Fabrication and properties of dispersed carbon nanotube–aluminum composites

Materials Science and Engineering A 508 (2009) 167–173

Contents lists available at ScienceDirect

Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea

Fabrication and properties of dispersed carbon nanotube–aluminum composites A.M.K. Esawi a,∗ , K. Morsi b , A. Sayed a , A. Abdel Gawad a , P. Borah b a Department of Mechanical Engineering and the Yousef Jameel Science and Technology Research Center, The American University in Cairo, AUC Avenue, P.O. Box 74, New Cairo 11835, Egypt b Department of Mechanical Engineering, San Diego State University, 5500 Campanile Drive, San Diego, CA 92182, USA

a r t i c l e

i n f o

Article history: Received 8 October 2008 Received in revised form 15 December 2008 Accepted 1 January 2009 Keywords: Carbon nanotubes Aluminium matrix composites Extrusion Ball milling

a b s t r a c t 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 implications on the final properties of the composite. Both un-annealed and annealed Al–2 wt% CNT composites were investigated. It was found that, ball-milled and extruded (un-annealed) samples of Al–2 wt% CNT demonstrated high notch-sensitivity and consistently fractured outside the gauge length during tensile testing. In contrast, extruded samples annealed at 400 and at 500 ◦ C for 10 h prior to testing, exhibited more 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-milling conditions used were found to result in the creation of a nanostructure in all samples produced, as shown by XRD and TEM analysis. Such nanostructure was retained after prolonged exposures to temperatures up to 500 ◦ C. The tensile testing fracture surfaces showed uniform dispersion and alignment of the CNTs in the aluminum matrix but also showed CNTs acting as nucleation sites for void formation during tensile testing. This has contributed to the observation of CNT pull-out due to the poor bond between the CNTs and the matrix. © 2009 Elsevier B.V. All rights reserved.

1. Introduction The interest in carbon nanotubes (CNTs) as super reinforcements for metallic matrices has been growing considerably over the past few years, largely focusing on investigating their contribution to the enhancement of the mechanical performance of the final composite. As in conventional composites, the orientation of the CNTs, homogeneity of the composite, nanotube matrix adhesion, nanotube aspect ratio and the volume fraction of nanotubes are expected to have significant influences on the properties of the nanocomposite. Controlling such factors to obtain an exceptional composite is very challenging. A number of reviews have been published on the subject, including ones by Harris [1], CarrenoMorelli [2] and Curtin and Sheldon [3]. Aluminum (Al) has benefited from intense research efforts in this area, where CNTs are hoped to provide substantial improvements to its properties [4–22]. One of the major obstacles to the effective use of carbon nanotubes as reinforcements in metal matrix composites is their agglomeration and poor distribution/dispersion within the metallic matrix.

∗ Corresponding author. Tel.: +20 2 2615 3102. E-mail address: a [email protected] (A.M.K. Esawi). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.01.002

A number of research groups, including the current authors, investigated 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 dispersed CNTs proved the process to be promising [9,14,16]. However, concerns about the possible damage and/or amorphization of the CNTs under the harsh milling conditions have been raised [9,22] so optimization of milling conditions is necessary. CNT contents up to 10 wt% have been investigated. Enhancements in mechanical properties due to CNT addition have been reported [6,9,21]. The optimum CNT content at which maximum enhancement was achieved 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 presence of unreacted and unmelted CNTs after exposure to the high temperatures associated with the plasma spraying process they used, most researchers have been cautious about processing CNT-composites at high temperatures so as to avoid CNT damage and possible adverse chemical interfacial reactions with the matrix. It has been argued, however, that the formation of Al4 C3 phase should not be viewed as detrimental since it can help to enhance the Al–CNT bond and can lock the nanotubes in place and hence contribute to the enhancement of the mechanical properties of the composites [6,12].

168

A.M.K. Esawi et al. / Materials Science and Engineering A 508 (2009) 167–173

In another study [18], the formation of an ultra thin ␤-SiC layer in an Al–Si alloy containing 10 wt% CNT synthesized by thermal spraying was observed to enhance the limited wettability and the interfacial bond between the CNTs and the aluminum matrix. One way to produce composites with uniform distribution of CNTs within the metallic matrix is by first dispersing them in metal powders then subsequently consolidating these Al–CNT composite powders. The ball-milling process, although effective, normally results in strain hardening of the powders, and this can have implications on the final properties of the bulk composite. Therefore the consolidation of heavily worked powders and their subsequent properties would be important from both a practical as well as a scientific standpoint. The present authors believe that results for CNT reinforced aluminum or even other metals should be viewed in light of the condition of the powders used to form the bulk samples (e.g. degree of strain hardening) and their subsequent process history (e.g. hot extrusion, annealing), a point that has not been given too much attention so far in the literature. Potential problems may also arise in terms of matrix/CNT interfacial reactions and CNT damage during high temperature processing or metal working. In previous studies by the authors, dispersion of 2 and 5 wt% CNTs in aluminum powders was achieved, and the effect of milling time on the morphological evolution of the Al–CNT composite powders was investigated [14,15]. This paper discusses the consolidation of ball-milled Al–CNT composite powders using hot extrusion and the evaluation of the mechanical response of the composites (both annealed and un-annealed) using tensile and hardness testing. 2. Experimental procedure 30 grams of Al (99.7% pure, −200 mesh, Aluminum Powder Company Ltd., UK) and 2 wt% multi-wall carbon nanotubes (average dimensions: 140 ± 30 nm outer diameter, 4–8 nm internal diameter, and 3–4 ␮m in length, MER corporation, USA) were placed in 250 ml stainless steel mixing jars together with 75 stainless steel milling balls (10 mm diameter); giving a ball-to-powder weight ratio of 10:1. The jars were filled with argon and were agitated using a planetary ball mill at 200 rpm for 3 and 6 h. 2 ml of methanol were added as a process control agent (PCA) in order to minimize cold welding of the Al particles and also to prevent powders sticking to the balls and the jar walls. In the case of pure Al, 2.5 ml of methanol were added. Approximately 26 g of the ball-milled powder mix were compacted in a 20 mm diameter compaction die at 475 MPa. Hot extrusion of the homogenized compact was conducted at 500 ◦ C using an extrusion ratio of 4:1, to produce 10 mm diameter extrudates. Tensile test samples were machined and super-finished out of the extrudates, with dimensions as shown in Fig. 1. The mechanical response of the processed materials was also characterized using both a Vickers Micro-hardness tester (Load = 300 g and Dwell time = 15 s) and Nanoindentation. Samples from the resulting extrudates were sectioned perpendicular to the extrusion direction and mounted using bakelite due to its lower

Fig. 1. Tensile test specimen dimensions.

elasticity compared to epoxy resins. The samples were ground and polished to 1 ␮m. Nanoindentation tests were performed on a Nanoindenter XP (MTS systems Co., Oak Ridge, TN, USA) with a diamond pyramid-shaped Berkovich-type indenter tip. The nanoindentation tests were conducted according to the Oliver & Pharr approach [23]. Continuous Stiffness method (CSM) was used. The maximum displacement of the tip into the surface was set to a high value (10 ␮m) in order to get the maximum possible indentation depth. Distance between adjacent indentations was set to 100 ␮m so as to avoid the effect of residual stresses around neighboring indentations. The thermal drift was kept below 0.05 nm/s. The reported hardness values are averages of at least ten indentations in the case of Vickers hardness testing and twenty indentations (5 × 4 matrix) for the nanoindentation experiments. For microstructural examination, the specimens were cut, ground and polished to 1 ␮m finish. Etching was performed in Keller’s solution, and the microstructure was characterized using a field emission scanning electron microscope FESEM (LEO supra 55). X-ray diffraction (XRD) (using Cu K␣, Panalytical Xpert Pro diffractometer) was also used for phase analysis and crystal size determination using the Scherer equation [24], and excluding instrumental effects on broadening. 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 and subsequent electropolishing techniques. 3. Results and discussion Fig. 2 shows the XRD scans for Al–CNT powders milled for 3 and 6 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 disappears after milling for 6 h. It is interesting to note that in previous published work that applied ball milling for 5 min, a clear peak at 26◦ was not observed, despite the TEM observation of the CNTs in the sample [16], albeit at ∼2 vol.% CNT which maybe at the limit of XRD resolution. In our case also, in spite of clear observation of the CNTs 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 powders ball-milled for 3 h. This could be due to a number of possible reasons; including the small amount of CNTs used, the well dispersed CNTs within the matrix, unfavorable strain/CNTs effect that reduces the CNT peak intensity, and amorphization of CNTs. Furthermore, the figure does not show any clear evidence of carbide phase formation under the current ball-milling conditions. Fig. 3 shows a clear distinction between the morphology of the Al–CNT powders milled for 3 and 6 h, where 6 h milling results in powders with bi-modal particle size distribution. In both cases of the 3 and 6 h milling, methanol was added as a process control agent

Fig. 2. XRD scans of Al–2 wt% CNT powders ball-milled for 3 and 6 h.

A.M.K. Esawi et al. / Materials Science and Engineering A 508 (2009) 167–173

169

Fig. 3. SEM micrographs of Al–2 wt% CNT powders ball-milled for (a) 3 h and (b) 6 h, respectively.

Fig. 4. (a) CNTs on the surface of aluminum powder after 0.5 h milling time and (b) CNTs embedded between the re-welded particles after 3 h milling time.

to limit the Al particle welding which tends to increase the particle size, as noted in our previous publication [14]. In addition to particle welding, the other competing process during ball milling is strain hardening of the powders, which leads to a decrease in ductility and eventual fracturing of the Al particles and thus a decrease in particle size. As the milling time increased from 3 to 6 h, strain/work-hardening of the powders is increased favoring more particles fracturing than re-welding, and leading to a refinement in the particle size. As reported in an earlier study by the authors [14], the ballmilling process was proven to be effective in dispersing the CNTs. Fig. 4(a) shows well-dispersed CNTs on the surface of ball-milled aluminum powders after 0.5 h of milling and Fig. 4(b) shows the CNTs embedded between the rewelded particles after 3 h of milling. All ball-milled Al samples tested exhibited around three times higher tensile strengths compared with un-milled Al samples. This

Fig. 5. Vickers micro-indentation and nano-indentation hardness of un-milled and milled pure Al and Al–2 wt% CNT extruded samples.

is, in fact, due to grain size reduction and crystallographic defects expected from the milling process. It is well established that ball milling can result in the development of a nanostructure (below 100 nm) in aluminum powders. The exact crystal size depends on the milling time, for example, Choi et al. [20], who milled the aluminum powders up to 48 h using an attritor mill at 550 rpm, observed a dependence of grain size on milling time which went down from 150 nm after 8 h to 48 nm after 48 h. In our case, where we have used a Ball-to-Powder weight ratio of 10:1 and a milling time of 6 h, the mean grain size, as found by XRD analysis, was 48.4 nm (for pure aluminium milled for 6 h and extruded at 500 ◦ C). Ball-milled samples containing CNTs exhibited high notch sensitivity and consistently fractured outside the gauge length, which was not the case for pure milled aluminum samples. Fig. 5 shows

Fig. 6. TEM image of the milled (MA = 6 h) pure Al matrix.

170

A.M.K. Esawi et al. / Materials Science and Engineering A 508 (2009) 167–173 Table 2 Mean crystal size (nm) of powders, extrudates and annealed extrudates of Al–2 wt% CNT estimated using XRD.

Al–2 wt% CNT (MA = 3 h) Al–2 wt% CNT (MA = 6 h)

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 times for ball-milled pure aluminium extrudates.

Powders

Extrudates (500 ◦ C, 30 min)

Annealed extrudates (500 ◦ C, 10 h)

87 45

– 56.5

93 72

nanoindentation and Vickers hardness results for extruded samples of un-milled and milled Al, in addition to milled Al–2 wt% CNT composites. The hardness is seen to increase significantly (∼3 times) for extrusions of ball-milled powders, due to excessive strain hardening. XRD analysis showed that the mean grain size of powders milled for 6 h is 45 nm. For extrudates of pure aluminum, the mean grain size was found to be 48.4 nm and for extrudates of CNT–Al, it is 56.5 nm which is only slightly greater than the grain size of ballmilled powders, thus suggesting that extrusion at 500 ◦ C retains the nanostructure. Fig. 6 shows a TEM image of a ball-milled extruded aluminum 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 over milled pure Al specimens. Fig. 7 shows that the CNTs also did not play a role in stiffening the heavily strain-hardened Al matrix. In an attempt to highlight the role of the CNTs and to reduce the notch-sensitivity of the samples containing CNTs, annealing experiments were conducted. Initial annealing experiments were performed on extruded pure aluminum that had undergone 6 h of ball milling. Vickers micro-hardness testing was used to assess the properties of the annealed samples under different annealing conditions.

Table 1 Tensile strength and ductility of extrudates with different process histories, % difference from pure Al “with same process history” is provided in brackets for CNT reinforced Al. Pre-extrusion condition

Aluminum

Aluminum + 2 wt% CNT Post-extrusion anneal (500 ◦ C, 10 h)

Post-extrusion anneal (400 ◦ C, 10 h)

Ball-milled (3 h) Ball-milled (6 h)

Post-extrusion anneal (400 ◦ C, 10 h)

Post-extrusion anneal (500 ◦ C, 10 h)

TS

EL%

TS

EL%

TS

EL%

TS

EL%

– 377.4

– 5.9

284.5 348.5

8.6 8.4

– 365.5 (−3.15%)

– 6.9

345 (+21.3%) 348 (−0.15%)

5.7 7.9

Note that the measured tensile strength for un-milled extruded Al is 130 MPa.

Fig. 9. XRD scans of extrudates with different process histories, showing predominantly aluminum peaks.

A.M.K. Esawi et al. / Materials Science and Engineering A 508 (2009) 167–173

Fig. 8 summarizes the results, which show a clear decline in hardness with increase in annealing temperature from 200 to 400 ◦ C. It is also clear that at a temperature of 400 ◦ C, annealing for 8 h or more provides the best annealing results under the investigated processing conditions. It must be mentioned that although higher temperatures may result in more ductile samples, it is not without a limit. This is especially true for Al–CNT composites, since interfacial reactions between the Al matrix and CNTs could result in carbide formation [12], which maybe excessive if amorphous carbon regions are present. For this reason, in this study the annealing temperatures investigated did not exceed 500 ◦ C. Bulk Al–CNT samples were produced and annealed at 400 ◦ C for 10 h prior to tensile testing. Although the samples exhibited an increased ductility (compared to those which did not undergo annealing), and failed in the gage length, no enhancement in strength due to the presence of CNTs was observed. Additionally, annealing at 500 ◦ C was performed. Table 1 shows the tensile strengths and ductilities (% elongation to fracture) of the specimens tested. It is clear that only for the 500 ◦ C annealed specimen (which was ball-milled for 3 h and extruded) there was an increase in strength of the CNT reinforced material compared with pure aluminum of similar process history. For all other specimens the addition of CNTs did not significantly change the properties of the un-reinforced counterpart, if, in fact slightly decreased. When comparing our tensile strength values obtained for Al–2 wt% CNT composites with previous published work, it is clear that values reported here are significantly higher than those previously published. For example, Esawi et al. [13], reported a tensile strength of 62 MPa for the same CNT content. Also, Kwon et al. [6] who used spark plasma sintering and subsequent extrusion to consolidate their 5 vol.% CNT–Al powders reported a tensile strength of 194 MPa. In both studies, however, ball-milling was not used to disperse the CNTs. Our current improved properties are mainly due to the better dispersion of CNTs provided by ballmilling in addition to the strain hardened powders contributing as a strengthening mechanism to the final strength of the composites. Although they used a slightly lower mass fraction of CNTs (i.e. 2 vol.% CNT), George et al. [16] reported a tensile strength of 138 MPa (still significantly lower than our reported values). In that study ball-milling was applied, but only for 5 min, so cold working of the powders may not be significant after this low milling time. Since it is well known that strain in ball-milled powders increases with milling duration [24], our milling times of 3 and 6 h result in more significant strain hardening of the powders. XRD analysis of extrudates milled for 6 h, extruded at 500 ◦ C and annealed for 10 h at 500 ◦ C showed that a slight increase in mean grain size occurred compared to milled powders and bulk extrudate (45, 56.5 and 72 nm, respectively). Reducing the milling time to 3 h gave a mean grain size of 87 nm for the ball-milled powders and 93 nm for extruded and annealed extrudates, as summarized in Table 2. This shows that the processing conditions used in the present study resulted in only a slight increase in the crystal size obtained through milling. The retention of the reduced crystal size resulted in the reported enhanced tensile strength of all the ballmilled samples. Fig. 9 shows the XRD scans of the specimens investigated, although no clear peak for aluminum carbide is observed, this does not rule out the formation of some carbides at the investigated annealing temperature especially at 500 ◦ C (the absence of carbide peaks could be due to the limitation of the resolution of X-ray diffraction). It is interesting to note that with regards to whether aluminum carbide forms or not, previous published work reported mixed observations with some researchers reporting aluminum carbide formation, for example, Deng et al. [11] who reported aluminum carbide formation at 656 ◦ C and others not observing any [9,22]; which seems to suggest the strong dependence on

171

the processing temperatures used as confirmed by Ci et al. [12] who investigated different annealing temperatures (450–950 ◦ C) and reported that Al4 C3 forms at the higher temperatures (starting 650 ◦ C). It must be added that the unidentified peaks observed in the XRD scans are from the mounting material used to mount the extrudate samples. It appears more for the 500 ◦ C annealed material since this was the smallest sample used and thus more contribution from the mounting material resulted. The results of the FESEM investigation of the tensile testing fractured surfaces of the 6 h ball-milled un-annealed, and the 6 and 3 h ball-milled and annealed (500 ◦ C, 10 h) samples are presented in

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.

172

A.M.K. Esawi et al. / Materials Science and Engineering A 508 (2009) 167–173

composite. Upon comparing the fracture surfaces of all samples after tensile testing (Fig. 10(a)–(c)) to those obtained by preferential etching of the aluminum matrix (Fig. 11(a) and (b)), it is clear that there are no voids surrounding the CNTs in samples not subjected to tensile testing. This confirms that void formation occurred only during tensile testing. Interestingly, inner tube slippage was also observed (Fig. 12) in the fractured surfaces, which may add to the ineffectiveness of the multi-wall nanotubes used in this study. Further work is being carried out to investigate the optimum ball-milling conditions (milling intensity, time and ball-to-powder ratio) in order to identify the conditions that lead to the effective dispersion of the CNTs in Al, minimum damage to the CNTs, and minimum strainhardening. The optimum annealing temperature and the effect of functionalization on the Al–CNT interfacial bonding will also be studied. 4. Conclusions

Fig. 11. (a) Deep etched surface of Al–CNT showing alignment of the CNTs in the extrusion direction, and the absence of void surrounding the nanotubes which have been noted in the fracture surfaces of tensile testing specimens. (b) Alignment of CNTs and lack of voids surrounding CNTs in a sample not subjected to tensile testing.

Fig. 10. CNTs were observed to be well dispersed in the matrix. What is also noticeable is the formation of voids around these nanotubes formed during the instability (necking) part of the deformation in tensile testing. It is assumed that the nanotubes provide nucleation sites for these voids that eventually contribute to fracture of the tensile testing specimen. In addition, CNT pullout is also observed in Fig. 10, which indicates a poor interfacial bond between the CNTs and the aluminum matrix. Fig. 11(a) and (b) shows the general alignment of the CNTs in the direction of extrusion in deep-etched samples of Al–2 wt% CNT extrudates. This may result in anisotropic properties of the final

Clustering of CNTs when added to Al powder has previously been reported as a major problem. Ball milling has been proven to be a promising technique for dispersing CNTs in the aluminum matrix. In this work, tensile strength enhancement of ∼21% was observed for 2 wt% CNT reinforced aluminum processed by cold compaction and hot extrusion. Such an enhancement in mechanical properties was only observed upon limiting the cold working effect of the ball-milling process (by reducing the milling time from 6 to 3 h and by subjecting the samples to a post-processing annealing treatment). Extrusion was also found to promote alignment of CNTs in the extrusion direction. Both XRD and TEM analysis showed that although a slight growth in mean crystal size took place after extrusion at 500 ◦ C and annealing at 500 ◦ C for 10 h, the nanostructure of the matrix was retained in the final products which contributed to the enhanced strength displayed by all samples compared to un-milled aluminum. CNTs have been found to act as nucleation sites for void formation during tensile testing. In addition, both CNT pullout and CNT inner tube slippage were observed in fractured surfaces. The present results show that the processing history of powders has a significant effect on the final CNT–Al composite properties and behaviour, a point receiving little attention so far in the literature. Ongoing and future work includes: (1) the use of electrically activated processes to promote simultaneous consolidation and annealing of powders, (2) studying the effect of the aspect ratio of CNT, (3) studying the effect of using higher weight fraction CNTs on the strengthening of the aluminum matrix, and (4) optimization of ball-milling conditions. Acknowledgments The authors wish to acknowledge the financial support of the US-Egypt Joint Science and Technology fund (grant number MAN11-011-007) and the National Science Foundation (Office of International Science and Engineering) under grant number 0710869, and the Yousef Jameel Science and Technology Research Center (STRC) at the American University in Cairo. Thanks also to Ms. Hanady Hussein and Mr. Rami Wasfi for technical assistance, and to Ms. Joan Kimbrough for assistance with X-ray diffraction. References

Fig. 12. FESEM image of fractured surface showing CNTs in which the individual CNT layers are observed to be slipping.

[1] P.J.F. Harris, Int. Mater. Rev. 49 (1 (February)) (2004) 31–43. [2] E. Carreno-Morelli, Carbon Nanotube–Metal Matrix Composites. The Dekker Encyclopedia of Nanoscience and Nanotechnology, Taylor and Francis, New York, 2006, pp. 1–9. [3] W.A. Curtin, B.W. Sheldon, Mater. Today 7 (11 (November)) (2004) 44–49.

A.M.K. Esawi et al. / Materials Science and Engineering A 508 (2009) 167–173 [4] [5] [6] [7] [8] [9]

[10] [11] [12] [13]

C.L. Xu, B.Q. Wei, R.Z. Ma, J. Liang, X.K. Ma, D.H. Wu, Carbon 37 (1999) 855–858. R. Zhong, H. Cong, P. Hou, Carbon 41 (2003) CO1-851. H. Kwon, M. Estili, K. Takagi, T. Miyazaki, A. Kawasaki, Carbon 47 (2009) 570–577. T. Kuzumaki, K. Miyazawa, H. Ichinose, K.J. Ito, Mater. Res. 13 (1998) 2445– 2449. S.-M. Zhou, X.-B. Zhang, Z.-P. Ding, C.-V. Min, G.-L. Xu, W.-M. Zhu, Compos. Part A: Appl. Sci. Manuf. 38 (2 (February)) (2007) 301–306. R. Perez-Bustamante, I. Estrada-Guel, W. Antunez-Flores, M. Miki-Yoshida, P.J. Ferreira, R. Martinez-Sanchez, J. Alloys Compd. 450 (1–2 (14 February)) (2008) 323–326. C.F. Deng, X.X. Zhang, D. Wang, Q. Lin, A. Li, Mater. Lett. 61 (2007) 1725–1728. C.F. Deng, D.Z. Wang, X.X. Zhang, A.B. Li, Mater. Sci. Eng. A 444 (1–2 (25 January)) (2007) 138–145. L. Ci, R. Zhenyu, N.Y. Jin-Phillipp, M. Rühle, Acta Mater. 54 (20 (December)) (2006) 5367–5375. A.M.K. Esawi, M.A. El Borady, Compos. Sci. Technol. 68 (2 (February)) (2008) 486–492.

173

[14] A.M.K. Esawi, K. Morsi, Compos. Part A: Appl. Sci. Manuf. 38 (2 (February)) (2007) 646–650. [15] K. Morsi, A. Esawi, J. Mater. Sci. 42 (2007) 4954–4959. [16] R. George, K.T. Kashyap, R. Rahul, S. Yamdagni, Scripta Mater. 53 (2005) 1159–1163. [17] T. Laha, A. Agarwal, Mater. Sci. Eng. A 480 (2008) 323–332. [18] T. Laha, S. Kuchibhatla, S. Seal, W. Li, A. Agarwal, Acta Mater. 55 (2007) 1059–1066. [19] T. Laha, A. Agarwal, T. McKechnie, S. Seal, Mater. Sci. Eng. A 381 (2004) 249–258. [20] H.J. Choi, S.W. Lee, J.S. Park, D.H. Bae, Scripta Mater. 59 (2008) 1123–1126. [21] H.J. Choi, G.B. Kwon, G.Y. Lee, D.H. Bae, Scripta Mater. 59 (2008) 360–363. [22] C. Edtmaier, E. Wallnoefer, A. Koeck, Proceedings of Powder Metallurgy World Congress PM2004, Vienna, Austria, October 17–21, 2004; Danninger, H. Ratzi, R. (Eds.), European Powder Metallurgy Association: Shrewsbury, UK, 2004; Vol. 1, 413–418. [23] W.C. Oliver, G.M. Pharr, Mater. J. Res. 19 (1) (2004) 3–20. [24] C. Suryanarayana, Prog. Mater. Sci. 46 (1–2 (January)) (2001) 1–184.