all samples using PSR model.

Sample a1 (N/mm) a2 (N/mm2) R2 HT (GPa)

A0 86.11 6825.36 0.996 12.66Al 40.09 9449.23 0.999 17.52A2 58.27 9482.53 0.998 17.58A3 72.32 8736.43 0.997 16.20

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crack length dependence of KR, following power law relation wasemployed [42]:

KR = k ∗ c˛ (6)

Fig. 8. Plot of P versus d according to M-PSR model for tested materials.

n HVT of CNT/Al2O3 nanocomposites, direct comparison was notossible.

Recently, Gong et al. [36] introduced an additional term to incor-orate residual stress effect due to machining in observed ISE ofaterials. The M-PSR model is expressed as:

= P0 + a1d + a2d2 (5)

here P0 accounts for residual stresses appeared in test specimenue to surface machining and physical significance of a1 and a2 arehe same as stated in Eq. (3). The second order polynomial fit ofP’ and ‘d’ data of all specimens are plotted in Fig. 8 and results areiven in Table 5. In all cases, polynomial fit resulted in R2 ≥ 0.999ndicating Eq. (5) was sufficiently suitable for describing experi-

ental load/size data of present ceramic matrix nanocomposites.rom P0 values given in Table 5 it is evident that all specimensad obvious effect of surface machining on ISE which were neg-tive in magnitude. However, since no regular trend in P0 wasbserved, therefore, direct correlation with microstructure and/orachining effect was not possible. Difference between HVT values

f specimens evaluated from M-PSR model and experimental HVatched closely with results from PSR model (Table 5). A well-

efined trend in (a1/a2) ratio of nanocomposites (Table 5), obtainedrom M-PSR model further indicates that with increase in con-entration of extremely elastic MWCNTs in Al2O3 matrix from

.15 vol.% to 1.2 vol.%, extent of post-indentation elastic recoveryy nanocomposites increased significantly than pure Al2O3 whilelastic deformation of nanocomposites reduced.

able 5egression analysis result of experimental data of all samples using M-PSR model.

Sample P0 (N) a1 (N/mm) a2 (N/mm2) R2 HT (GPa) (a1/a2)

A0 −0.81 125.99 6491.64 0.999 12.04 0.019A1 −0.29 56.02 9298.77 0.999 17.24 0.006A2 −2.71 189.74 8312.56 0.999 15.41 0.023A3 −4.50 283.91 6920.84 0.999 12.83 0.041A4 −2.81 263.77 5346.70 0.999 9.91 0.049

Fig. 9. Plot of log(C) versus log(P) of all specimen for applied load range from 4.9 Nto 98 N; slope and square of correlation coefficient (R2) of linear fit to experimentaldata points are given adjacent to each specimen identification.

3.3. Fracture toughness

KR values of all specimens at each indentation load were cal-culated using direct crack measurement (DCM) method. It wasobserved that irrespective of specimen type, at low loads (<19.6 N)most of the cracks generated from corners of Vickers impressionwere Palmqvist in nature having crack length (c) to half of theimpression diagonal (a) ratio below 2.5 whereas at 49 N or 98 N,cracks were mainly median type (c/a > 2.5). The log–log plot ofmeasured ‘c’ versus ‘P’ of all specimens is shown in Fig. 9. Linearfit to this experimental data revealed that slopes obtained for allspecimens (0.664–0.693) were matched well with standard slopeof 2/3 of log(c) versus log(P) plot according to indentation frac-ture mechanics and other reports [38,39]. Therefore, to measureKR values of present samples at each applied load, two equationsproposed by Niihara et al. [40,41] which are separately applicablefor Palmqvist and median crack systems were utilized. Average KR

values of all samples are plotted in Fig. 10 against ‘c’ values. To ana-lyze the so-called R-curve behavior of present nanocomposites, i.e.

Fig. 10. Fracture resistance (KR) as a function of crack length.

66 S. Sarkar, P.Kr. Das / Materials Science an

Fig. 11. Fracture toughness (KIC) obtained by SENB technique versus specimen type.

Fig. 12. Flexural strength (�FS) of all specimens as a function of test temperature.

Fig. 13. Extent of change in flexural strength of nanocomposites compared to pureAl2O3 as a function of test temperature.

d Engineering A 531 (2012) 61– 69

where k and are constants. The toughness exponent ‘˛’ is an esti-mation of sensitivity towards R-curve behavior. It may be seen fromFig. 10 that ‘˛’ values obtained after plotting experimental KR and‘c’ of all specimens using Eq. (6) lay well within available ‘k’ and‘˛’ values reported for various Al2O3 and Al2O3 matrix compos-ites [42–45]. It may be seen from data given in Fig. 10 that althoughnanocomposites containing MWCNTs up to 0.6 vol.% showed eithersame or improved crack growth resistance behavior compared toA0 ( = 0.092), 1.2 vol.% MWCNT/Al2O3 nanocomposite had lowerR-curve sensitivity than pure Al2O3. Besides toughening effectoffered by matrix Al2O3 grains, additional increase in R-curvebehavior of A1, A2 and A3 specimens was achieved by satisfac-tory crack deflection/bridging and higher pull-out resistance bywell-dispersed bamboo structured MWCNTs [46,47] in these hightemperature sintered nanocomposites. In view of effect of grainsize and thermal residual stresses on R-curve behavior of materialsincluding various alumina matrix composites [48–51], the declinedR-curve sensitivity of high MWCNT loaded nanocomposites may beexplained by following features:

(i) Presence of increasing concentration of clustered CNTs actedas potential defect sites in composite structures offering notoughening effect and this adverse effect of agglomerated CNTswas much pronounced at higher indentation loads.

(ii) Large difference in coefficient of thermal expansion (CTE)between MWCNTs and Al2O3 also played important rolein lowering compressive thermal residual stresses in bridg-ing Al2O3 grains particularly at higher MWCNT content innanocomposites because MWCNTs have small and negativeCTE and is essentially isotropic in nature [52,53] while CTE ofAl2O3 is large, positive and anisotropic in nature.

(iii) Reduced matrix Al2O3 grain size with increased MWCNT con-tent resulted in: (a) reduced involvement of bridging Al2O3grains along crack path which were under compressive resid-ual stress and responsible for KR enhancement; (b) reducedthermal residual stresses in matrix grains that offered lowerfrictional forces during grain pull-out; and (c) reduction inextent of crack branching through which additional energydissipation can be attained.

Fracture toughness values obtained from SENB method of allspecimens are shown in Fig. 11 where vertical lines at the topof each bar represent standard deviation. KIC increased up to0.3 vol.% MWCNT and A2 specimen had the highest enhancement(∼34%) in toughness than A0 (3.84 ± 0.21 MPa m0.5). On contrary,although, nanocomposites with 0.6 vol.% and 1.2 vol.% MWCNToffered higher KIC compared to A0, extent of toughness enhance-ment decreased with increasing MWCNT content and finally, A4nanocomposite had KIC value of (4.06 ± 0.21) MPa m0.5 which wasonly ∼6% higher than that of pure Al2O3. Comparison o f KR valuesdetermined by DCM method (Fig. 10) and KIC values evaluatedfrom SENB technique (Fig. 11) indicated that although, percentageimprovement in toughness of present nanocomposites than pureAl2O3 obtained from both the methods matched closely, individualKR values obtained from DCM technique were slightly deviatingfrom KIC data obtained from SENB method since KR was responsiblefor crack arrest and KIC for fast crack propagation. Further, ahigher wake controlled crack tip shielding effect in indentationinduced short-cracks possibly resulted in higher toughness valuecompared to KIC evaluated from SENB technique [54]. Differencebetween fracture toughness data evaluated by different techniquesnamely DCM, SENB, single edge V-notched beam (SEVNB), double

cantilever beam (DCB) and surface crack in flexure (SCF) has alsobeen reported for various advanced ceramics [55–60]. However,for many industrial applications, fracture toughness evaluation bySENB, SEVNB or even SCF technique may not be feasible since they

S. Sarkar, P.Kr. Das / Materials Science and Engineering A 531 (2012) 61– 69 67

Fig. 14. (a–c) Fractographic observations of flexure tested specimens; scale 10 �m. (d) Fractographic observations of flexure tested specimens; scale 20 �m. (e) Fractographicobservations of flexure tested specimens; scale 30 �m. (f) Fractographic observations of flexure tested specimens; scale 50 �m. (g) Fractographic observations of flexuretested specimens; scale 10 �m (h–j) Fractographic observations of flexure tested specimens; scale 1 �m.

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8 S. Sarkar, P.Kr. Das / Materials Scien

ll require large specimen size and machining for precise notchreparation as well as large material volume to get reproducibleata, whereas, toughness evaluation by DCM technique requiresnly small specimen size to get averaged toughness value fromarge data points and easy experimental procedure [59–61]. Fur-her, in fracture toughness evaluation by DCM technique a series ofimple equations available for median and Palmqvist crack systemsrom which one can choose the most appropriate equation to

easure toughness value of a material. Thus, from applicationoint of view, fracture toughness of materials evaluated by DCMechnique as well as R-curve behavior are equally important asoughness and R-curve obtained from SENB technique.

.4. Flexural strength

Four-point flexural strength of nanocomposites and pure Al2O3t room temperature (∼40 ◦C) and 600 ◦C, 900 ◦C and 1100 ◦C inaboratory atmosphere are plotted in Fig. 12. Except nanocom-osites with 1.2 vol.% and 2.4 vol.% MWCNT content, other threeanocomposites showed improved �FS at each test tempera-ure. At room temperature, the highest �FS was obtained for1 (�FS = 266.03 ± 26.01 MPa) which was ∼20% higher than A0.

ncrease in �FS of A2 and A3 was ∼12% and ∼6%, respectively,han that of pure Al2O3. Although, a continuous increment in �FSf nanocomposites might be expected at higher MWCNT contentince as mentioned in Section 3.1 that matrix grain size decreasedontinuously with increasing CNT concentration, decline in extentf strength improvement of nanocomposites was observed pri-arily due to increased concentration of MWCNT agglomerates at

igher nanotube content which acted as potential strength limit-ng defect sites and poor sinterability of nanocomposites beyond.6 vol.% MWCNT that led to inadequate and uneven interface per-ormance and as a result, A4 and A5 had �FS of ∼122 MPa and64 MPa, respectively, which were ∼45% and ∼70% lower thanonolithic Al2O3. Fig. 13 showed variation in percent change in

FS of MWCNT/Al2O3 nanocomposites compared to pure Al2O3 as function of test temperature. It may be seen from the figure thatp to 600 ◦C, percent change in strength of all nanocompositeshan pure Al2O3 was the same as obtained from room temper-ture strength analysis. However, as test temperature increased,ll curves of Fig. 13 moved upward suggesting that high tempera-ure strength retention of present MWCNT/Al2O3 nanocompositesere much better than monolithic Al2O3. In addition, high temper-

ture strength data also indicated that MWCNTs present withinulk of nanocomposite specimens were successfully preservedy surrounding Al2O3 matrix and thus, were able to offer ade-uate reinforcing effect to retain �FS of nanocomposites. As aesult, at 1100 ◦C, A1 offered ∼50% higher �FS compared to A0�FS, 1100 ≈ 135 MPa) while �FS of A4 was only ∼9% lower than thatf A0. Fractographic analysis revealed that in pure Al2O3 pres-nce of individual or regions of large grains (Fig. 14a–c) mostlycted as strength limiting sites while for nanocomposites, fractureainly initiated from machining induced surface flaws (Fig. 14d).owever, surface voids (Fig. 14b), cavitations (Fig. 14c), porous

eams (Fig. 14e) and sagging (Fig. 14f) also played critical role inowering room temperature as well as high temperature flexuraltrength of pure Al2O3. Further analysis revealed intergranular asell as transgranular crack extension in A0 (Fig. 14g), whereas,

n nanocomposites, intergranular crack path was mainly detectedFig. 14h). This suggested that MWCNT/Al2O3 interface regionsere effective in successful crack deflection around the interface.

resence of bridging and pulled-out MWCNTs was also detected inexure tested MWCNT/Al2O3 nanocomposites (Fig. 14i). It was also

ound during fractographic analysis that few Al2O3 grains experi-nced shear induced slip band formation (Fig. 14j).

[

[[[

d Engineering A 531 (2012) 61– 69

4. Conclusions

It is possible to fabricate structural MWCNT/Al2O3 nanocompos-ites by cost-effective process involving wet-mixing of as-receivedprecursors followed by pressureless sintering in inert atmospherehaving superior mechanical performance than pure Al2O3. How-ever, to avoid presence of large CNT agglomerates and to achievebetter reinforcing effect, MWCNT concentration in such nanocom-posites should be kept below 0.6 vol.%. Improvement in mechanicalproperties of MWCNT/Al2O3 nanocomposites was achieved bymatrix grain refining effect of CNTs and formation of effectiveCNT/Al2O3 interface that helped in successful load transfer frommatrix to filler and facilitated crack deflection/bridging and energydissipation through CNT pull-out. To analyze ISE on hardness ofpresent nanocomposites, the M-PSR model was found to be mosteffective due to presence of “machining induced surface residualstress” term that had significant contribution on Vickers hardnessespecially for nanocomposites with porous microstructure andinsufficient interface performance having high CNT agglomerateswith no load bearing capacity. The crack growth resistance behav-ior with increasing indentation load was also the same, where,besides porous microstructure and poor interface performanceof high MWCNT containing nanocomposite, the reduced matrixgrain size and compressive thermal residual stress around bridg-ing Al2O3 grains also played important role in lowering R-curvesensitivity. However, as far as high temperature strength retentionwas concerned, MWCNT/Al2O3 nanocomposites showed betterperformance than pure Al2O3 by retaining structural integrity ofreinforcing MWCNTs by surrounding Al2O3 matrix even up to1100 ◦C in laboratory environment.

Acknowledgements

The authors express their sincere gratitude to the Director, CSIR-Central Glass and Ceramic Research Institute (CSIR-CGCRI), India forhis kind permission to publish this work. The authors are also grate-ful to the members of Analytical Facility Division of CGCRI, for theirextensive help in carrying out the microstructure analysis. The firstauthor acknowledges the financial support of the Council of Scien-tific and Industrial Research (CSIR), India.

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