Strong and ductile carbon nanotube/aluminum bulk nanolaminated composites with two-dimensional alignment of carbon nanotubes Lin Jiang, Zhiqiang Li, ⇑ Genlian Fan, Linlin Cao and Di Zhang ⇑ State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai 200240, PR China Received 22 September 2011; accepted 18 November 2011 Available online 26 November 2011 In order to combine high tensile strength and ductile behavior, carbon nanotube (CNT)/Al nanolaminated composites with alter- nating layers of Al (400 nm) and CNTs (50 nm) were fabricated by flake powder metallurgy. Compared with conventional homogeneous nanocomposites composed of the same constituents, the final bulk products with high level ordered nanolaminates exhibited both greatly improved tensile strength of 375 MPa and plasticity of 12%, mainly because they enabled enhanced disloca- tion storage capability and two-dimensional alignment of CNTs. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Aluminum; Carbon nanotubes; Powder consolidation; Layered structures; Ductility With a Young’s modulus higher than 1 10 12 Pa and tensile strength about 50 times that of steel, quasi- one-dimensional carbon nanotubes (CNTs) can still undergo plastic deformation to failure of 20–30% [1–3]. Thus, incorporating CNTs into light metals and alloys, such as Al, was supposed to make a new class of engineer- ing nanocomposites with both high strength and good ductility, which is suitable for the automotive or aero- space industries [4]. However, for lack of effective fabrica- tions, the as-prepared CNT/Al nanocomposites usually failed to fully realize the strengthening efficiency of CNTs and achieve satisfying ductility, which is an obvious draw- back to their practical applications [3,5]. Several methods, such as high energy ball milling (HEBM) [6], an in situ approach [7] and nanoscale dispersion (NSD) [8], have been exploited to supply homogeneous CNT/Al compos- ite powders for the powder metallurgy (PM) route. In these approaches, most emphasis has been placed on breaking up agglomeration of CNTs to homogeneously disperse them in three-dimensional (3-D) spherical Al powders [6–8], whereas little attention has been paid to the control of shape and stacking modes of composite powders, so that CNT distribution and composite structure have been left to evolve randomly. Such a homo- geneous structure with 3-D randomly distributed CNTs, however, is unfavorable to enable longitudinal properties of one-dimensional (1-D) CNTs and improve dislocation storage capability, which is the prerequisite for compati- ble plastic strains of metals. Moreover, structural defects were caused either by the severe mechanical damage to CNTs in HEBM [9] or by the unstable growth of CNTs during in situ reactions. Therefore, the ductility of these as-fabricated nanocomposites was disappointing. For example, HEBM CNT/Al nanocomposite shows strength of 345 MPa but only 4.7% for elongation [6], while the in situ CNT/Al nanocomposite shows strength of 398 MPa but only 2% for elongation [7]. Thus, a strategy to tailor the structure of CNT/Al nanocomposites and thus endowing such materials with both high strength and good ductility is in great demand. Recently, nanolaminate architecture has become a prevailing model to develop a new generation of high strength metal/metal or metal/ceramic bimaterials with considerable ductility [10,11], due to their remarkable ability of energy-absorbing and dislocation storage [11,12]. Meanwhile, there is an increasing effort to con- trol two-dimensional (2-D) distribution of CNTs and thus create a CNT/metal laminate structure. Zhu et al. [13] fabricated CNT/Cu laminated composites by re- peated cold rolling of CNT films sandwiched between Cu thin foils and annealing, and Kim et al. [14] obtained a similar laminate microstructure by selective dip- coating of CNTs and Cu. In these attempts, both strength and ductility enhancement were realized by the laminate structure with 2-D distribution of CNTs. Nevertheless, these methods are restricted to the 1359-6462/$ - see front matter Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2011.11.023 ⇑ Corresponding authors. Tel.: +86 21 3420 2584; fax: +86 21 3420 3913; e-mail addresses: [email protected]; [email protected]Available online at www.sciencedirect.com Scripta Materialia 66 (2012) 331–334 www.elsevier.com/locate/scriptamat
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Available online at www.sciencedirect.com
Scripta Materialia 66 (2012) 331–334
www.elsevier.com/locate/scriptamat
Strong and ductile carbon nanotube/aluminum bulk nanolaminatedcomposites with two-dimensional alignment of carbon nanotubes
Lin Jiang, Zhiqiang Li,⇑ Genlian Fan, Linlin Cao and Di Zhang⇑
State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai 200240, PR China
Received 22 September 2011; accepted 18 November 2011Available online 26 November 2011
In order to combine high tensile strength and ductile behavior, carbon nanotube (CNT)/Al nanolaminated composites with alter-nating layers of Al (�400 nm) and CNTs (�50 nm) were fabricated by flake powder metallurgy. Compared with conventionalhomogeneous nanocomposites composed of the same constituents, the final bulk products with high level ordered nanolaminatesexhibited both greatly improved tensile strength of 375 MPa and plasticity of 12%, mainly because they enabled enhanced disloca-tion storage capability and two-dimensional alignment of CNTs.� 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
With a Young’s modulus higher than 1 � 1012 Paand tensile strength about 50 times that of steel, quasi-one-dimensional carbon nanotubes (CNTs) can stillundergo plastic deformation to failure of 20–30% [1–3].Thus, incorporating CNTs into light metals and alloys,such as Al, was supposed to make a new class of engineer-ing nanocomposites with both high strength and goodductility, which is suitable for the automotive or aero-space industries [4]. However, for lack of effective fabrica-tions, the as-prepared CNT/Al nanocomposites usuallyfailed to fully realize the strengthening efficiency of CNTsand achieve satisfying ductility, which is an obvious draw-back to their practical applications [3,5]. Several methods,such as high energy ball milling (HEBM) [6], an in situapproach [7] and nanoscale dispersion (NSD) [8], havebeen exploited to supply homogeneous CNT/Al compos-ite powders for the powder metallurgy (PM) route. Inthese approaches, most emphasis has been placed onbreaking up agglomeration of CNTs to homogeneouslydisperse them in three-dimensional (3-D) spherical Alpowders [6–8], whereas little attention has been paid tothe control of shape and stacking modes of compositepowders, so that CNT distribution and compositestructure have been left to evolve randomly. Such a homo-geneous structure with 3-D randomly distributed CNTs,however, is unfavorable to enable longitudinal properties
1359-6462/$ - see front matter � 2011 Acta Materialia Inc. Published by Eldoi:10.1016/j.scriptamat.2011.11.023
of one-dimensional (1-D) CNTs and improve dislocationstorage capability, which is the prerequisite for compati-ble plastic strains of metals. Moreover, structural defectswere caused either by the severe mechanical damage toCNTs in HEBM [9] or by the unstable growth of CNTsduring in situ reactions. Therefore, the ductility of theseas-fabricated nanocomposites was disappointing. Forexample, HEBM CNT/Al nanocomposite shows strengthof 345 MPa but only 4.7% for elongation [6], while thein situ CNT/Al nanocomposite shows strength of398 MPa but only 2% for elongation [7]. Thus, a strategyto tailor the structure of CNT/Al nanocomposites andthus endowing such materials with both high strengthand good ductility is in great demand.
Recently, nanolaminate architecture has become aprevailing model to develop a new generation of highstrength metal/metal or metal/ceramic bimaterials withconsiderable ductility [10,11], due to their remarkableability of energy-absorbing and dislocation storage[11,12]. Meanwhile, there is an increasing effort to con-trol two-dimensional (2-D) distribution of CNTs andthus create a CNT/metal laminate structure. Zhu et al.[13] fabricated CNT/Cu laminated composites by re-peated cold rolling of CNT films sandwiched betweenCu thin foils and annealing, and Kim et al. [14] obtaineda similar laminate microstructure by selective dip-coating of CNTs and Cu. In these attempts, bothstrength and ductility enhancement were realized bythe laminate structure with 2-D distribution of CNTs.Nevertheless, these methods are restricted to the
332 L. Jiang et al. / Scripta Materialia 66 (2012) 331–334
fabrication of miniature sized sheets and, more impor-tantly, they can hardly reduce the laminate thicknessto submicron or nanoscale.
In the current study, CNT/Al bulk nanolaminatedcomposites were fabricated by a flake powder metal-lurgy (flake PM) route. Through morphology transfor-mation from 3-D sphere to 2-D nanoflake and surfacemodification with poly(vinyl alcohol) (PVA), the Alnanoflakes can uniformly anchor 1-D CNTs onto theirsurface. The as-obtained composite powders could beforced to stack in an orderly way into nanolaminateswith 2-D alignment of CNTs through powder compact-ing, followed by hot extrusion. The high level orderingnanolaminate architecture and the structure integrityof CNTs were achieved in the final products. Thus,excellent mechanical performance of the final productswas promised.
In a typical flake PM, three steps were involved, asshown in Figure 1b.
(i) Nanoflake powder preparation. The near-spheri-cal powder (10 lm in diameter, and 99.5 wt.% inpurity) with 1 wt.% stearic acid can quickly reacha flake thickness of several hundred nanometersby ball-milling in an attritor at 423 rpm at roomtemperature for only 1–2 h.
(ii) Adsorption of CNTs. Mutil-wall CNTs (30–50 nm in diameter, �3 lm in length, functional-ized with carboxyl groups (–COOH)) were dis-persed into water by ultrasonicating for 2 h andthe as-prepared Al nanoflakes were surface mod-ified by PVA with 1700–1800 repeat units. ThePVA-coated nanoflakes were added into waterto make powder slurry, and then the CNT aque-ous dispersion was added in drop by drop. Themixed slurry was mechanically stirred until itscolor changed from black to transparent, then fil-tered and rinsed with deionized water to get theCNT/Al nanocomposite powders.
(iii) Nanoflake powder alignment and consolidation.Before compacting, the CNT/Al nanocompositepowders were heated in flowing Ar atmosphereat 500 �C for 2 h to remove PVA from the nano-composite powder. Compacting was used to alignthe nanoflakes into column (U40 mm � 30 mm)under 500 MPa pressure. Sintering in flowingAr atmosphere at 550 �C for 2 h and hot extru-sion at 440 �C with an extrusion ratio of 20:1 ata ram speed of 0.5 mm min�1 were conductedto consolidate the nanoflakes. For comparison,we also prepared the CNT/Al nanocomposites
Figure 1. Fabrication procedures for CNT/Al nanocom
by conventional PM, in which slurry blending isfirst used to disperse CNT in spherical Al pow-ders with a 3-D spherical morphology (Fig. 2a)and then HEBM (initial ball-to-powder weightratio of 20:1, 300 rpm/6 h) was used to co-millCNTs and Al powders, as illustrated in Figure1a. The consolidation parameters of conventionPM were the same as that of flake PM.
The structure parameters of samples were character-ized by field emission scanning electron microscopy(FESEM) using a LEO Supra 55 FESEM and transmis-sion electron microscopy (TEM) in a Philips CM200microscope operated at 200 kV. To evaluate the tensilestrength, specimens were machined from the extrudedrods with the tensile axis parallel to the extrusion direc-tion. The gauge length of the specimens was 25 mm, andthe diameter was 5 mm. The tensile strength wasmeasured by a universal testing machine at an initialstrain rate of 5 � 10�4 s�1 at room temperature (AUTO-GRAPH AG-I 50 KN, Shimadzu Co. Ltd., Japan). Thedislocation density of samples was analyzed by X-raydiffraction (XRD, Rigaku D/max-2550/PC) with a CuKa radiation source and calculated as a function ofthe contribution of micro-strain and crystallite size usingthe following relationships [15]:
qe ¼ ð3Ke2=D2b2Þ1=2 ð1Þwhere e is micro-strain, K is the factor related toGaussian strain distribution, D is crystallite size and bis the Burgers vector.
As seen in Figure 2b, the as-prepared nanoflakes havea 2-D planar morphology with an average diameter of45 lm and thickness of 300–500 nm, thus assuring alarge aspect ratio (diameter to thickness), which is help-ful for the alignment of flake powders under compact-ing. After PVA modification, these 2-D nanoflakes canuniformly adsorb CNTs on their surface (Fig. 2c) dueto the formation of hydrogen bonding between the –OH group of the PVA membrane and the –COOHgroup of the functionized CNTs [16]. After compacting,the CNT/Al powder compact exhibits a structure with astrikingly strong nanoflake alignment (Fig. 2d) and 2-Ddistributed CNTs are confined by the nanoflakes (insetof Fig. 2d). After extrusion, the dense nanocompositesfabricated by flake PM show that all the Al plateletswith aligned CNTs are organized with their facesparallel to extrusion direction (Fig. 3a). The extrudedmultilayer structure with alternating Al (�400 nm) andCNT (�50 nm) layers can be clearly observed in Figure3b, while, as obviously seen in Figure 3c and d, CNTs
posites: (a) conventional PM and (b) flake PM.
Figure 2. FESEM of: (a) spherical Al powders; (b) Al nanoflakes, insetshows the thickness of the nanoflakes; (c) nanoflake with uniformlyadsorbed CNT on its surface; (d) CNT/Al nanolaminates aftercompacting, inset shows the CNTs confined by nanoflakes.
L. Jiang et al. / Scripta Materialia 66 (2012) 331–334 333
exhibit disorientation and form a 3-D random distribu-tion in the conventional PM nanocomposites.
The resulting tensile properties of CNT/Al nanocom-posites fabricated by conventional PM and flake PM areshown in Figure 4a. The flake PM CNT/Al nanocom-posite shows a plasticity of 12% at a tensile strength of375 MPa, while the conventional PM nanocompositeexhibits only a plasticity of 6% at a tensile strength of330 MPa. Therefore, nanocomposites with the high levelordering nanolaminates exhibited greatly improvedtensile strength and plasticity compared with nanocom-posites composed of the same constituents but randomlydistributed CNTs.
The increased strength was mainly attributed to thenanolaminate structure, which enhanced the 2-D align-ment of CNTs through geometric confinement and thus
Figure 4. Tensile properties of CNT/Al nanocomposites (a) loading with 1shows the relevant strengthening efficiencies of CNTs); (b) fabricated by vari
Figure 3. Microstructure of CNT/Al nanocomposites: (a) optical microscopTEM of conventional PM material; (e) Raman spectrum of raw CNTs, flak
enabled the improved longitudinal properties of CNTsin matrix. The strengthening efficiency of CNTs (R)can be expressed as:
R ¼ ðrc � rmÞ=V f rm ð2Þ
where rc and rm are the tensile strength of nanocompos-ites and matrix respectively, and Vf is volume fraction ofCNTs. Using the generalized shear-lag model, whenfibrous reinforcement has perfect alignment in the load-ing direction, reinforcement has the strengthening effi-ciency R, which can be expressed as S/2, where S isthe aspect ratio of reinforcement, i.e. the ratio betweenthe diameter and length of reinforcements [17]. In theCNT/Al nanocomposite by flake PM, the enhancedalignment of CNTs enables the R of 27, which is closeto the S/2 (�35 in this study) and much higher thanthe R (7.5) of conventional PM nanocomposites with3-D randomly distributed CNTs. Moreover, the otherfactor responsible for the enhanced strength is the main-taining of structure integrity of CNTs in flake PM nano-composites. As shown in Figure 3e, the relative intensityratio of D band to G band (ID/IG) of conventional PMnanocomposite increased to 1.91 while was only 0.7 forraw CNTs, which implies that the amount of defects inthe CNTs apparently increased. However, there was nodifferences in either the magnitude or the shape of thepeaks between the raw CNTs and flake PM CNT/Alnanocomposite. This implies that well-maintained integ-rity of CNTs appeared in flake PM nanocomposites andas a result, the potential of CNTs as reinforcements insuch materials is promising.
A better appreciation of the unique mechanicalproperties of the CNT/Al nanolaminated composites
vol.% CNTs and fabricated by conventional PM and flake PM (insetous methods; the data was drawn based on recent reviews (Refs. [4,5]).
y and (b) TEM of flake PM material; (c) optical microscopy and (d)e PM and conventional PM CNT/Al nanocomposites.
Table 1. Structure parameters and tensile properties of pure Al and CNT/Al nanocomposites.
Ref. Component Grainsize (nm)
Strength(MPa)
Uniformelongation (%)
Totalelongation (%)
Dislocationdensity (m�2)
Process
[18] Pure Al 180–200 334 1.8 7 1.33 � 1014 ARB[18] Pure Al 180–200 350 <1 <1 0.33 � 1014 ARB & annealingThis study CNT/Al 215 330 4 6 5.7 � 1014 Conventional PMThis study CNT/Al 230 375 8 12 11.5 � 1014 Flake PM
334 L. Jiang et al. / Scripta Materialia 66 (2012) 331–334
can be gained by comparing with those fabricated byother processes. As shown in Figure 4b, most CNT/Alnanocomposites with randomly distributed CNTs havea strength–ductility trade-off; that is, high strengthaccompanied with low ductility (points inside the purpleregion). However, the flake PM nanocomposites (a redpoint outside the purple region) exhibit both highstrength and good ductility, indicating the possibilityof retaining good ductility in CNT/Al nanolaminatedcomposites. This is mainly due to the nanolaminatestructure, which enables the full potential of CNTs tohinder recovery and thus keeps dislocations inside tinygrains. As shown in Table 1, the grain size of conven-tional PM and flake PM nanocomposite is respectively215 and 230 nm. Generally, such ultrafine-grained Alhas a very low uniform elongation (strain before neck-ing). Huang et al. [18] observed that ultrafine-grainedAl obtained by accumulative roll bonding (ARB) has astrength of �334 MPa, but a very low uniform elonga-tion less than 2%, and when such ARB Al was annealedat 150 �C for 30 min, the total elongation decreasedmarkedly to less than 1%, making the material almostbrittle. This is mainly because of a result of diminishingstrain hardening capacity due to the intrinsic difficulty inkeeping dislocations inside the tiny grains after anneal-ing or hot-deformation [11,18].
Our strategy is to make full use of the CNTs thathomogeneously distributed within and between metalnanolaminates to initiate, drag and pin dislocations,such that dynamic recovery could be reduced after hotextrusion. As shown in Table 1, the dislocation densityof flake PM nanocomposite (11.5 � 1014 m�2) is muchhigher than that of pure Al (1.33 � 1014 m�2) with sim-ilar grain size and about twice that of convention PMnanocomposite, indicating that the nanolaminate com-posites with 2-D aligned distribution of CNTs have ahigher recovery-hindering efficiency than that of com-posites with 3-D random distribution of CNTs. Signifi-cant dislocation storage is required for compatibleplastic strains of metal materials, allowing a highstrain-hardening rate, which leads to larger uniformstrains while maintaining a high level of strength[18,19]. Thus, the increased ductility was achieved inthe flake PM CNT/Al nanocomposites.
In conclusion, CNT/Al nanocomposites with nanola-minate structure over large-scale dimensions were fabri-cated by flake PM. The as-fabricated nanolaminateswith CNTs aligned between interlayers enabled longitu-dinal properties of 1-D CNTs and improved dislocationstorage capability, which assured both enhancedstrength and ductility over the nanocomposites com-posed of the same constituents but randomly distributedCNTs. Thus, the flake PM and relevant nanolaminatedesign were proved to be a practical and effective solu-
tion for strong CNT/Al nanocomposites with remark-able ductility. Finally, it is worth mentioning that theflake PM strategy toward creating nanolaminates canbe applied to other metallic or ceramic based bimaterialsreinforced with a variety of nanofibers or nanosheetssuch as graphene, while it can still be used as a referenceand for the enlightenment on the current advocation ofmicrostructural reinforcement.
The authors would like to acknowledge thefinancial support of the National Basic Research Pro-gram of China (973 Program, No. 2012CB619600), theNational High-Tech R&D Program (863 Program,No. 2012AA030611), the National Natural ScienceFoundation (Nos. 51071100, 51131004, 50890174), theInternational S&T Cooperation Program of China(Nos. 2010DFA52550, 2009DFA52410) and ShanghaiScience & Technology Committee (No. 11JC1405500).
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