Accepted Manuscript Creation of individual few-layer graphene incorporated in an aluminum matrix Weiwei Zhou, Yuchi Fan, Xiaopeng Feng, Keiko Kikuchi, Naoyuki Nomura, Akira Kawasaki PII: DOI: Reference:
S1359-835X(18)30236-7 https://doi.org/10.1016/j.compositesa.2018.06.008 JCOMA 5069
To appear in:
Composites: Part A
Received Date: Revised Date: Accepted Date:
3 November 2017 15 May 2018 6 June 2018
Please cite this article as: Zhou, W., Fan, Y., Feng, X., Kikuchi, K., Nomura, N., Kawasaki, A., Creation of individual few-layer graphene incorporated in an aluminum matrix, Composites: Part A (2018), doi: https://doi.org/10.1016/ j.compositesa.2018.06.008
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1
Creation of individual few-layer graphene incorporated in an aluminum matrix
Weiwei Zhou a,*,Yuchi Fan b, Xiaopeng Feng a, Keiko Kikuchi a, Naoyuki Nomura a, Akira Kawasaki a,*
a
Department of Materials Processing, Graduate School of Engineering, Tohoku University, Sendai
980-8579, Japan b
State Key Laboratory for Modification of Chemical Fibers and Polymer Materials, College of
Materials Science and Engineering, Donghua University, Shanghai 201620, China
ABSTRACT 3D-networks of few-layer graphene (FLG) platelets at grain boundaries, sandwiched between thin amorphous Al2O3 layers, were fabricated by spark plasma sintering (SPS) of graphene oxide (GO)/Al mixed powders. The GO was prepared by a modified Hummers’ method, and was thermally reduced to FLG simultaneously during SPS densification. Subsequent plastic flow of the Al matrix during the hot extrusion process caused the destruction of this structure, rearranged the FLG platelets individually into the uniaxial direction, and made them incorporate in the Al matrix. Observations by high-resolution transmission electron microscopy proved the existence of a direct-contact interface between the FLG and the Al matrix without any interfacial compounds, and revealed that the Al matrix featured a fairly low dislocation density. Consequently, the mechanical strength of Al matrix was noticeably enhanced by FLG incorporation, agreeing with the potential strengthening effect predicted by the load transfer mechanism.
KEYWORDS: A. Metal matrix composites (MMCs); B. Graphene; C. Aluminum; D. Interfaces.
*Corresponding authors. Email addresses:
[email protected] (W.W. Zhou),
[email protected] (A. Kawasaki) 1
1. INTRODUCTION Because of its large specific surface (~ 2630 m2g-1) [1], high fracture strength (~ 130GPa) [2], high Young’s modulus (~ 1TPa) [2], high thermal conductivity (~ 5300 Wm-1 K-1) [3], and unique electrical conductivity [4], graphene is considered an ideal reinforcement for composite materials. The lightweight metals such as Al strengthened by graphene are expected to become new high specific-strength materials with good electrical conductivity, showing a promising application prospect
for
new electric wires
[5].
So far,
however,
limited
researches regarding
graphene-reinforced metal-matrix composites (MMCs) has been reported because of the great difficulty in dispersing the individual graphene uniformly and homogeneously within the metal matrix [6, 7]. The mean particle size of commercially available metal powders is usually several tens of micrometers, while single-layer graphene (SLG) is 0.4 nm thick and several micrometers in size. This size difference makes difficult to achieve homogeneous dispersion of SLG in a metal matrix. In addition, the SLG is easily prone to be tangled as a result of the strong Van der Waal forces [8] and winkled because of its flexibility, which prevents the sheet-like form from being retained during processing. Furthermore, SLG has limited availability in bulk quantities. Instead, graphene platelets (GPLs) consisting of multiple layers of graphene (~ 50–200 layers) were prepared for use in previous studies of composite strengthening [5]. The GPL/Al composites have been almost exclusively fabricated by powder metallurgy techniques [9-20], because of the density discrepancy and the poor wettability between Al and GPL. To address the dispersion issue, a high-energy ball-milling (HEBM) technique has been mostly employed to exfoliate GPLs from graphite flakes and to prepare the initial GPL/Al mixed powders [5, 21]. Subsequently, the powder mixtures were mainly densified using hot rolling [9, 13], hot pressing [17, 22], or spark plasma sintering (SPS) [23-25]. For instance, Li et al. [17] attempted to incorporate GPLs into the Al matrix using HEBM, followed by hot pressing at 883 K for 4 h. However, the disadvantage of HEBM includes the fact that the structural integrity of the GPLs was severely damaged by the high-energy balls. In many prior studies [10, 11, 14, 15, 19, 21, 26-28], the existence and dispersion of GPL in MMCs were merely confirmed via energy dispersive spectroscopy (EDS) analysis along with 2
scanning electron microscopy (SEM) and/or transmission electron microscopy (TEM) to determine the existence of elemental carbon. In these studies, it was difficult to clarify unambiguously exactly how the GPLs were incorporated into the metal matrix. There are few reports revealing the evidence of the individually dispersed GPLs in an Al matrix and discussing the interface between the GPLs and the Al matrix [29-33]. Moreover, the GPLs usually included more than 150 layers of graphene, and the differences between their properties and those of graphite were too ambiguous to clarify [18, 19, 29]. Recently, Li et al. [12] reported the fabrication of few-layer graphene (FLG)/Al composites with a nanolaminated structure. Shin et al. [13] reported the processing of FLG/Al composites by HEBM and hot rolling. The reported results were interesting; however, the information about the incorporation of FLG in the Al matrix and the direct-contact interface between FLG and the Al matrix was still insufficient. A better understanding of the interface is needed to fully exploit the excellent properties of graphene. From the discussion above, we can see that it is imperative to fabricate a novel graphene/Al composite that addresses the great challenge of uniform dispersion of graphene in the Al matrix. This task is quite important before one may discuss the role of the interface between graphene and the Al matrix on the mechanical and physical properties, as well as on the interfacial reactions in Al-matrix composites. The fabrication process must guarantee that the GPL will be as thin as possible, that is, FLG. Also, the fabrication process must avoid structural damage to the FLG and must be a surfactant-less process to avoid residual contamination in the mixture of FLG and Al powders. The modified Hummers’ method is a common approach for the cost-effective mass production of graphene oxide (GO) sheets [34]. Carbon atoms in the basal plane of graphene are decorated with oxygen-containing groups, which tend to be hydrophilic. Hence, graphite flakes are exfoliated into few-layer GO and even single layer of GO in liquid media under moderate sonication [35]. Compared with the graphene sheets fabricated by chemical vapor deposition (CVD) or mechanical exfoliation [36], it is feasible and easier to incorporate defective GO into a metal matrix. More importantly, via a simple thermal reduction, few-layer GO can be easily reduced to few-layer reduced graphene oxide (r-GO) [37]; and Fan et al. [38] reported that r-GO retains excellent mechanical properties (e.g., Young’s modulus ~ 460 GPa). In this work, few-layer GO sheets synthesized by a modified Hummers’ method were employed 3
as the initial precursors of reinforcements. Although they are not single layer nor reduced into perfect graphene, we have attempted to fabricate a novel FLG/Al composite through the combination of a hetero-agglomeration method for mixing GO and Al powders, densification, and in-situ reduction by SPS and subsequent hot extrusion (HE). The high-resolution TEM (HRTEM) analysis proved the existence of FLG in the Al matrix and has shown that individually dispersed FLG platelets were incorporated in the Al matrix. The interface between FLG and the Al matrix has been studied in detail by HRTEM to verify the direct contact of FLG with Al and its interfacial reaction. The strengthening mechanism of FLG/Al composites was also discussed based on tensile properties.
2. EXPERIMENTAL PROCEDURES 2.1 Preparation of GO colloid GO was synthesized by modified Hummers’ method reported elsewhere [34] using purified natural graphite powders (SP-1, Bay Carbon, MI). Briefly, 1 g graphite was added to a flask and filled with 25 mL H2SO4 at room temperature, followed by addition of KMnO4 (3.5 g) gradually at 273 K in the ice-water bath. The mixture was stirred by magnetic stirring bar for 2 h under a temperature of 313 K. Then, 200 mL deionized water was added into this mixture, followed by the slow incorporation of 5 mL H2O2 (30 wt.%) to remove the excess KMnO4. The as-prepared GO quickly precipitated due to the strong acid environment. The sediment mixture was washed with 150 mL HCl solution by vacuum filtration. After dispersed into 200 mL water, the mixture was dialyzed for one week to remove the residual metal species. 1 L water was added to the purified mixture followed by sonication to form GO colloid. Finally, the colloid was centrifuged at 4000 rpm for 0.5 h to remove graphite and large GO sheets. The stable GO colloid with concentration of 4.95×10-4 g/mL [38] was obtained and used to prepare all the composites.
2.2 Preparation of GO/Al mixed powders The commercially available Al powders (Ecka granule Japan Co. Ltd.) fabricated by the gas atomization, with a purity of 99.85 % and an average particle size of ~5.5 µm, were used as the starting material in this study. 15 g Al powders were directly poured into a beaker and filled with 200 ml water. After the mechanical stirring and ultrasonication in the ice-water bath for 2 h, certain 4
amount of GO colloid (25, 50, 100, 200, 300, 500 mL) was added into the Al suspension by titration. Finally, the GO/Al hybrids were separated by filtration followed by drying at 343 K in vacuum environment. Correspondingly, the GO/Al mixed powders with the GO contents of 0.08 wt.%, 0.16 wt.%, 0.33 wt.%, 0.66 wt.%, 0.98 wt.% and 1.62 wt.% were prepared.
2.3 Consolidation of GO/Al mixed powders The mixed GO/Al powders were sintered by SPS (Dr. Sinter S511, Sumitomo Coal Mining Co. Ltd.), which was performed under a pressure of 50 MPa, sintering temperature of 873 K, holding time of 0.33 h and heating rate of 1200 K/h. The SPS-sintered (SPSed) compacts had a diameter of 15 mm and length of 30 mm. The SPSed bulks were hot extruded (UH-500kN1, Shimadzu Corporation, Japan) at 773 K with an applied pressure of 500 kN. The extrusion velocity and extrusion ratio were 0.06 m/h and 20, respectively. For comparison, a pure Al sample was prepared under the same procedures of powder mixing and drying, SPS and HE.
2.4 Characterization The apparent densities of composites were measured by Archimedes’s principle. The amounts of carbon element in composites after HE were evaluated by LECO (TC-444, LECO Co. Ltd. USA). The microstructures of as-prepared GO, GO/Al mixed powders and r-GO/Al matrix composites were analyzed by field emission scanning electron microscope (FESEM; JSM-6500F, JEOL, Japan), high-resolution transmission electron microscope (HRTEM; HF-2000EDX, Hitachi, Japan) and the electron back-scattered diffraction (EBSD) (OIM ver. 6, TSL Solutions, Japan). The TEM bulk specimens were prepared by grinding to the thickness of 50 μm, and thinned by the ion milling method (GATAN PISP Model 691, Gatan Inc.) under a voltage of 4 kV. The element distribution in the r-GO/Al composite was performed by electron probe micro-analysis (EPMA) (JEOL JXA-8530F) and JEOL JEM-ARM200F electron microscope. The thermo gravimetric analysis (TGA, SDT Q600, Hitachi, Japan) for GO was conducted in Ar protective atmosphere with a heating rate of 1200 K/h up to 873 K. The functional groups in the GO and r-GO were confirmed by X-ray photoelectron spectroscopy (XPS) with PHI5600 equipment (Ulvac Phi, Kanagawa, Japan). The crystallinity of GO and that in composites was evaluated by Raman spectroscopy (SOLAR TII Nanofinder, Tokyo 5
Instruments Co. Ltd., Japan) with 532 nm wavelength incident laser light. The tensile properties of pure Al and r-GO/Al composite were analyzed by tensile tests (AUTOGRAPH AG-I 50 kN, Shimadzu Co. Ltd., Japan at room temperature) with a strain rate of 1.67 × 10-3 s-1. The dogbone-shaped tensile samples having 2.5 mm in diameter and 10 mm in gauge length were directly machined from the hot-extruded bulks. Three tensile tests were carried out for each set of samples. The electrical conductivities of hot-extruded composites were evaluated using an Ulvac ZEM-3 system at room temperature.
3. RESULTS AND DISCUSSION 3.1 Characterization of prepared GO Fig. 1a shows an as-prepared GO prepared by the modified Hummer’s method. The GO is sheet-like and its edges tend to scroll and fold slightly (see the red dot in Fig. 1a). In addition, some crumples are observed (marked with a black dashed circle Fig. 1a), showing its flexibility. Fig. 1b shows the selected-area electron diffraction (SAED) pattern of the GO sheet taken from the position marked by the yellow square in Fig. 1a. This diffraction pattern shows the typical six-fold symmetry expected for graphene [39, 40]. Horiuchi et al. [41] reported that the layer number of graphene could be estimated from the diffracted intensity ratio. They mentioned I{1100}/I{2110} > 1 in the case of a monolayer, and I{1100}/I{2110} < 1 for multilayers with Bernal stacking. Labeling the peaks with the Miller-Bravais indices, Fig. 1b shows the (1-210) - (0-110) - (-1010) - (-2110) axis for the SAED pattern. Fig. 1c shows the diffraction intensity taken along the (1-210) to (-2110) axis for the SAED pattern in Fig. 1b. The {1100} inner peaks appear to be more intense than the {2110} outer ones; that is, I{1100}/I{2110} > 1. Thus, we may conclude that this GO sheet is a monolayer. On the other hand, TEM observations revealed that the prepared GO was composed of monolayers and few-layer (< 10) structures. Centrifugation was performed to remove thick GO flakes and thereby guarantee consistently thin GO platelets in this study. The average number of graphene layers will be discussed later, based on HRTEM analysis using r-GO/Al composites. It was confirmed by means of dynamic light-scattering (DLS) measurements that the average lateral size of the GO was around 700 nm [42]. This is much smaller than the Al powders used in this study (~ 5.5 µm).
6
3.2 GO/Al powder mixture and densification by SPS It is known that the surface of GO is functionalized by tertiary alcohols, ethers, and carboxyl groups [43]. Because of the abundance of these functional groups, GO exhibits a significantly negative charge which makes it hydrophilic, resulting in the easy exfoliation of graphite into GO sheets when it is dispersed in water. In general, the surface of gas-atomized Al powders is covered with very thin aluminum oxide [44]. Since the surface of alumina (Al2O3) is positively charged in aqueous media, Al powders are repulsively dispersed under mechanical stirring. Therefore, GOs and Al powders are attracted to each other when GO colloid is added drop-by-drop into an Al suspension (Fig. S1 of Supporting information). Fig. 2 shows FESEM images of GO/Al powder mixtures. GO can be observed on the surface of Al powders with many crumples (the black dashed circle in Fig. 2b). Few GO clusters were observed in GO/Al powder mixtures at a GO concentration of 0.66 wt.% (Fig. 2c). However, at 1.62 wt.% of GO, the aggregation of GO was largely present among Al particles (Fig. 2d). One of the reasons is the limited contact surface of Al particles compared to the total surface area of GO sheets at this concentration during powder mixing. Moreover, the GO sheets are likely to be crumpled and winkled. Fig. 3a shows a typical example of a GO-wrapped structure. As mentioned above, there is a significant size difference between GO sheets (~ 700 nm) and Al powders (~ 5.5 μm). Hence, an Al particle is covered with several GO sheets under electrostatic attraction during hetero-agglomeration. Thanks to the remarkable flexibility and ultra-thin thickness, some GO sheets are completely attached and fitted to a spherical Al particle (Fig. 2b). Meanwhile, some GO sheets are bridging two neighboring Al particles (Fig. 3b), so that the GO sheets are overlapped and attached to each other, forming a network structure. In contrast, the GPLs generally keep a flaky shape due to the high out-plane stiffness, thus they are not easily deformed or bridged to Al particles. Correspondingly, this unique network structure has never before been observed in GPL/Al powder mixtures [29]. Furthermore, the potential contribution of the GO network structure to the comprehensive performance of MMCs, e.g., hindering the grain growth during sintering, pinning the movement of dislocation or grain boundaries under loading, as well as improving the electrical or thermal conductivity of the matrix, should be systematically investigated in future work.
7
The relative densities of the SPSed composites with GO concentrations of 0.08-0.98 wt.% were 98-99 % (Table 1). Fig. 4 shows the Raman spectra of GO and SPSed composites. The Raman spectra include the D and G bands; the D band at 1350 cm-1 is usually associated with the presence of lattice disorders, while the G band at 1580 cm-1 strongly corresponds to in-plane C-C symmetric stretching vibrations in graphene sheets. The degree of structural defects or disorders in sp 2-based carbon can be evaluated by the intensity ratio of the D and G bands (ID/IG) in the spectra [45]. The ID/IG ratio increases as structural defects increase. However, it is known that too many structural defects cause D and G peaks to broaden, and the ID/IG ratio of GO then becomes nearly unity, even though the GO includes an abundance of structural defects [46]. As a consequence of this fact, the ID/IG ratio would first increase during reduction of GO. Huh [47] reported the variation of the ID/IG ratio during thermal reduction of GO; the ID/IG ratio increased gradually with the reduction temperature in the range from 473 to 1273 K, and subsequently decreased in the range from 1273 K to 2273 K. The cracking of two bonds of C-O-C probably causes the simultaneous reconstruction of adjacent C=C aromaticity during the thermal reduction. Thus, the ID/IG ratio increased slightly from 0.97 to 1.09 after SPS (Fig. 4). It is thought that the recovery of GO during SPS gives rise to this increased ID/IG ratio. In addition to the G and D bands, two Raman bands with weaker intensity, namely 2D and D+G at 2600-3000 cm-1 were detected. The 2D band is sensitive to the aromatic carbon structure whereas the combination mode of D+G is induced by lattice disorder in crystalline graphitic materials [47]. The 2D band of SPSed composites reappeared with an apparently increased I2D/ID+G ratio, because of the recovery of graphitic electronic conjugation. It was found that GO was effectively recovered and became reduced-GO (r-GO) after SPS, which is consistent with the results of TGA and XPS analysis (Fig. S2 and S3 of Supporting information). Recently, Zandiatashbar et al. [48] reported a simple and non-destructive approach to quantifying the mechanical properties of graphene by measuring the Raman spectrum. In this work, therefore, the young’s modulus and fracture strength of r-GO were estimated to be 320 GPa and 23.2 GPa, respectively. Fig. 5b shows HRTEM-EDS mapping at a grain boundary taken from the area inside the white square in Fig. 5a. The element C is correlated with r-GO. Al containing oxygen corresponds to Al2O3 that originates from the thin Al2O3 layer (~ 5 nm) on the raw Al powders induced by gas atomization [44]. The crystallinity of the Al2O3 layer was determined from the corresponding SAED pattern [see 8
the upper inset in (c)]. Since the employed beam size (several tens of nanometers in diameter) exceeded the thickness of the nano-Al2 O3 layer (~ 5 nm in average thickness), the diffracted spots of the Al matrix are also included in the SAED pattern. However, it can be clearly observed that the Al2O3 layer possesses the typical amorphous feature of a hollow pattern (see the dotted circle in 5c). Thus, as shown in Fig. 5c and 5d, the r-GO platelets were sandwiched between the thin amorphous Al2O3 layers. This is clear evidence that r-GO platelets are located at the grain boundary. The detailed analysis at the red lines in Fig. 5c and 5d allows us to determine the number of graphene layers at the boundaries. As illustrated in the insets, the r-GO platelets contain 5 layers of graphene in Fig. 5c and 3 layers of graphene in Fig. 5d. The interlayer distance was 0.4~0.5 nm, which is larger than the d-spacing of graphite (0.334 nm) but smaller than that of GO (0.7 nm) [49]. Since platelets with more than 10 layers of graphene have been barely observed by HRTEM, we may deduce that the r-GO platelets in grain boundary are FLG platelets. Fig. 5a shows that FLG continually exists at grain boundaries and the FLG is fitted to curved grain boundaries, indicating its flexibility. Fig. 5d shows that the interface between the FLG and Al2O3 layer is intimate. Fig. 6 shows the microstructure of the transverse cross section of a 0.4 vol.% FLG/Al composite. The average grain size is 4.8 μm, which is almost the same as the mean particle size of the Al powders, indicating less grain growth. That was possibly due to the pinning effect of the incorporated FLG and the remaining Al2O3 layers [29]. In addition, the grain size of the Al matrix was independent of FLG concentration (Table 1 and Fig. S4 of supporting information). Different from the SPSed pure Al (Fig. S5 of supporting information), the particle boundaries of composite, including FLG platelets, shown in Fig. 5 are observed as continuous white outlines in Fig. 6. These outlines surround all the grains with a practically uniform thickness. Fig. 7 shows the EPMA mapping, and those lines consist of elemental carbon. Therefore, it is noted that FLG was uniformly distributed at the grain boundaries in the SPSed compacts and FLG has been formed in 3D-networks along grain boundaries. The FLG platelets are probably overlapped and attached to each other.
3.3 Microstructure of the FLG/Al composites after SPS and HE
9
The combination of SPS and HE has been very effective for preparing fully dense FLG/Al composites, and all the extruded samples possessed a relative density greater than 99.5 % (Table 1). Fig. 8 shows the EBSD mappings of the 0.4 vol.% FLG/Al composite. The primary Al particles are elongated into spindle-like shapes and oriented along the extrusion direction providing a strong fiber texture to the Al matrix (Fig. 8b and Fig. S6 of supporting information). The average grain size in the transverse cross section of the extruded composite is 2.2 µm (Fig. 8a), so that the grain size in the transverse cross section is smaller than that of the SPSed one (~ 4.6 µm). Therefore, the distance between two adjacent FLGs is seen to decrease by 50 % in the transverse cross section. The HRTEM observations shown in Fig. 9a and 9c clarified that individual FLG platelets are aligned and incorporated in the Al matrix in the direction parallel to the HE direction. However, their angles of rotation are unknown. Fig. 9b shows that a FLG is in close contact with Al2O3 on one side; however, the FLG is in direct contact with the Al matrix on other side. Fig. 9d shows a FLG in which both sides of the FLG are in direct contacted with the Al matrix. It is deduced that the plastic flow of the Al matrix during HE process destroyed the FLG network at the grain boundary and rearranged individual FLG platelets along the direction of plastic flow. The plastic flow may contribute to stretch the winkled FLG platelets and also to flattening. At the same time, the thin Al2O3 layer on the surface of the Al powders was broken by the plastic flow, which produced a new surface of Al and caused direct contact between the FLG and the Al matrix. The Raman spectrum analysis revealed that the quality of the FLG was retained after HE (see Fig. 4). The unidirectional arrangement of FLG platelets may cause anisotropic mechanical properties of the FLG/Al composite. A related experiment is now being conducted in our laboratory; the results and a discussion will appear in our next report. It is known that the mismatches in the coefficients of thermal expansion between the Al (23.6 × 10-6 K-1) and graphene (1 × 10-6 K-1) may cause the prismatic punching of dislocations at the interface [21]. However, few dislocation lines were detected by HRTEM at the interfaces between the FLG and the Al matrix. Moreover, a very low dislocation density was observed in the Al matrix, primarily owing to the dynamic recovery of the Al matrix during the high-temperature extrusion process (~ 773 K). Fig. 4, Fig 9b and Fig 9d suggest that the formation of Al4C3 was unlikely to occur at the FLG-Al interfaces. It has been demonstrated that the fabricated FLG/Al composites are fully dense, with individual 10
FLG platelets aligned mostly in the extrusion direction, having a fairly low dislocation density and a direct FLG/Al interface without any interfacial chemical compound. Such a clean system allows us to discuss the direct contribution of FLG on the tensile responses of Al matrix and the exploitation of the strengthening capability of FLG.
3.4 Mechanical and electrical properties of FLG/Al composites after SPS and HE Fig. 10 (a) shows the typical nominal tensile stress-tensile strain curves of the FLG/Al composites and of pure Al. The UTS and YS of Al matrix proportionally increased with the volume fraction of FLG platelets up to 0.2 vol.%; subsequently, they increased slowly up to 0.4 vol.%, reaching the maximum values of 173.0 MPa and 146.2 MPa, respectively. However, beyond 0.4 vol.%, the UTS and YS appeared to be saturated and independent of the FLG concentration. In general, the MMCs can be enhanced by a synergetic mechanism of grain refinement [6], geometrically necessary dislocation formation (due to different Young’s modules and coefficients of thermal expansion between the matrix and reinforcement) [50], work hardening [6, 29], Orowan strengthening [29] and load transfer if the aspect ratio of reinforcement and interfacial bonding are sufficient [9]. Concerning this FLG/Al composite, the strength enhancement by the refinement of Al grain (Table 1) and by the dislocation generation remained minimal (Fig. 9b and 9d). Moreover, there was a fairly low dislocation density in the Al matrix after hot extrusion while the unidirectionally aligned 2D FLG platelets are unlikely to prevent the dislocation pile-up and accumulation (Fig. 9a and 9c). Thus, it seems that the load transfer strengthening should be the dominating strengthening mechanism for the FLG/Al composite. In this case, the shear lag model is usually used to estimate the strength of the composites [6, 13, 50, 51]. In this model, it is assumed that the fiber reinforcements perfectly contacted with the matrix are aligned to a single direction [51]. Kelly-Tyson’s formula is written by [51] c f V（f 1
c fVf
c ） m (1 V f ) ( ≧ c ) 2
m (1 V f ) 2 c
11
( c )
(1) ( 2)
where
c and m are the yield strength of the composite and the matrix, respectively; f , V f ,
d f and are the strength, volume fraction, and average length of the fiber; c is the critical 2 m fiber length, indicating the minimum in which the fiber is loaded up to its fracture strength; d is the fiber diameter;
m
is the interfacial shear strength (~ 0.5
m ). However, FLG is a two-dimensional
thin sheet, so that appropriate modification of c is necessary. Shin et al. [13] recently demonstrated this modification and derived the following equation:
c f
A mS
(3)
where A( wt) is the cross-sectional area of FLG; S [（2 w t )] is the interfacial area; w and t are the width and thickness of FLG, respectively. In this work, the FLG platelet is assumed to be a square-like sheet, we have w = 700 nm [42]. Meanwhile, the average thickness of FLG, t, is considered as 2.5 nm (~ 5 layers of graphene). According to Eq. (3), the critical FLG length,
c , is
calculated to be 524 nm, considering the value of f is 23.2 GPa as beforementioned. Thus, in this FLG/Al composite, c . As demonstrated for this FLG/Al composite, individual FLG platelets have an intimate interface with Al matrix, and are mostly aligned along the extrusion direction (parallel to the tensile direction), indicating the requirement of the assumption is satisfied. The estimated yield strength of the FLG/Al composites using Eq. (1) is shown in Fig. 10b by the red dashed line. Interestingly, the experimental YS values (see the black triangles in Fig. 10b) are well consistent with the estimated values in composites with up to 0.2 vol.% of FLG. This result indicated the tensile load could be transferred effectively from the matrix to the FLG platelets, suggesting the interface close to a perfect one. The intimate interface between the FLG and Al matrix is probably attributed to the existence of native oxygen on FGL surface [52]. The oxygen atoms are expected to form strong covalent bonding with Al atoms, improving the interfacial adhesion energy [52]. The deviation of the experimental values at the FLG concentration beyond 0.2 vol.% may be attributed to the agglomeration of FLG platelets
12
(Fig. S7 of supporting information), preventing formation of intimate interfaces with the matrix, which is a prerequisite for load transfer. On the other hand, the advantage of FLG over other reinforcements can be discussed by comparing the strengthening efficiency, R, as follows [29]:
R
c m Vf m
( 4)
The FLG/Al composite showed a high R of ~120. This value is higher than R of 10–60 reported in GPL/Al composites [9, 11, 12, 17, 19, 26, 53], and R of 1–30 for many commonly used reinforcements (e.g., nanoceramic [54] and carbon nanotube [9, 55, 56]). This result shows the potential of individual FLG platelets as a reinforcement in MMCs. Considering the application of electric wires, the room-temperature electrical conductivities of pure Al and FLG/Al composites were measured. Table 2 shows that the electrical conductivity of Al matrix was slightly decreased with increasing the FLG contents. That is because the FLG platelets by Hummer’s method and subsequent thermal reduction contain abundant structural disorders or oxygen-containing functional groups, possessing a relatively lower electrical conductivity (~ 8×103 S.m-1 [57]) than pure Al. Nevertheless, the FLG/Al composites could still remain high electrical conductivities, e.g., the 0.4 vol.% FLG/Al composite has an electrical conductivity of 3.43×107 S.m-1 (~ 95% of pure Al), simultaneously showing a UTS of 173 MPa and failure elongation of 16.2 % (see Table S1 of Supporting information). Accordingly, our results strongly suggest that the FLG/Al composite has a great application prospect in electric conductor fields.
4. CONCLUSIONS Individual FLG platelets were incorporated in an Al matrix by means of SPS sintering, followed by a hot-extrusion process. The GO sheets were prepared by a modified Hummers’ method, and they were mixed with the Al powders in an aqueous solution utilizing a hetero-agglomeration method without the assistance of surfactants. This strategy produced GO platelets that covered the surface of the Al particles and interconnected with each other as a network structure. A 3D-network of FLG platelets in the grain boundaries, sandwiched between thin amorphous Al2O3 layers, was created by spark plasma sintering of GO/Al mixed powders. Subsequent plastic flow of the Al matrix by the hot 13
extrusion process caused the destruction of this structure, rearranged the FLG platelets individually in the uniaxial direction, and made them incorporate into the Al matrix. The HRTEM observations proved the existence of a direct-contact interface between the FLG and the Al matrix without formation of Al4C3 impurities, and revealed that the Al matrix had a fairly low dislocation density. Therefore, we have successfully fabricated a novel FLG/Al composite that is fully dense, with individual FLG platelets aligned mostly in the extrusion direction, having low dislocation density and a direct FLG/Al interface without Al4C3 formation. Furthermore, the strength of FLG/Al composites increased proportionally with the volume fraction of FLG up to ~0.2 vol.%. The experimental yield strength was well consistent with the estimated values by shear lag model. It can be concluded that the strength enhancement of FLG/Al is mainly ascribed to the load transfer strengthening; the promising load-bearing capability of the FLG platelets has been exploited effectively during loading.
ACKNOWLEDGMENT The authors would like to thank Dr. Takamichi Miyazaki for his sophisticated skill and beneficial discussion in TEM. We also appreciate the generous helps from Dr. Kosei Kobayashi, Pavlina Mikulova, Xiaohao Sun in Tohoku University.
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19
FIGURE CAPTIONS
Fig. 1 (a) TEM image of as-prepared GO on a copper grid; (b) SAED pattern of GO taken from the position of yellow square in (a); (c) diffracted intensity taken along the 1-210 to -2110 axis for the pattern in (b).
Fig. 2 FESEM images of GO/Al mixed powders with different GO contents: (a) 0 wt.%; (b) 0.16 wt.%; (c) 0.66 wt.%; (d) 1.62 wt.%. The black dashed circle in (b) shows the crumpled GO on an Al surface; the white arrow in (d) displays the aggregation of GO sheets among Al particles.
Fig. 3 TEM images of the 0.66 wt.% GO/Al mixed powders. The black arrow in (b) shows the GO sheets bridging two Al particles.
Fig. 4 Raman spectra of GO, SPS-sintered composite and the composite after SPS and HE.
Fig. 5 (a) TEM image and (c, d) HRTEM images for the 0.4 vol.% r-GO/Al composite after SPS; (b) a bright field (BF) TEM image taken from the position of white square in (a) and the corresponding EDS mappings of elemental C, O, Al in this area. Insets in (c) show the SAED pattern of Al2O3-layer taken from yellow spot and the profile plot along red line; insets in (d) show the SAED pattern of Al matrix taken from the yellow spot and the profile plot along red line.
Fig. 6 FESEM images for the transverse cross section of SPS-sintered 0.4 vol.% FLG/Al composite.
Fig. 7 Backscattered electron (BSE) image and EPMA mappings of elemental C, O, Al for the transverse cross section of SPS-sintered 0.4 vol.% FLG/Al composite.
Fig. 8 Inverse polar figure (IPF) mappings for (a) transverse and (b) longitudinal cross section of 0.4 vol.% FLG/Al composite after hot-extrusion.
20
Fig. 9 (a, c) TEM images and (b, d) HRTEM images of the hot-extruded 0.4 vol.% FLG/Al composite. (b) is taken from the position of white square in (a); insets in (b) show the EDS analysis of Al2O3-layer taken from red spot, the SAED pattern of Al matrix taken from yellow spot and the profile plot along white line; insets in (d) show the Fast Fourier transform (FFT) pattern of Al matrix taken from yellow spot and the profile plot along white line.
Fig. 10 (a) Nominal tensile stress-strain curves of pure Al and FLG/Al composites; (b) the experimental YS values versus the volume fraction of FLG, and the estimated YS values by using shear lag model. The inset in (b) shows the schematic depiction of FLG platelet under loading.
TABLE
CAPTIONS
Table 1 Characteristics of GO/Al mixed composites, SPS-sintered r-GO/Al composites and the composites after SPS and HE. The volume fractions of r-GO in the composites were calculated based on the density of r-GO (2.28 gcm-1) and Al (2.7 gcm-1).
Table 2 Electrical conductivities of pure Al, and FLG/Al matrix composites with various FLG concentrations.
21
Fig. 1 (a) TEM image of as-prepared GO on a copper grid; (b) SAED pattern of GO taken from the position of yellow square in (a); (c) diffracted intensity taken along the 1-210 to -2110 axis for the pattern in (b).
22
Fig. 2 FESEM images of GO/Al mixed powders with different GO contents: (a) 0 wt.%; (b) 0.16 wt.%; (c) 0.66 wt.%; (d) 1.62 wt.%. The black dashed circle in (b) shows the crumpled GO on an Al surface; the white arrow in (d) displays the aggregation of GO sheets among Al particles.
23
Fig. 3 TEM images of the 0.66 wt.% GO/Al mixed powders. The black arrow in (b) shows the GO sheets bridging two Al particles.
24
Fig. 4 Raman spectra of GO, SPS-sintered composite and the composite after SPS and HE.
25
Fig. 5 (a) TEM image and (c, d) HRTEM images for the 0.4 vol.% r-GO/Al composite; (b) a bright field (BF) TEM image taken from the position of white square in (a) and the corresponding EDS mappings of elemental C, O, Al in this area. Insets in (c) show the SAED pattern of Al2O3-layer taken from yellow spot and the profile plot along red line; insets in (d) show the SAED pattern of Al matrix taken from the yellow spot and the profile plot along red line.
26
Fig. 6 FESEM images for the transverse cross section of SPS-sintered 0.4 vol.% FLG/Al composite.
27
Fig. 7 Backscattered electron (BSE) image and EPMA mappings of elemental C, O, Al for the transverse cross section of SPS-sintered 0.4 vol.% FLG/Al composite.
28
Fig. 8 Inverse polar figure (IPF) mappings for (a) transverse and (b) longitudinal cross section of 0.4 vol.% FLG/Al composite after hot-extrusion.
29
Fig. 9 (a, c) TEM images and (b, d) HRTEM images of the hot-extruded 0.4 vol.% FLG/Al composite. (b) is taken from the position of white square in (a); insets in (b) show the EDS analysis of Al2O3-layer taken from red spot, the SAED pattern of Al matrix taken from yellow spot and the profile plot along white line; insets in (d) show the Fast Fourier transform (FFT) pattern of Al matrix taken from yellow spot and the profile plot along white line.
30
Fig. 10 (a) Nominal tensile stress-strain curves of pure Al and FLG/Al composites; (b) the experimental YS values versus the volume fraction of FLG, and the estimated YS values by using shear lag model. The inset in (b) shows the schematic depiction of FLG platelet under loading.
31
Table 1 Characteristics of GO/Al mixed composites, SPS-sintered r-GO/Al composites and the composites after SPS and HE. The volume fractions of r-GO in the composites were calculated based on the density of r-GO (2.28 gcm-1) and Al (2.7 gcm-1). Content of GO in
Relative density
Relative density
Concentration
Grain size
Grain size
mixed powders
after SPS
after HE
of r-GO
after SPS
perpendicular
(wt.%)
(%)
(%)
(vol.%)
(µm)
to HE (µm)
0
100
100
-
4.80±1.43
2.88 ± 0.94
0.08
99.84
99.60
0.05
-
2.46 ± 0.83
0.16
99.56
99.58
0.1
-
-
0.33
99.98
99.99
0.2
4.47 ± 1.47
2.14 ± 0.74
0.66
99.66
99.70
0.4
4.58 ± 1.50
2.17± 0.70
0.98
98.07
99.70
0.6
-
2.38 ± 0.85
1.62
96.74
99.69
1.0
4.51 ±1.44
-
32
Table 2 Electrical conductivities of pure Al, and FLG/Al matrix composites with various FLG concentrations. Concentration 0
0.05
0.1
0.2
0.4
1.0
3.61
3.57
3.53
3.50
3.43
3.22
of r-GO (vol. %) Electrical conductivity (×107 S.m-1)
33
