Materials Science and Engineering A 508 (2009) 167–173
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Materials Science and Engineering A
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abrication and properties of dispersed carbon nanotube–aluminum composites
.M.K. Esawia,∗, K. Morsib, A. Sayeda, A. Abdel Gawada, P. Borahb
Department of Mechanical Engineering and the Yousef Jameel Science and Technology Research Center, The American University in Cairo,UC Avenue, P.O. Box 74, New Cairo 11835, EgyptDepartment of Mechanical Engineering, San Diego State University, 5500 Campanile Drive, San Diego, CA 92182, USA
r t i c l e i n f o
rticle history:eceived 8 October 2008eceived in revised form 15 December 2008ccepted 1 January 2009
Powder metallurgy techniques have emerged as promising routes for the fabrication of carbon nanotube(CNT) reinforced metal matrix composites. In this work, planetary ball milling was used to disperse 2 wt%MWCNT in aluminum (Al) powder. Despite the success of ball milling in dispersing CNTs in Al powder,it is often accompanied with considerable strain hardening of the Al powder, which may have implica-tions on the final properties of the composite. Both un-annealed and annealed Al–2 wt% CNT compositeswere investigated. It was found that, ball-milled and extruded (un-annealed) samples of Al–2 wt% CNTdemonstrated high notch-sensitivity and consistently fractured outside the gauge length during tensiletesting. In contrast, extruded samples annealed at 400 and at 500 ◦C for 10 h prior to testing, exhibitedmore ductile behavior and no notch sensitivity. Under the present investigated processing conditions,ball milling for 3 h followed by hot extrusion and annealing at 500 ◦C resulted in enhancements of around
21% in tensile strength compared with pure aluminum with the same process history. The ball-millingconditions used were found to result in the creation of a nanostructure in all samples produced, as shownby XRD and TEM analysis. Such nanostructure was retained after prolonged exposures to temperaturesup to 500 ◦C. The tensile testing fracture surfaces showed uniform dispersion and alignment of the CNTsin the aluminum matrix but also showed CNTs acting as nucleation sites for void formation during tensiletesting. This has contributed to the observation of CNT pull-out due to the poor bond between the CNTs and the matrix.
. Introduction
The interest in carbon nanotubes (CNTs) as super reinforce-ents for metallic matrices has been growing considerably over
he past few years, largely focusing on investigating their contri-ution to the enhancement of the mechanical performance of thenal composite. As in conventional composites, the orientation ofhe CNTs, homogeneity of the composite, nanotube matrix adhe-ion, nanotube aspect ratio and the volume fraction of nanotubesre expected to have significant influences on the properties of theanocomposite. Controlling such factors to obtain an exceptionalomposite is very challenging. A number of reviews have beenublished on the subject, including ones by Harris [1], Carreno-orelli [2] and Curtin and Sheldon [3]. Aluminum (Al) has benefited
rom intense research efforts in this area, where CNTs are hoped
o provide substantial improvements to its properties [4–22]. Onef the major obstacles to the effective use of carbon nanotubess reinforcements in metal matrix composites is their agglomer-tion and poor distribution/dispersion within the metallic matrix.
A number of research groups, including the current authors, inves-tigated the use of ball milling as a mechanical dispersion technique[8,9,11,14,16,22]. Different milling conditions (energy and time)were investigated. SEM and TEM images showing well dispersedCNTs proved the process to be promising [9,14,16]. However, con-cerns about the possible damage and/or amorphization of theCNTs under the harsh milling conditions have been raised [9,22]so optimization of milling conditions is necessary. CNT contentsup to 10 wt% have been investigated. Enhancements in mechan-ical properties due to CNT addition have been reported [6,9,21].The optimum CNT content at which maximum enhancement wasachieved varied depending on mixing and preparation techniques.In general, the enhancements were much lower than expected[11,13,16,21]. Although Laha et al. [17–19] confirmed the presenceof unreacted and unmelted CNTs after exposure to the high temper-atures associated with the plasma spraying process they used, mostresearchers have been cautious about processing CNT-compositesat high temperatures so as to avoid CNT damage and possible
adverse chemical interfacial reactions with the matrix. It has beenargued, however, that the formation of Al4C3 phase should not beviewed as detrimental since it can help to enhance the Al–CNT bondand can lock the nanotubes in place and hence contribute to theenhancement of the mechanical properties of the composites [6,12].
phase formation under the current ball-milling conditions. Fig. 3shows a clear distinction between the morphology of the Al–CNTpowders milled for 3 and 6 h, where 6 h milling results in pow-ders with bi-modal particle size distribution. In both cases of the3 and 6 h milling, methanol was added as a process control agent
68 A.M.K. Esawi et al. / Materials Scien
n another study [18], the formation of an ultra thin �-SiC layer in anl–Si alloy containing 10 wt% CNT synthesized by thermal sprayingas observed to enhance the limited wettability and the interfacialond between the CNTs and the aluminum matrix.
One way to produce composites with uniform distribution ofNTs within the metallic matrix is by first dispersing them in metalowders then subsequently consolidating these Al–CNT compos-
te powders. The ball-milling process, although effective, normallyesults in strain hardening of the powders, and this can have impli-ations on the final properties of the bulk composite. Thereforehe consolidation of heavily worked powders and their subsequentroperties would be important from both a practical as well as a sci-ntific standpoint. The present authors believe that results for CNTeinforced aluminum or even other metals should be viewed in lightf the condition of the powders used to form the bulk samples (e.g.egree of strain hardening) and their subsequent process historye.g. hot extrusion, annealing), a point that has not been given too
uch attention so far in the literature. Potential problems may alsorise in terms of matrix/CNT interfacial reactions and CNT damageuring high temperature processing or metal working.
In previous studies by the authors, dispersion of 2 and 5 wt%NTs in aluminum powders was achieved, and the effect of millingime on the morphological evolution of the Al–CNT compositeowders was investigated [14,15]. This paper discusses the con-olidation of ball-milled Al–CNT composite powders using hotxtrusion and the evaluation of the mechanical response of theomposites (both annealed and un-annealed) using tensile andardness testing.
. Experimental procedure
30 grams of Al (99.7% pure, −200 mesh, Aluminum Powder Com-any Ltd., UK) and 2 wt% multi-wall carbon nanotubes (averageimensions: 140 ± 30 nm outer diameter, 4–8 nm internal diameter,nd 3–4 �m in length, MER corporation, USA) were placed in 250 mltainless steel mixing jars together with 75 stainless steel millingalls (10 mm diameter); giving a ball-to-powder weight ratio of0:1. The jars were filled with argon and were agitated using a plan-tary ball mill at 200 rpm for 3 and 6 h. 2 ml of methanol were addeds a process control agent (PCA) in order to minimize cold weldingf the Al particles and also to prevent powders sticking to the ballsnd the jar walls. In the case of pure Al, 2.5 ml of methanol weredded.
Approximately 26 g of the ball-milled powder mix were com-acted in a 20 mm diameter compaction die at 475 MPa. Hotxtrusion of the homogenized compact was conducted at 500 ◦Csing an extrusion ratio of 4:1, to produce 10 mm diameter extru-ates. Tensile test samples were machined and super-finished outf the extrudates, with dimensions as shown in Fig. 1.
The mechanical response of the processed materials waslso characterized using both a Vickers Micro-hardness testerLoad = 300 g and Dwell time = 15 s) and Nanoindentation. Samplesrom the resulting extrudates were sectioned perpendicular to thextrusion direction and mounted using bakelite due to its lower
Fig. 1. Tensile test specimen dimensions.
Engineering A 508 (2009) 167–173
elasticity compared to epoxy resins. The samples were groundand polished to 1 �m. Nanoindentation tests were performed ona Nanoindenter XP (MTS systems Co., Oak Ridge, TN, USA) with adiamond pyramid-shaped Berkovich-type indenter tip. The nanoin-dentation tests were conducted according to the Oliver & Pharrapproach [23]. Continuous Stiffness method (CSM) was used. Themaximum displacement of the tip into the surface was set to a highvalue (10 �m) in order to get the maximum possible indentationdepth. Distance between adjacent indentations was set to 100 �mso as to avoid the effect of residual stresses around neighboringindentations. The thermal drift was kept below 0.05 nm/s. Thereported hardness values are averages of at least ten indentations inthe case of Vickers hardness testing and twenty indentations (5 × 4matrix) for the nanoindentation experiments. For microstructuralexamination, the specimens were cut, ground and polished to1 �m finish. Etching was performed in Keller’s solution, and themicrostructure was characterized using a field emission scanningelectron microscope FESEM (LEO supra 55). X-ray diffraction (XRD)(using Cu K�, Panalytical Xpert Pro diffractometer) was also usedfor phase analysis and crystal size determination using the Schererequation [24], and excluding instrumental effects on broaden-ing. TEM analysis was conducted on a JEOL 2010 analytical TEM(80–200 kV), having a LaB6 electron gun and a resolution of 0.19 nm.Sample preparation was conducted using conventional cutting andsubsequent electropolishing techniques.
3. Results and discussion
Fig. 2 shows the XRD scans for Al–CNT powders milled for 3 and6 h, respectively. A small peak is observed at 26◦ corresponding to(0 0 2) of graphite for powders ball-milled for 3 h, but then disap-pears after milling for 6 h. It is interesting to note that in previouspublished work that applied ball milling for 5 min, a clear peak at26◦ was not observed, despite the TEM observation of the CNTs inthe sample [16], albeit at ∼2 vol.% CNT which maybe at the limit ofXRD resolution. In our case also, in spite of clear observation of theCNTs in the fractured surfaces of all samples, as will be shown later,only a small peak for CNT (at 2 theta = 26◦) was observed for pow-ders ball-milled for 3 h. This could be due to a number of possiblereasons; including the small amount of CNTs used, the well dis-persed CNTs within the matrix, unfavorable strain/CNTs effect thatreduces the CNT peak intensity, and amorphization of CNTs. Fur-thermore, the figure does not show any clear evidence of carbide
Fig. 2. XRD scans of Al–2 wt% CNT powders ball-milled for 3 and 6 h.
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Fig. 3. SEM micrographs of Al–2 wt% CNT powders ball-milled for (a) 3 h and (b) 6 h, respectively.
and (
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time of 6 h, the mean grain size, as found by XRD analysis, was48.4 nm (for pure aluminium milled for 6 h and extruded at 500 ◦C).
Fig. 4. (a) CNTs on the surface of aluminum powder after 0.5 h milling time
o limit the Al particle welding which tends to increase the parti-le size, as noted in our previous publication [14]. In addition toarticle welding, the other competing process during ball milling
s strain hardening of the powders, which leads to a decrease inuctility and eventual fracturing of the Al particles and thus aecrease in particle size. As the milling time increased from 3 to 6 h,train/work-hardening of the powders is increased favoring morearticles fracturing than re-welding, and leading to a refinement inhe particle size.
As reported in an earlier study by the authors [14], the ball-illing process was proven to be effective in dispersing the CNTs.
ig. 4(a) shows well-dispersed CNTs on the surface of ball-milledluminum powders after 0.5 h of milling and Fig. 4(b) shows theNTs embedded between the rewelded particles after 3 h of milling.
All ball-milled Al samples tested exhibited around three timesigher tensile strengths compared with un-milled Al samples. This
ig. 5. Vickers micro-indentation and nano-indentation hardness of un-milled andilled pure Al and Al–2 wt% CNT extruded samples.
b) CNTs embedded between the re-welded particles after 3 h milling time.
is, in fact, due to grain size reduction and crystallographic defectsexpected from the milling process. It is well established that ballmilling can result in the development of a nanostructure (below100 nm) in aluminum powders. The exact crystal size depends onthe milling time, for example, Choi et al. [20], who milled thealuminum powders up to 48 h using an attritor mill at 550 rpm,observed a dependence of grain size on milling time which wentdown from 150 nm after 8 h to 48 nm after 48 h. In our case, wherewe have used a Ball-to-Powder weight ratio of 10:1 and a milling
Ball-milled samples containing CNTs exhibited high notch sen-sitivity and consistently fractured outside the gauge length, whichwas not the case for pure milled aluminum samples. Fig. 5 shows
Fig. 6. TEM image of the milled (MA = 6 h) pure Al matrix.
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Fig. 7. Indentation modulus for un-milled pure Al and milled pure and Al–2 wt%CNT samples.
Fig. 8. Vickers Hardness (GPa) at the investigated annealing temperatures and timesfor ball-milled pure aluminium extrudates.
Table 2Mean crystal size (nm) of powders, extrudates and annealed extrudates of Al–2 wt%CNT estimated using XRD.
Powders Extrudates◦
Annealed extrudates◦
Table 1Tensile strength and ductility of extrudates with different process histories, % difference fAl.
Note that the measured tensile strength for un-milled extruded Al is 130 MPa.
Fig. 9. XRD scans of extrudates with different process hi
(500 C, 30 min) (500 C, 10 h)
Al–2 wt% CNT (MA = 3 h) 87 – 93Al–2 wt% CNT (MA = 6 h) 45 56.5 72
nanoindentation and Vickers hardness results for extruded samplesof un-milled and milled Al, in addition to milled Al–2 wt% CNT com-posites. The hardness is seen to increase significantly (∼3 times)for extrusions of ball-milled powders, due to excessive strain hard-ening. XRD analysis showed that the mean grain size of powdersmilled for 6 h is 45 nm. For extrudates of pure aluminum, the meangrain size was found to be 48.4 nm and for extrudates of CNT–Al, itis 56.5 nm which is only slightly greater than the grain size of ball-milled powders, thus suggesting that extrusion at 500 ◦C retains thenanostructure. Fig. 6 shows a TEM image of a ball-milled extrudedaluminum sample confirming the nanostructure.
The addition of CNTs in Al–CNT ball-milled samples, however,does not seem to present any advantage in terms of hardness overmilled pure Al specimens. Fig. 7 shows that the CNTs also did notplay a role in stiffening the heavily strain-hardened Al matrix.
In an attempt to highlight the role of the CNTs and to reducethe notch-sensitivity of the samples containing CNTs, annealingexperiments were conducted. Initial annealing experiments were
performed on extruded pure aluminum that had undergone 6 hof ball milling. Vickers micro-hardness testing was used to assessthe properties of the annealed samples under different annealingconditions.
rom pure Al “with same process history” is provided in brackets for CNT reinforced
since this was the smallest sample used and thus more contributionfrom the mounting material resulted.
The results of the FESEM investigation of the tensile testing frac-tured surfaces of the 6 h ball-milled un-annealed, and the 6 and 3 hball-milled and annealed (500 ◦C, 10 h) samples are presented in
A.M.K. Esawi et al. / Materials Scienc
Fig. 8 summarizes the results, which show a clear decline inardness with increase in annealing temperature from 200 to00 ◦C. It is also clear that at a temperature of 400 ◦C, annealingor 8 h or more provides the best annealing results under the inves-igated processing conditions. It must be mentioned that althoughigher temperatures may result in more ductile samples, it is notithout a limit. This is especially true for Al–CNT composites, since
nterfacial reactions between the Al matrix and CNTs could result inarbide formation [12], which maybe excessive if amorphous car-on regions are present. For this reason, in this study the annealingemperatures investigated did not exceed 500 ◦C.
Bulk Al–CNT samples were produced and annealed at 400 ◦Cor 10 h prior to tensile testing. Although the samples exhibitedn increased ductility (compared to those which did not undergonnealing), and failed in the gage length, no enhancement intrength due to the presence of CNTs was observed. Additionally,nnealing at 500 ◦C was performed.
Table 1 shows the tensile strengths and ductilities (% elonga-ion to fracture) of the specimens tested. It is clear that only forhe 500 ◦C annealed specimen (which was ball-milled for 3 h andxtruded) there was an increase in strength of the CNT reinforcedaterial compared with pure aluminum of similar process history.
or all other specimens the addition of CNTs did not significantlyhange the properties of the un-reinforced counterpart, if, in factlightly decreased. When comparing our tensile strength valuesbtained for Al–2 wt% CNT composites with previous publishedork, it is clear that values reported here are significantly higher
han those previously published. For example, Esawi et al. [13],eported a tensile strength of 62 MPa for the same CNT content. Also,won et al. [6] who used spark plasma sintering and subsequentxtrusion to consolidate their 5 vol.% CNT–Al powders reported aensile strength of 194 MPa. In both studies, however, ball-millingas not used to disperse the CNTs. Our current improved properties
re mainly due to the better dispersion of CNTs provided by ball-illing in addition to the strain hardened powders contributing
s a strengthening mechanism to the final strength of the com-osites. Although they used a slightly lower mass fraction of CNTsi.e. 2 vol.% CNT), George et al. [16] reported a tensile strength of38 MPa (still significantly lower than our reported values). In thattudy ball-milling was applied, but only for 5 min, so cold workingf the powders may not be significant after this low milling time.ince it is well known that strain in ball-milled powders increasesith milling duration [24], our milling times of 3 and 6 h result inore significant strain hardening of the powders.XRD analysis of extrudates milled for 6 h, extruded at 500 ◦C and
nnealed for 10 h at 500 ◦C showed that a slight increase in meanrain size occurred compared to milled powders and bulk extrudate45, 56.5 and 72 nm, respectively). Reducing the milling time toh gave a mean grain size of 87 nm for the ball-milled powdersnd 93 nm for extruded and annealed extrudates, as summarizedn Table 2. This shows that the processing conditions used in theresent study resulted in only a slight increase in the crystal sizebtained through milling. The retention of the reduced crystal sizeesulted in the reported enhanced tensile strength of all the ball-illed samples.Fig. 9 shows the XRD scans of the specimens investigated,
lthough no clear peak for aluminum carbide is observed, this doesot rule out the formation of some carbides at the investigatednnealing temperature especially at 500 ◦C (the absence of car-ide peaks could be due to the limitation of the resolution of X-rayiffraction). It is interesting to note that with regards to whether
luminum carbide forms or not, previous published work reportedixed observations with some researchers reporting aluminum
arbide formation, for example, Deng et al. [11] who reportedluminum carbide formation at 656 ◦C and others not observingny [9,22]; which seems to suggest the strong dependence on
Engineering A 508 (2009) 167–173 171
the processing temperatures used as confirmed by Ci et al. [12]who investigated different annealing temperatures (450–950 ◦C)and reported that Al4C3 forms at the higher temperatures (start-ing 650 ◦C). It must be added that the unidentified peaks observedin the XRD scans are from the mounting material used to mount theextrudate samples. It appears more for the 500 ◦C annealed material
Fig. 10. Fracture surfaces of tensile specimens tested (a) 6 h ball-milled un-annealed,(b) 6 h ball-milled and annealed (500 ◦C for 10 h) and (c) 3 h ball-milled and annealed(500 ◦C for 10 h) showing individual CNTs dispersed in the Al matrix.
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ig. 11. (a) Deep etched surface of Al–CNT showing alignment of the CNTs in thextrusion direction, and the absence of void surrounding the nanotubes which haveeen noted in the fracture surfaces of tensile testing specimens. (b) Alignment ofNTs and lack of voids surrounding CNTs in a sample not subjected to tensile testing.
ig. 10. CNTs were observed to be well dispersed in the matrix. Whats also noticeable is the formation of voids around these nanotubesormed during the instability (necking) part of the deformation inensile testing. It is assumed that the nanotubes provide nucleationites for these voids that eventually contribute to fracture of theensile testing specimen. In addition, CNT pullout is also observedn Fig. 10, which indicates a poor interfacial bond between the CNTs
nd the aluminum matrix.
Fig. 11(a) and (b) shows the general alignment of the CNTs inhe direction of extrusion in deep-etched samples of Al–2 wt% CNTxtrudates. This may result in anisotropic properties of the final
ig. 12. FESEM image of fractured surface showing CNTs in which the individualNT layers are observed to be slipping.
Engineering A 508 (2009) 167–173
composite. Upon comparing the fracture surfaces of all samplesafter tensile testing (Fig. 10(a)–(c)) to those obtained by preferentialetching of the aluminum matrix (Fig. 11(a) and (b)), it is clear thatthere are no voids surrounding the CNTs in samples not subjectedto tensile testing. This confirms that void formation occurred onlyduring tensile testing.
Interestingly, inner tube slippage was also observed (Fig. 12)in the fractured surfaces, which may add to the ineffectivenessof the multi-wall nanotubes used in this study. Further work isbeing carried out to investigate the optimum ball-milling condi-tions (milling intensity, time and ball-to-powder ratio) in order toidentify the conditions that lead to the effective dispersion of theCNTs in Al, minimum damage to the CNTs, and minimum strain-hardening. The optimum annealing temperature and the effect offunctionalization on the Al–CNT interfacial bonding will also bestudied.
4. Conclusions
Clustering of CNTs when added to Al powder has previously beenreported as a major problem. Ball milling has been proven to be apromising technique for dispersing CNTs in the aluminum matrix.In this work, tensile strength enhancement of ∼21% was observedfor 2 wt% CNT reinforced aluminum processed by cold compactionand hot extrusion. Such an enhancement in mechanical proper-ties was only observed upon limiting the cold working effect ofthe ball-milling process (by reducing the milling time from 6 to3 h and by subjecting the samples to a post-processing annealingtreatment). Extrusion was also found to promote alignment of CNTsin the extrusion direction. Both XRD and TEM analysis showed thatalthough a slight growth in mean crystal size took place after extru-sion at 500 ◦C and annealing at 500 ◦C for 10 h, the nanostructure ofthe matrix was retained in the final products which contributedto the enhanced strength displayed by all samples compared toun-milled aluminum.
CNTs have been found to act as nucleation sites for void for-mation during tensile testing. In addition, both CNT pullout andCNT inner tube slippage were observed in fractured surfaces. Thepresent results show that the processing history of powders hasa significant effect on the final CNT–Al composite properties andbehaviour, a point receiving little attention so far in the literature.
Ongoing and future work includes: (1) the use of electricallyactivated processes to promote simultaneous consolidation andannealing of powders, (2) studying the effect of the aspect ratioof CNT, (3) studying the effect of using higher weight fraction CNTson the strengthening of the aluminum matrix, and (4) optimizationof ball-milling conditions.
Acknowledgments
The authors wish to acknowledge the financial support ofthe US-Egypt Joint Science and Technology fund (grant num-ber MAN11-011-007) and the National Science Foundation (Officeof International Science and Engineering) under grant number0710869, and the Yousef Jameel Science and Technology ResearchCenter (STRC) at the American University in Cairo. Thanks also toMs. Hanady Hussein and Mr. Rami Wasfi for technical assistance,and to Ms. Joan Kimbrough for assistance with X-ray diffrac-tion.
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