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Effect of Interface Structure on the Mechanical Properties of Graphene Nanosheets Reinforced Copper Matrix Composites Xiang Zhang, Chunsheng Shi, Enzuo Liu, Naiqin Zhao, and Chunnian He ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b09799 • Publication Date (Web): 09 Oct 2018 Downloaded from http://pubs.acs.org on October 10, 2018

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ACS Applied Materials & Interfaces

Effect of Interface Structure on the Mechanical Properties of Graphene Nanosheets Reinforced Copper Matrix Composites Xiang Zhang †, Chunsheng Shi †, Enzuo Liu †, ‡, Naiqin Zhao †, ‡ *, Chunnian He †, ‡ * †

School of Materials Science and Engineering and Tianjin Key Laboratory of Composites and Functional Materials, Tianjin University, Tianjin, 300072, P. R. China

‡

Collaborative Innovation Center of Chemical Science and Engineering, Tianjin 300072, China

* Corresponding author. E-mail addresses: [email protected] (C.N. He), [email protected] (N.Q. Zhao)

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ABSTRACT: Currently, seldom studies have paid close attention to the impact of the defects and oxygen-containing functional groups on the surface of the graphene for composite applications. In this work, two typical graphene materials, namely graphene nanosheets synthesized by an in-situ catalytic reaction and reduced graphene oxide (RGO), were adopted to fabricate reinforced copper matrix composites by spark plasma sintering. A harmful transitional interfacial layer made up of Cu/CuOx/amorphous carbon/RGO, resulted from interfacial reaction between Cu and RGO, were observed in the RGO/Cu composite. In contrast, the in-situ synthesized graphene with fewer defect and oxygen level can realize clean graphene-Cu interface with Cu-O-C bonding and thus lead to much improved interface bonding and superior yield strength and tensile ductility. These results imply that the in situ synthesized graphene is more favorable for achievement of robust interfacial bonding for enhancing the mechanical properties of the graphene-Cu composites.

KEYWORDS: metal matrix composites (MMCs), copper, graphene, interface, mechanical properties

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

INTRODUCTION Graphene, certified with exceptional physical and mechanical properties associated

with its unique single layer of hexagonal lattice structure, is reckoned as the ideal reinforcement for producing metal matrix composites (MMCs).

1-4

Recently, there have

been a lot of works reported on the fabrication of graphene reinforced MMCs (Al, Cu, Mg, Fe and etc.) by powder metallurgy, laser sintering, and so on.

5-17

However, major

challenges in the synthesis of graphene-reinforced MMCs lie in the difficulty of incorporating and distributing evenly graphene in the matrix. In addition, the large specific area and high aspect ratio endow graphene a large contact surface with metals, therefore the chemistry, atomic structure and the bonding at the interface between graphene and the matrix plays a crucial role in determining the overall properties of the composite material. 9 Stiffening and strengthening rely on the mechanical load transfer from the matrix to higher strength reinforcement across the interface,

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while toughness is influenced by crack

deflection or bridging at the interface and ductility is affected by relaxation of peak stresses near the interface.

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For achieving better mechanical properties in the MMCs, the metal

matrix should possess strong interface bonding with the reinforcing phase in order to avoid microscale cavity formation and enable robust interface adherence to restrain any delamination at the interface, and thereby allow for effective load transfer.

20, 21

From this

point of view the interfacial structure of the composites, which can tailor the microstructures with specific mechanical properties, is highly needed to be investigated and controlled. On the other hand, in the respect of practical applications, the single-layered or few-layered graphene with approximately perfect structure produced by the physical

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exfoliation or chemical vapor deposition grown on metal (such Cu, Ni, etc) foil method cannot satisfy the needs of mass production due to the highly time-consuming and expensive fabrication process. While in the recent reported work, the easily and massively prepared multi-layered graphene derivatives, namely reduced graphene oxides (RGO), 15, 22

graphene nanoplates,

nanosheets

10

6, 12, 14, 23

graphene nanoribbons,

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8, 9,

in-situ synthesized graphene

and etc. exhibit outstanding prospect of fabricating MMCs. However, these

graphene materials for composite applications possess defects and oxygen-containing functional groups (OCFGs) introduced from raw materials or preparation process, which would make them more reactive than perfect graphene

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and thus impacts the interfacial

structure and bonding state between graphene and metals remarkably.

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During the

high-temperature fabrication process of MMCs, the activated oxygen atoms obtaining energy from Joule heat are driven out of the graphene surface and move towards interface even possibly diffuse into the metal bulk as a result of differential oxygen concentration between metal and graphene.

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At the meantime, the local stress resulted from the

restriction of matrix plastic deformation by the graphene reinforcement also makes the interfacial condition even more complex.

15-17

To clarify the interfacial bonding between

these "defected graphene" and metal matrix, a comprehensive consideration of the existence type and formation mechanisms of the defects and oxygen element in the graphene and their potential interfacial reaction is very necessary. Taking one of the most promising matrix copper as an example, however, until now there has been only very few research reported in this area: Huang et al. 9 used a molecular-level mixing combined spark plasma sintering (SPS) to make RGO/Cu composites. They asserted that Cu-O-C bond on the interface improves the load transfer efficiently, but detailed studies were not conducted

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to ascertain interfacial oxygen type; Chu et al.

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discovered nanometer-sized amorphous

carbon layer on the interface between SPSed RGO/Cu and attributed the amorphization to the structure damage of RGO during the harsh SPS condition; coincidentally, Cho et al.

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also reported the existence of amorphous interfacial layer in the SPSed multi-layered carbon nanotubes/Cu composites. They interpreted the transition layer as the reaction product of the copper oxide outside Cu powder with the active amorphous carbon on the carbon nanotube surface. In addition, according to the research by Hui et al., 29 the OCFGs on the surface of graphene have great potential to oxidize the contacted copper film in which process the law of thermodynamics dominate. Following this concept, the interfacial state in the bulk composites would change significantly once the oxidation products formed, due to the barrier of the direct contact between graphene and copper. The results above definitely confirm that the defects and OCFGs in the graphene are highly active during the course of compositing with copper. Unfortunately, these preliminary studies have not provided a global understanding of the characteristics and possible influence of interfacial composition and structure. Moreover, the fundamental insights related to the formation mechanisms responsible for the interface characteristics and the influence of different graphene structure on the interface also remain absent in the literature. In view of the above issues, we adopted two typical graphene reinforcements to produce reinforced copper matrix composites: one is graphene nanosheets synthesized by an in-situ catalytic reaction; the other is a widely used RGO powders. The former has less defects and lower oxygen level on the graphene surface benefited by metal catalyst, while the latter is characterized with more defects and higher oxygen level derived from the strong chemical exfoliation process. The two groups of composites (namely 0.5GN-SPS

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and 0.5RGO-SPS) with the same reinforcement content of 0.5 vol. % were fabricated by a modified molecular-level mixing method and rapid SPS consolidation to achieve fully densified composites and to avoid damage to the graphene nanosheets during the compositing and sintering processes. Through the careful comparisons based on high resolution transmission electron microscopy (HRTEM), electron energy loss spectroscopy (EELS), energy-dispersive X-ray spectroscopy (EDXS) characterizations and in-situ TEM tensile tests, in this work we first ascertained the formation of different interface structure (clean GN-Cu interface for 0.5GN-SPS and complex interface containing transition layer for 0.5RGO-SPS) and their effects on the mechanical behaviors of the graphene/copper composites. In addition, the catalytic synthesis process of the GN reinforcement and the formation mechanism of the interface in the composites were investigated in detail. Moreover, the calculation based on first-principle theory was conducted to qualitatively discuss the effect of interfacial oxygen level on the binding energy of the interface to further verify the experimental results.



EXPERIMENTAL SECTION Materials and Reagents. Cupric nitrate trihydrate (Cu(NO3)2·3H2O), glucose

(C6H12O6), sodium chloride (NaCl) (the three regents above for the preparation of the GN anchored

with

Cu

nanoparticles

(GN@Cu

NPs)

powder),

cupric

acetate

(Cu(CH3COO)2·H2O), and ammonia (NH3·H2O) were purchased from Tianjin Kemiou Chemical Reagent Co., Ltd, purity≥99.5%. The graphene oxide nanosheets were purchased from Nanjing XFNano Material Tech Co. Ltd, purity>98%. All of these materials and reagents were used without further purification.

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Preparation of GN@Cu@Cu and RGO@Cu Composite Powders. Firstly, the GN anchored with Cu nanoparticles (GN@Cu NPs) (the GN content in the GN@Cu powders was calculated to be ~28.6 wt.% according to Thermogravimetric Analysis (TGA)) were in-situ synthesized by a NaCl template-assisted high-temperature calcination strategy, the detailed introduction of the preparation process could be found in our previous work. 10 The second step is to further coat the GN@Cu with copper through an impregnation-reduction process in order to form GN@Cu coated with impregnated-reduced copper NPs (indicated as GN@Cu@Cu) as well as tune the ratio of reinforcement phase to Cu. To make 40.0 g 0.5 vol.% GN@Cu@Cu powders, 124.242 g Cu(CH3COO)2·H2O was dispersed into 292 mL NH3·H2O under magnetic stirring, then 00.173 g GN@Cu powders were added in the solution, followed by ultrasonificaiton for 0.5h. Then the solution was vaporized with magnetic stirring at 95 °C and dried at 200 °C in a drying oven. Afterwards, the dried and grounded fine powders were firstly calcinated at 400 °C for 1h under Ar and then reduced at 400 °C for 2 h under H2 to form GN@Cu@Cu powders. The powders after reduction were immediately vacuumed into the sealed bags before transferring to graphite mould for SPS consolidation. So the oxidation of the small-sized Cu powders could be avoided which guarantees their sintering capability in the SPS process. The experiments to fabricate reduced graphene oxide (RGO) wrapped by copper (indicated with RGO@Cu), were conducted under a similar impregnation-reduction method using graphene oxide (GO) (0.050 g to make 40.0 g RGO/Cu powders) as the raw materials. During the processing, the GO nanosheets were reduced at the same time with CuO by H2. The GO powder used in this work was purchased from Nanjing XFNano Material Tech Co. Ltd., China. The volume fraction of graphene reinforcement is a controllable nominal

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content without considering the loss during preparation. In fact, it is very difficult to accurately calculate the exact content of nanocarbon reinforcement in the graphene/Cu composite powders as well as bulk materials due to the small volume fraction (0.5 vol. %) of graphene. Fabrication of the GN/Cu and RGO/Cu Bulk Composites. Typically, the GN@Cu@Cu or RGO@Cu powders were placed in a graphite mold and then were consolidated into bulk composites with a pressure of 50 MPa at 600 °C for 10 min under vacuum of 10-1 Pa level. During the heating process, the pressure was applied at 400 °C and slowly increased to 50 MPa with the temperature raising until 600 °C. In this way, the adsorbed gas on the powder surface could be easily escaped so that the densification of the bulk composites could be achieved (the relative density of 0.5GN-SPS and 0.5RGO-SPS are 98.5% and 97.0%, respectively). The heating rate of SPS was 50 °C min-1. The consolidated 3D GN/Cu bulk composites was 30 mm in diameter and 2.5 mm in thickness. Characterization. Scanning electron microscopy (SEM) (Hitachi S-4800) and TEM (JEOL JEM-2100F) were utilized to observe the microstructures of the powders and composites. High angle Annular Dark Field-Electron energy loss spectroscopy (HAADF-EELS) was done on a FEI Tecnai F20 TEM, to track the structure change of carbon and oxygen elements. Raman spectroscopy (Renishaw inVia Raman Microscope) with 532 nm Ar+ laser was performed to characterize the graphene materials. Thermogravimetry (TG) and differential scanning calorimetry (DSC) analyses (TG-DSC) (TA Perkin-Elmer) were performed up to 800 °C at a heating rate of 5 °C min-1 in Ar to study the transformation process of the in situ synthesized GN@Cu. X-ray photoelectron spectroscopic (XPS) measurement of the GN@Cu, GO powders and GN/Cu, RGO/Cu bulk

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composites were carried out on a PHI 1600 ESCA system. Electron Backscattered Diffraction (EBSD) analysis was carried out by using a HKL Channel 5 system attached to SEM (FEI Nova NanoSEM 430) to evaluate the average grain size and the Taylor factor distribution of the composites. For tensile testing, the obtained bulk samples were cut and polished to a dog-bone shape with a gauge length of 10 mm, a gauge width of 3 mm and a thickness of 2 mm (the illustration is shown in Figure S17). For compressive testing, the samples had a cylindrical disc shape with 2 mm in height and 1.5 mm in diameter. Tensile and compressive testing experiments were performed by a standard mechanical tester (Lloyd (AMETEK) EZ 20) with a crosshead speed of 0.5 mm min-1 and 0.2 mm min-1 at room temperature, respectively. In-situ TEM tensile tests were conducted by single tilt strain holder (Gatan 654), in TEM system (JEOL JEM-2100) with accelerating voltage of 200 kV, at a displacement rate of ~1 μm s-1. The samples were machined to have a dimension of 6.0 mm in gauge length, 2.5 mm in width, and 60 μm in thickness, mechanically polished and then thinned by ion milling. Density Functional Theory (DFT) Calculation Details. To study the influence of the matrix-side interfacial oxygen on the bonding between graphene and copper, we conducted DFT calculation using the projector augmented wave (PAW) formalism

30, 31

of density

functional theory as implemented in the Vienna ab initio simulation package. 32, 33 A kinetic energy cutoff of 400 eV was used with a plane-wave basis set. In order to model the structure of single-layered graphene and Cu matrix with interfacial oxygen, we chose a 2×2 hexagonal supercell of copper (111) of six layers (The copper atoms in the bottom three layers were fixed in XYZ directions) and graphene of one layer as the simulation cell. In the direction perpendicular to the copper, a vacuum region of 20 Å was used to eliminate

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the interactions between the neighboring images. Spin polarization is considered and the exchange–correlation interactions between electrons were treated within the local (spin) density approximation (LDA). The stable geometries of graphene and Cu matrix with one or two interfacial oxygen atoms or without oxygen were obtained firstly by allowing the atomic positions to vary. Then, the models were optimized until the maximum force acting on each atom converged to 0.01 eV/Å or less. The Brillouin zone was sampled using centered scheme with 3×3×1 k-point grid for the structural optimization. The tetrahedron method together with Blöchl corrections 34 was used, and the convergence criterion was set to 10-5 eV. 

RESULTS Microstructure of the Two Types of Reinforcement and Their Composite Powders.

The morphology of the in-situ synthesized GN anchored with Cu NPs was firstly explored via SEM and TEM. As shown in Figure 1a, it exhibits a microscopically porous network of GN which perfectly replicates the flat surface of the assembled NaCl templates. The porous network feature which is actually made up of interconnected pores with the size of 0.5-1 μm, could be identified from the higher magnified SEM image in Figure 1b. Besides, the wall of GN is observed to be homogenously anchored with small-sized Cu NPs. TEM characterization in Figure 1c is in good agreement with the SEM analysis results. The Cu NPs with the measured size of 30-100 nm from Figure 1d retain a good stability on the wall surface of GNs even during a severe TEM specimen preparation process, indicating the robust contact between GN and Cu NPs. The HRTEM image in Figure 1g of the edge area from Figure 1c confirms the thickness of the walls of GN to be about 1~3 nm (3~8 layers). After a coating of copper by an impregnation-reduction process, the GN was thoroughly 10


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encapsulated by big-sized Cu particles of 50-150 nm, verified by both the SEM and TEM images from Figure 1e and 1f. The in-situ decoration of Cu NPs on the GN wall surface facilitates the adhesion of Cu2+ in the impregnation step and furthermore promote the nucleation of Cu particles during the reduction process. Characterized with a large specific surface area, i.e. the lateral plane size is about 500 nm on average and the sheet thickness of 1-5 nm respectively measured from the SEM images in Figure 1h-i, the obtained GO nanosheets used in this work have got a geometric size matched well with that of GN nanosheets. Thanks to the abundant OCFGs on the surface of GO naonosheets, the chemical bonds formed between OCFGs and the impregnated Cu ions promotes the combination of RGO and Cu particles during the subsequent reduction process. As a result, the Cu coated on the RGO surface display an even distribution but a non-uniform size from 40 nm to 300 nm (Figure 1k). So in this way, the homogeneous coating of Cu NPs on the surface of graphene by impregnation-reduction process can ensure the uniform distribution of GN and RGO in the Cu matrix composites fabricated by subsequent SPS consolidation.

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Figure 1. (a, b) SEM and (c, d) TEM images of the GN@Cu at different magnifications; (e) SEM and (f) TEM image of the impregnated-reduced GN@Cu@Cu hybrid powders. The dashed circle highlighted the exposed location of GN in the GN@Cu@Cu. (g) HRTEM image of the edge area of the GN in c; (h) SEM image and (i) TEM image of GO; (j) HRTEM image of the edge area of the GO in i; (k) SEM image of the impregnated-reduced RGO@Cu hybrid powders. Microstructure of the Bulk Composites. Figure S1 demonstrates the typical microstructure of the SPSed 0.5 vol. % GN/Cu and 0.5 vol. % RGO/Cu bulk composites. It could be seen from Figure S1a and d that, both the two types of graphene/Cu composites have fine-grained equiaxed microstructures. The HRTEM images in Figure S1b-c present the detailed information about the distribution of graphene. For 0.5GN-SPS, the in-situ synthesized GN wrap the outer contour of the Cu grains, in which way they could prevent 12


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the contact of the nearby Cu grains and refine the grain size effectively. While for the SPSed RGO/Cu, the RGO nanosheets in the composites were discovered with a separated island-like microstructure (Figure S1e and f). In both cases, the graphene reinforcement has a similar intergranular distribution type. In addition, we validated through a designed experiment that the in-situ synthesized Cu NPs on the surface of GN in GN@Cu has better wettability between graphene and the Cu matrix compared with RGO (as discussed in Figure S2 of supporting materials). EBSD tests were further carried out to find out the difference in grain size and its distribution characteristics of the two types of composite samples (Figure S3a and b, the corresponding distribution analysis shown in Figure S3c and d). As can be seen, both of the samples display the typical characteristics of equiaxed grains, and the statistical result represents that the average grain sizes of the GN/Cu (0.753 μm) are smaller than RGO/Cu (1.125 μm). Thanks to the lower sintering temperature, more rapid heating rate and shorter holding time of SPS method compared with the traditional sintering routes such as hot pressing, the tendency of grain growth in the composites was effectively inhabited, thereby giving rise to a much refined microstructure. The distribution of grain boundary (GB) misorientation in Figure S3e derived from Figure S3a and b exhibits that a slightly larger proportion (97.8%) of high angle GBs (HAGBs, θ≥15 °) and Σ3 twin boundaries (TBs, θ=60°) are identified in GN/Cu bulk composites in comparison with those of RGO/Cu composites (92.1%(SPS)). The large percentage HAGBs and TBs may act as barriers for dislocation slip and thus increase the strain hardening capability of the materials. 11 On the whole, the EBSD results above display that the copper matrix composites reinforced by the two different graphene reinforcements with the same volume fraction exhibit a similar

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matrix grain microstructures. XPS spectroscopy is an effective surface-sensitive method to identify the composition and chemical state of materials. With the fact that massive oxygen atoms are introduced on the surface of the GO nanosheets during the chemical exfoliation process, we are particularly interested in the chemical state difference of oxygen between the two graphene reinforcements in the original powder form as well as in the bulk composites. The C1s spectra for GO and GN@Cu in Figure 2a demonstrate the existence of oxygen-containing group in both GO and GN. For the GO nanosheets as the raw materials, the three major components in the C1s spectrum can be identified as C=C or C-C (284.4- 284.5 eV), C-O (286.1 eV) and C=O (288.8 eV). The high relative content of C=O (28.1%) and C-O (5.7%) suggests the existence of abundant oxygen on the GO surface. Here the O 1s spectra were also applied to track the evolution of the oxygen. It should be pointed out that the O 1s spectra we obtained are surface-focused due to the limitation of the much lower sampling depth of O 1s than the relevant C 1s spectra. 35 According to Figure 2b, the O 1s spectrum of GO can be deconvoluted into: C=O (531.8 eV), C-O (532.6 eV). The detected higher content of C=O than C-O could be attributed to the fact that the O 1s data obtained incorporates parts of the contribution from trapped CO2 molecules. As for the bulk composites, it can be seen that the C1s spectrum of the 0.5RGO-SPS presents a remarkable reduction of the peak intensities of C=O and C-O (relative content of 2.6% and 14.3% respectively) which indicates the effective transformation into RGO during reduction and a good structure maintaining after SPS (see Figure 2c). While significant difference could be easily identified from the C1s and O1s spectra of GN@Cu and 0.5GN-SPS. There is only a small C-O peak (relative content of 9.2%) in the C1s spectrum of GN@Cu and no evidence

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of C=O peak. The relatively low oxygen level of GN is due to the catalytic effect of the Cu NPs anchored on the surface, which contributes to a good graphitization during the high-temperature calcination process.

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While there is no big change in the relative

intensity of C1s spectrum for 0.5GN-SPS with a C-O relative content of 3.5%, suggesting that the composition of GN remains nearly the same during the multi-step processing. On the basis of above evidence, it can be inferred that the in-situ catalytic synthesis process of GN@Cu effectively eliminates most of the OCFCs in the solid carbon source. In this case, the surface of GN is decorated with small percent of C-O bonds, which could facilitate the formation of oxygen mediated carbon/Cu bonding (C-O-Cu) and thereby improve the interfacial bonding strength between Cu and graphene. 9, 22 The O1s XPS spectra could also be adopted to track the oxidation of Cu as presented in Figure 2b and d. For the GN@Cu powders, the peak corresponding to CuOx (530.2 eV) could be found due to the possible oxidation of the small-sized Cu NPs on the GN surface (see Figure 2b). Quite unexpectedly, the conspicuous CuOx peak was only found in the O1s of 0.5RGO-SPS but manifest no obvious signal in 0.5GN-SPS as displayed in Figure 2d. The identification of copper oxides in the RGO/Cu composites prompts us to make a further step into the micro-characterization of the bulk materials. Figure 2e-h represent the STEM/EDXS line-scan results of the interfacial area for 0.5RGO-SPS and 0.5GN-SPS. As can be concluded from Figure 2f and h, for the 0.5RGO-SPS sample, the O level is relatively higher than that of 0.5GN-SPS. Besides, the O element distribution for 0.5RGO-SPS has an up-stream trend close to the copper matrix while there is no significant change for that of 0.5GN-SPS. The evidence above suggests that the CuOx characterized in 0.5RGO-SPS is not the outcome from the reaction between copper powders and the

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exotically-adsorbed oxygen on the surface or the inevitable and native oxygen inside Cu powders as no CuOx was found in 0.5GN-SPS. To sum up, all the evidences above suggest a different interfacial composition between the two types of graphene/copper composites due to the different content and existence form of oxygen.

Figure 2. C1s XPS spectrum of (a) GO and GN@Cu composite powder and (c) 0.5RGO-SPS and 0.5GN-SPS samples; O1s XPS of (b) GO and GN@Cu composite powder and (d) 0.5RGO-SPS and 0.5GN-SPS samples; STEM image and the corresponding EDXS line scan spetra of Cu, C and O elements of (e, f) 0.5RGO-SPS and (g, h) 0.5GN-SPS. ◆ is the start point and ● is the end point of EDXS line scan.

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Interface Structure. The detailed interface structure analyses of the RGO/Cu and GN/Cu bulk composites fabricated by SPS were carried out by HRTEM (Figure 3 and 4). Figure 3c and d present the HRTEM images of interface structure of the magnified selected areas I and Ⅱ of RGO/Cu composites from Figure 3a. Both of the images exhibit seamless interface without voids and inclusions. The few-layered graphene is identified with a lattice fringe with a 0.34 nm spacing of (002) plane which is distinct from the Cu matrix. The inset of Figure 3c and d shows the fast Fourier transformation (FFT) image of the selected area by which the phase and lattice plane could be confirmed. Unexpectedly, we identified cubic structured Cu and Cu2O clearly for the two cases with different zone axis of [110] and

[001] due to the distinguishable difference in the lattice distance between Cu {1 1 1} (0.208 nm) and Cu2O {1 1 1} (0.243 nm); Cu {0 0 2} (1.804 nm) and Cu2O {0 0 2} (0.212 nm); Cu {2 2 0}(1.277 nm) and Cu2O {2 2 0}(1.486 nm) (Cu: JCPDS 04-0836; Cu2O: JCPDS 34-1354). The orientation relationships of the Cu2O with the Cu matrix in Figure 3c are determined as:







(00 2)Cu2O P(00 2)Cu ,



(111)Cu2O P(111)Cu ; while another special 



orientation is also confirmed in Figure 3d: (110)Cu2O P(110)Cu . The corresponding inverse Fast Fourier transform (IFFT) in Figure 3c and d marks the interface between Cu2O and Cu clearly. Besides, a special 6×7 coincidence site lattice (CSL) is confirmed for both cases, which provides a minimum coincidence misfit of 1.22% between the two phases (illustrated in Figure 3e and f). The result above which manifests that Cu2O on the interface have a ″cube-on-cube epitaxy″ relationship with the Cu matrix is a good evidence of the possible oxidation of the Cu matrix by the unreduced oxygen-containing functional groups on the surface of RGO during processing.

36, 37

In order to verify this hypothesis, we performed 17



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EELS spectrum line scan to study the structure and composition change of the C and O elements across the interface in Figure 3g (the magnified selected areas Ⅲ in Figure 3a). We take a series of 9 equally-spaced EELS spectra of which the positions are marked in the HAADF image in Figure 3h. The high-resolution EELS spectrum from carbon K-edge regions in Figure 3i demonstrates two peaks appearing at ~285 eV corresponding to transitions from 1s to the π* state of the sp2-bonded atoms and at ~291 eV related to transitions from 1s to σ* state at both sp2- and sp3-bonded atoms.

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The fine structure of

the C K-edge displays a notable change by changing the position across the interface. At the beginning site near the Cu matrix, the π* peak is hardly seen and the σ* peak has a broad and rounded shape which is correspond to disordered amorphous carbon; with the position moving towards the inner area of RGO, both the π* and σ* become sharper, showing the character of graphene with good crystallinity. The fine structure of O K-edge in Figure 3i derived from the same EELS spectrum of the marked position has a distinct peak at 538 eV and an extended broad peak at 560 eV. The two peaks of O K-edge become weaker with the position approaching the RGO and are almost invisible to the end. The opposite changing trend of the characteristic peaks of O K-edge with C K-edge illustrates that the oxygen atoms associated with the functional groups on the surface of RGO might have a diffusion process towards the interface between RGO and Cu matrix and thus finally move into the lattice of the front-end copper. This could explain the occurrence of well lattice-matched Cu2O on the interface. The interface structure of GN/Cu composites consolidated by the same SPS method exhibits a different result. From Figure 4b and c of the magnified selected areas (A and B) of GN/Cu composites in Figure 4a, a clean and tightly bonded interface could be observed

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between GN and Cu matrix. The selected area electron diffraction (SAED) patterns in Figure 4d associated with the interfacial area in Figure 4c confirms the polycrystalline structural feature of GN. The IFFT images of the front-end interface, indicated by white arrows respectively in the inset of Figure 4b and c, provide us with strong evidence that the interface structure of GN/Cu are composed of GN and pure Cu. At the meantime, a special 

orientation relationship of (002)GN P(111)Cu could be obtained by measuring the lattice distance. In this case, however, the Cu2O in the front-end of the interface existed in the RGO/Cu was not characterized in GN/Cu composites, which is supposed to open a window for us to comparatively study the difference in the interface structure between the GN/Cu and RGO/Cu.

19



Figure 3. (a) TEM image of 0.5RGO-SPS bulk composites; (b) The schematic of the location of the oxygen diffusion zones in (a); (c) and (d) are the HRTEM images of the selected area I and Ⅱ in (a); Insets of A and B are the magnified and the IFFT images of the areas marked with red box, respectively; (e) and (f) are the corresponded schematics of the

"cube-on-cube epitaxy" relationship between Cu and Cu2O in the axis zones of [110]

and [001] . The picture shows a special 6×7 coincidence site lattice (CSL) for both cases; (g) the bright-field STEM image and (h) the corresponding HAADF image of the RGO/Cu interface, the arrow in yellow indicates the direction of the EELS spectrum line scan; The 20


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fine structure of (i) C K-edge and (j) O K-edge measured from the EELS spectrum line scan in (h).

Figure 4. (a) TEM image of 0.5GN-SPS bulk composites; (b) and (c) are the HRTEM images of the selected area of A and B in (a), showing the interface structure between GN and Cu. The insets are the IFFT image of the interface indicated by arrows; (d) is the 

corresponding SAED patterns of the region (c) in the axis zone of [110] Cu.

Mechanical Behaviors. The great difference in microstructure results in different mechanical properties. The tensile properties of the GN/Cu and RGO/Cu bulk composites fabricated by SPS are demonstrated in Figure 5a, and the key mechanical data obtained from the tensile tests are summarized into Table 1. For even such a low graphene content of 21



Page 22 of 48

0.5 vol. %, the yield strength (YS) as well as the ultimate tensile strength (UTS) manifest considerable improvement compared with pure Cu, as revealed in Figure 5a. The 0.5GN-SPS composite is shown to have a yield strength of 332±7 MPa, an ultimate tensile strength of 335±11 MPa, ~39 % and ~ 18 % higher than those of the pure Cu matrix; The 0.5RGO-SPS, by contrast, displays a slightly lower strengthening capability but an obviously decreased fractural elongation of only 8%, which is about 1/3 of that of GN/Cu-SPS. The toughness modulus based on the integral area of the engineering stress-strain curve was also calculated to evaluate the toughness of the composites.

39

The

0.5 GN-SPS bulk has a toughness modulus of 79.8±9.7 MJ m-3, which is 163% higher than 0.5RGO-SPS (30.4±5.6 MJ m-3). Relatively speaking, the 0.5RGO-SPS bulk did not show big advantage in strengthening over RGO/Cu in comparison with noticeable improvement in toughening modulus. This may be related with the relatively small volume fraction of 0.5 vol. % used to guarantee a homogenous distribution of reinforcement in the Cu matrix, considering the fact that graphene is very hard to disperse in metal matrix due to their high aspect ratio. Given that the geometric size of the RGO nanosheets is similar to that of GN nanosheets, the small graphene content did not cause much difference in the matrix grain structure of the two types of graphene/Cu composites (as seen in Figure S3) and thus lead to insignificant increase of both yield strength and tensile strength of 0.5GN-SPS than 0.5RGO-SPS. The normalized strain hardening rates (NSHRs) curves are plotted in Figure 5c to compare the strain hardening behaviors of the two samples. The NSHR is defined as =   /  g /  , where

 is the true stress and  is the true strain, of the GN/Cu

composites and the Cu matrix, versus the corresponding true strain. According to the

22


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Considère criterion that   1 corresponds to the plastic instability of the bulk matrix, 40 a conclusion could be easily reached that despite of obtaining a similar  descending trend (   1% ) in the initial stage of tensile deformation, the 0.5GN-SPS sample holds a much longer steady  which is of vital importance to achieve a higher ductility of the matrix compared with 0.5RGO-SPS. This can also be verified from the normalized strain hardening rate vs. true stress curves in the inset of Figure 5c. The compressive behavior of 0.5GN-SPS and 0.5RGO-SPS are also shown in Figure 5b. Because of the good deformation capability of copper matrix, both the bulk composites did not suffer from collapse even after 45% descending amount. However, the compressive strength reveals a trend that is consistent with that of the tensile properties. The compressive yield strength for 0.5GN-SPS is 319±10 MPa, while that for the 0.5RGO-SPS is 274±5 MPa. It can be seen from the above analysis that the GN/Cu composite has a better overall mechanical properties (yield strength and tensile ductility) than those of the RGO/Cu composite. The different strengthening and toughening effects of the two types of graphene/Cu composites should be considered from two aspects: The first is the intrinsic matrix strengthening and toughening related to the matrix microstructure (i.e. grain structure and dislocation density) change affected by the distributed graphene;

39

the second is the

composite strengthening and toughening associated with the different interface structure. Taking the strengthening mechanism as an example, the dislocation density data of the two types of graphene/Cu composites, which is quantitatively measured by the XRD analysis of bulk materials based on a Williamson-Hall method,

41

demonstrate insignificant difference

for both the deformed and unreformed samples (as calculated from the XRD patterns in Figure S4 of supporting materials). It should be pointed out that the different average grain

23



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size (GN/Cu: 0.753 µm, RGO/Cu: 1.125 µm) may result in different grain refinement strengthening effects according to the classic Hall-Petch relationship. The calculation result (as discussed in the Supporting Materials in detail) suggests that GN/Cu has a relatively higher grain refinement strengthening (  H  P ) compared with RGO/Cu, but the absolute values of  H  P are small with little gap between two types of graphene/Cu composites. This suggest that matrix strengthening mechanisms are not supposed to account for the difference in strengthening effects. As for the toughening mechanisms, the intrinsic toughening is mainly relevant to matrix grain deformation. The strain hardening capability, which is used to estimate the grain deformability of the matrix, could be reflected from the measurement of their Vickers hardness before and after tensile tests.

42

Figure 5d presents

the variation of Vickers hardness plotted as a function of the relative distance from the center position along the thickness direction (0~450 μm) of the tensile specimen gauge area before and after tensile deformation (AT). The RGO/Cu exhibits a slightly higher increase in Vickers hardness (15% for 0.5GN-SPSAT and 21% for 0.5RGO-SPSAT), indicating more significant strain hardening during tensile deformation. However, considering the fact that the 0.5GN-SPS demonstrates a better overall performance both in tensile yield strength and fractural ductility, the strain hardening is not the prominent mechanism in explaining the different toughening effects in the mechanical properties of the two composites. On the basis of above analysis, we can speculate that the markedly different interface structure should of first concern in affecting the mechanical behaviors of the two composites.

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Figure 5. (a) Tensile and (b) compressive engineering stress-strain curves of pure Cu, 0.5GN-SPS and 0.5RGO-SPS bulk composites; (c) The normalized hardening rate versus true strain and true stress (inset) curves derived from the stress-strain curves in (a); (d) The Vickers hardness test of the GN/Cu and RGO/Cu samples before and after tensile test.

Table 1. Summary of the tensile mechanical properties of Cu-SPS, 0.5GN-SPS and 0.5RGO-SPS samples. Average Materials

Total YS

UTS

Hardening

(MPa)

(MPa)

Exponent

Grain size

elongation

(μm) Cu-SPS

2.502

(%) 238±5

283±7

0.188±0.020

25


37±7

Toughness -3

(MJ m ) 100.0±21.3


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0.5GN-SPS

0.753

332±7

335±11

0.123±0.013

23±3

79.8±9.7

0.5RGO-SPS

1.125

326±10

334±13

0.073±0.028

8±4

30.4±5.6



DISCUSSIONS The Transformation Process of the In Situ Synthesized GN Reinforcement. It is well

known that copper itself has been proved as a good catalyst for graphene growth. In the widely investigated field of chemical vapor deposition-grown graphene by using gas-state or solid-state carbon source, copper facilitates the rearrangement of pyrolized carbon atoms into sp2 hybridization and eliminates the amorphous carbon effectively in H2 atmosphere. 43-45

Therefore, the clean and robust interface structure between GN and Cu should be

closely associated with the unique in situ catalytic reaction. On the other hand, according to our previous work, 46 during the in-situ synthesis process the 3D interconnected cubic NaCl particles were adopted to offer a 2D-confined space between adjacent NaCl surfaces to achieve the in situ generation of graphene nanosheets. In this way, under high temperature GN and Cu NPs anchored on its surface could reach a stable interface state within the limited interval reaction space between NaCl particles. So for the purpose of figuring out the origin of the clean and robust GN/Cu interface, it is necessary to have a deep and total understanding of the transformation process of the in situ GN reinforcement. Here, a series of thermal analysis experiments were performed to explore the transformation process from hybrid precursors to GN@Cu in detail. Three groups of designed TGA/DSC experiments (Figure 6), namely carbon nanosheets (CN) precursor [C6H12O6+NaCl], Cu/C precursor [C6H12O6+Cu(NO3)2·3H2O] and GN@Cu precursor 26


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[C6H12O6+Cu(NO3)2·3H2O+NaCl], were carried out to discuss the microscopic role of NaCl and Cu NPs during heat-treatment. On the basis of the TGA/DSC experiments of Figure 6, the schematic illustration of the overall transformation process and the molecular level reaction products of three different stages related to the different temperatures was demonstrated in Figure 7. As presented in Figure 6a and Figure 7, during the stage Ⅰ (20 °C to about 400°C) of the heating process, the carbon nanosheets, Cu/C and GN@Cu precursor firstly undergo a dehydration process under the temperature below 100 °C, accompany with slight weight loss in TGA curves and small endothermic peaks at about 80-90 °C, respectively. After that, C6H12O6 starts to melt at 160 °C and decompose at about 180 °C, which is verified by the DTA/DSC result of freeze-dried C6H12O6 (as seen in Figure S5a). For the TGA/DSC curves of Cu/C precursor, the exothermic peak at 188.4 °C could be attributed to the reduction of copper salt.

47

While in TGA/DSC curves of the

GN@Cu precursor, a less obvious peak at about 300 °C to 360°C (designated as 319.1 °C) is related to a reduction process. Different from the heating process of the Cu(NO3)2·H2O+NaCl precursor, the copper ions could be reduced directly at a rather low temperature without decomposition into Cu2O or CuO, of which the reaction starts at 300-350 °C as observed in Figure S5b. To verify this speculations, the sample of CN@Cu-400 °C was prepared and characterized, with the results displayed in Figure 6b and e, Figure S6a and c. The SEM images in Figure 6b and e depict an irregular and collapsed structure of carbon skeleton uniformly decorated with small particles on its surface. The XRD result in Figure S7a confirms the full conversion of Cu2+ to pure Cu at 400 °C. Due to the relatively large thickness of the skeleton walls, the Cu NPs are actually wrapped by the carbon layer with low quality (ID/IG=0.89 verified by the Raman spectrum

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results in Figure S7b) after carbonization (as seen in Figure 6h, S6a and c). In the subsequent higher temperature range of Stage Ⅱ (400 °C-600 °C)，two weak exothermic peaks at 425.5 °C and 605.7 °C for Cu/C precursor and an obvious peak at 529.7 °C for GN@Cu precursor are believed to be associated with the reaction that amorphous carbon starts to be graphitized under the catalytic reaction of its tightly bound Cu NPs.

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In order

to verify this reaction, the sample of GN@Cu-600 °C was also prepared. As can be seen in Figure 6c and f, a regular network-like and transparent carbon skeleton certifies the much thicker walls compared with those of GN@Cu-400 °C while the Cu NPs are of a similar size. The captured TEM images in Figure 6i, S6b and d manifest an obvious reaction trace of carbon nanosheets with discernable graphene layers around catalytic Cu NPs. However, the Raman spectrum results (Figure S7b) suggest that the carbon nanosheets are of relatively low quality at 600 °C calcination with a high ID/IG ratio of 0.80 and a broad 2D peak. While, the GN@Cu-750 °C are characterized with much smaller ID/IG ratio of 0.75 and a sharp 2D peak. The results imply that during the Stage Ⅲ (600 °C-750 °C), the thick-layered carbon nanosheets undergo further catalytic graphitization to few-layered graphene (Figure 6d, g and j). The products obtained from the calcination of carbon nanosheet precursor at 750 °C has a similar network-like structure but a larger wall thickness of more than 4.5 nm (Figure S8), demonstrating the important role of catalytic Cu NPs in eliminating the redundant amorphous carbon. As talked above, the space-confined effect of NaCl particle assemblies and the in-situ catalytic reaction play a key role in achieving the tightly bonded Cu NPs/GN interface. The markedly lagged peaks of Cu reduction and Cu catalytic graphitization reaction of GN@Cu precursor than Cu/C precursor in Figure 6a are supposed to be caused by the space barrier

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and confinement effect of NaCl particles. Without using NaCl, the calcinated Cu/C precursor at 750 °C yields microscale carbon blocks embedded with aggregated Cu NPs of about 20-50 nm (Figure S9), indicating that the assembled NaCl particles in the GN@Cu precursor not only act as hard templates but also provides the reaction space, thereby leading to a steadier catalytic graphitization process and a close contact between Cu NPs and GN upon calcination process. For better understand the merits of this in situ synthesis, the product of Cu NPs/RGO composite powders were also prepared with a similar process to that of the GN@Cu but replacing the carbon source with GO. As can be seen in Figure S10, the large Cu NPs (>200 nm) lying on the surface of RGO nanosheets demonstrate an inconspicuous interaction with the substrate as a result of aggregation. In this case, these large Cu NPs could not catalyze the GO, in which the C atoms are in well-ordered configuration beyond rearrangement via catalysis, to form a tight interface bonding and remove amorphous carbon from RGO. As a result, our in situ strategy is more favorable to obtain graphene nanosheets with high quality and achieve robust interface interaction between graphene and metal NPs, which would be very beneficial for the effective fabrication of graphene/Cu bulk composites.

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Figure 6. (a) TGA/DSC curves of carbon nanosheet precursor [C6H12O6+NaCl], Cu/C precursor

[C6H12O6+Cu(NO3)2·3H2O]

and

GN@Cu

precursor

[C6H12O6+Cu(NO3)2·3H2O+NaCl]; SEM and TEM images of (b, e and h) CN@Cu-400 °C, (c, f and i) GN@Cu-600 °C and (d, g and j) GN@Cu-750 °C.

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Figure 7. Schematic illustration of the transformation process of the in situ synthesized GN anchored with Cu NPs.

Interfacial Reaction Between RGO and Copper Matrix. Unlike the in-situ synthesized GN with high crystallinity and low oxygen content, the RGO nanosheets which are reduced at 400 °C in H2 atmosphere have some amount of oxygen containing functional groups and also large parts of sp3 carbon associated with the surface defects caused by the severe exfoliation process to obtain GO (the ID/IG ratio 0.97, as shown in Figure S11). The identification of Cu2O phase on the interface of RGO/Cu encourages us to take a further step into the interfacial reaction between RGO and copper matrix. According to the thermodynamics data, the activation energy for forming Cu2O is much lower than energy 31



barrier for the oxidation of graphene.

29

This means that once active oxygen atoms are

released from the RGO surface, the oxidation of copper would occur in priority. Considering the unique mechanism of the SPS consolidation, a high level of pulsed DC current (more than 1 000 amperes and a few volts) is sent directly to the powder bed during sintering,

49

therefore it is believed that the high local sintering temperature offers enough

motivation for oxygen atoms in the RGO to get rid of the constraint from RGO and diffuse to the interfacial area to react with copper. As discussed above, compared with the clean interface between Cu and GN, the interfacial reaction product Cu2O is not supposed to be caused by oxidation during powder preparation. Instead, the typical cube-on-cube orientation relationship of Cu2O with copper supports the concept of oxygen diffusion to the matrix lattice (The schematic illustration of the process can be seen in Figure 8a). A good evidence of the reaction between RGO and copper could be observed from one typical HRTEM image in Figure 8c (magnified selected interfacial area in Figure 8b). A distinctive amorphous layer (FFT image in the inset A) in several nanometers width exists between well-crystalized RGO nanosheets (FFT image in the inset B) and the copper matrix. According to the EELS spectrum line scan result in Figure 3g to j, this area contains sp3 hybridized carbon and higher amount of oxygen than the inner side of RGO. The Raman spectrum in Figure S12 also displays that the ID/IG ratio of 0.5RGO-SPS (0.98) is much higher than that of 0.5GN-SPS (0.72), validating that more defects exist in the SPSed RGO/Cu bulk composites. The formation of the amorphous layer is a combined result of the structure damage caused by the high sintering temperature of SPS and the increase of sp3 dangling bond due to the loss of oxygen containing functional groups on the edge and basal plane of RGO.

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Figure 8. (a) Schematic illustration of the interfacial reaction between RGO and copper matrix during SPS process. (b) TEM image of a typical microstructure of 0.5RGO/Cu with RGO distributed in the grain boundary; (c) The magnified HRTEM image of the marked area in (a), showing the transitional amorphous layer between RGO and Cu matrix. Inset A and B are the corresponding FFT image of area A and B in (c), respectively.

33



Effect of Interface on Mechanical Behaviors of the Composites. In view of the fact that no significant difference has been found for the matrix grain structure, the different interfacial structure is the definitely the primary factor which affects the different deformation and fracture behavior. To qualitatively study the role of the interface in contribution to the mechanical properties of the composites, here we conducted in-situ TEM straining to visualize the interaction between the graphene reinforcement phase and cracks in the dynamic fracture process of the distinctive 0.5RGO-SPS and 0.5GN-SPS composites. Figure 9a-d are the screenshots taken from a typical local fractural area for RGO/Cu composites (the complete process can be watched in video S1). The EDXS spectrum in Figure 9e confirms the existence of RGO nanosheets in Figure 9a. During the in-situ TEM tensile process, crack propagates perpendicular to the tensile direction and almost along the RGO/Cu interface (Figure 9b). As the deformation proceeds, the RGO nanosheets peeled off from the upper part of the interface in a second without displaying obvious interaction with the cracks (Figure 9c, Figure S13). After that, the two parts of the matrix on both sides separated quickly and finally went through final fracture (Figure 9d). The observed phenomenon verified the weak bonding force between RGO and the matrix due to the complex interface structure of Cu/CuOx/amorphous carbon/RGO (illustrated in Figure 9f), which weakens the load transfer and fail to act as effective crack barrier during facture.

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Figure 9. Screenshots of the in situ TEM tensile test of 0.5RGO-SPS composites sample, (a-d) (Video S1†) showing the crack propagation in the RGO/Cu interface; (e) EDXS spectrum of (d); (f) The schematic of the complex interface for RGO/Cu bulk composites.

The clean and robust interface between GN and Cu not only guarantees the effective load transfer, but also possibly contributes to the extrinsic toughening, which is associated with the crack propagation and final facture process. The most widely accepted theory is crack deflection and bridging, which are two effective ways to block the crack tip from moving forward.

19

From the screenshots of the in situ TEM tensile test of 0.5GN-SPS in Figure

10a-d (Video S2), dislocations emitted from the crack tip moves ahead quickly and piled up at the graphene-Cu interface. As the crack came close to GN, a clear pull-out phenomenon was observed which indicates a comparably large interfacial bonding between GN and copper (Figure 10b). The tightly bonded interface acts as obstacle for dislocations and thus improves the resistance of crack propagation in its original direction.

50

As a result, crack

deflects to the pure matrix side (Figure 10d) and in this way more energy is consumed to pass through the GN-Cu interaction zones. The magnified HRTEM image of the GN-Cu area in Figure 10e clearly shows that no voids or gaps exist between two phases. The 35



EDXS characterization (as presented in Figure 10f) validates the existence as well as the relatively intact structure remaining of GN in the pull-out area in Figure 10e. From another sequence of screenshots in Figure S14, we also confirmed the effective crack bridging effect of GN in copper, which could also serve as an effective toughening mechanism to prevent the crack propagation and connection. Due to the limitations of the plane stress condition of the TEM foil specimen, it is impossible to directly observe the three-dimensional crack propagation path and the interaction of the GN and copper. Due to the limitation of equipment and methods, we could not accurately measure the interfacial bonding strengths of the two different interface structure. Based on the obvious different fracture behaviors, however, it is still safe to draw the conclusion that the 0.5GN-SPS has relatively stronger interfacial bonding strength between GN and Cu and this results in superior strengthening and toughening effect to 0.5RGO-SPS excluded difference in matrix grain structure.

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Figure 10. Screenshots of the in situ TEM tensile test of 0.5GN-SPS composite sample, (a-d) (Video S2†) showing the hindrance on crack propagation by crack deflection; (e) The magnified HRTEM image of the selected area in (d); (f) EDXS spectrum of (e); (g) The schematic illustration of the in-situ tensile test process of (a)-(d).

DFT Calculations. The above experimental results have confirmed the dynamic process that oxygen atoms derived from the OFCGs on the graphene surface have potential to

37



diffuse into the Cu matrix lattice when obtaining energy from Joule heat during SPS consolidation. In order to further verify the important role of the interfacial oxygen played in affecting the interfacial bonding strength and thus influencing the mechanical properties of the composites, the computational studies of interfacial interaction between graphene and copper with different interfacial oxygen level using first-principles total energy calculations was performed, with the results presented in Figure 11. For simplicity, the supercell contains single layer graphene with perfect structure and six layers of copper (in the (1 1 1) direction) with substitutional oxygen atoms on the top layer. Considering the fact that the Cu, O ratio in the characterized interfacial product Cu2O is 2:1, here the 1LG-Cu-1O (Cu:O=3:1) and 1LG-Cu-2O (Cu:O=1:1) models were built to study the contribution of the interfacial oxygen level on the bonding between graphene and copper. First of all, the Cu-1O and Cu-2O models without graphene were built and verified with a tendency that oxygen atoms move into the interior of the Cu substrate (Figure S15). Then we performed calculations by optimizing the atomic position based on three different stacking geometry between graphene and the copper below, namely topfcc, hcpfcc and tophcp.

51

For both cases of 1LG-Cu-1O and 1LG-Cu-2O, the copper matrix with

substitutional oxygen moving into interiors attracts the graphene layer moving closer to its side, suggesting an enhanced interaction between oxygen mediated copper and graphene (Figure S16a and b). To quantitatively characterize the interfacial mechanical properties, the binding energy between the graphene layer and copper substrate is calculated as

Eb  EG  M  ( EG  EM ) ,

52

where EC  M , EG and EM are energies of the hybrid systems,

isolated graphene single layer and copper substrate, respectively. As shown in Table S1, the topfcc-1O and hcpfcc-2O cases each have the lowest binding energies of -1.34417 eV and

38


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-1.13559 eV. The universally higher binding energies calculated for 1LG-Cu-1O than those of 1LG-Cu-2O cases reflects a less stabilized state of energy for the graphene-copper system with more oxygen in the matrix-side interface. We next chose topfcc-1O and topfcc-2O to compare the difference of electronic structure of the two cases (the energy minimized structure depicted in Figure 11a and b, Figure 11c and d, respectively). As can be deduced from Figure 11e and g, with the assistance of oxygen, an ionic bond-like feature could be observed between interfacial copper and carbon atoms, suggesting an effective interaction between graphene and copper substrate. The partial density of states (PDOS) analysis of two models shown in Figure 11f and h clearly illustrates the difference in the electronic structures for the two models. The 2p electron states of the oxygen atom of topfcc-1O (Figure 11f, the atomic relative position indicated by Figure 11b) has strong overlap with Cu-3d and C 2p states in the energy interval near the Fermi level and also from -6.5 eV to -5.5 eV, accounting well for the enhanced Eb between graphene and oxygen mediated copper substrate. 25 For the topfcc-2O case in Figure 11h (atomic position illustrated in Figure 11d), there is no obvious hybridization near the Fermi level but only a small overlap between O 2p, Cu 3d and C 2p states in the energy interval from -7 eV to 5.5 eV. The electronic structure analysis also verified the weaker interfacial bonding for 1LG-Cu-2O than that of 1LG-Cu-1O which is probably due to the weakending of decentralized electron interaction between oxygen and the adjacent copper atoms. In this way, we verified the assumption that the increasing oxygen level in the matrix-side interface may be less effective in improving the Cu-graphene interfacial bonding. Experimentally, we validated that 0.5GN-SPS with low oxygen level on the interface which benefits the formation of Cu-O-C bond and thus improve the bonding strength between Cu

39



and GN

53

The topfcc-1O configuration of which the oxygen level is less than the Cu2O

stoichiometry could be applied there to explain the good combination. While for the case of 0.5RGO-Cu, we can comparatively analyze the calculation result of topfcc-2O configuration that, the high level of interfacial oxygen atoms which diffused from RGO weaken the interfacial bonding. Therefore, the DFT calculation result matches well with our experimental results and offers new insights into the interface structure desgin of graphene/metal composites.

Figure 11. Energy minimized structure of (a) the topfcc-1O configuration from the side view, (b) the topfcc-1O configuration from the top view, (c) the topfcc-2O configuration from the side view and (d) the topfcc-2O configuration from the top view; Electron density contours (in unit of e/Å3) of (e) topfcc-1O and (g) topfcc-2O, the black short dot dash line in (b) and (d) marking the separate position of the observed sections; Density of state for (f) topfcc-1O and (h) topfcc-2O. The copper (in the first layer), carbon and oxygen indices correspond to those in (b) and (d), respectively.

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

CONCLUSION Two types of graphene/copper composites, reinforced by in-situ catalytically

synthesized GN with fewer defects and lower oxygen level and RGO nanosheets with more defects and higher oxygen level were successfully fabricated by a modified MLM and SPS method. The interfaces of the GN/Cu and RGO/Cu bulk composite and their effect on the mechanical properties were explored in detail. The measured overall mechanical properties (strength and ductility) of the GN/Cu composites were found to be better than the RGO/Cu composites. Excluded the influence of matrix grain structure, the improved mechanical properties of the GN/Cu composites compared with RGO/Cu composites can be primarily attributed to the effect of interfaces. A complex interfacial structure of Cu/CuOx/amorphous carbon/RGO, which is the result of interfacial reaction between Cu and RGO, was ascertained in the RGO/Cu composite. Moreover, the Cu2O phase as a result of oxygen diffusion into Cu lattice was confirmed in the front site of RGO/Cu interface and has a typical ″cube-on-cube epitaxy″ orientation with Cu. In contrast, the GN/Cu composite demonstrates a clean and robust interface structure consisting of pure Cu and graphene due to the formation of high-quality graphene through in situ catalytic synthesis method, in which process the catalytic effect of Cu NPs facilitates GN to effectively remove the amorphous carbon on the surface and reduce the oxygen content. The strong interface bonding between the GN and Cu matrix results in effective load transfer from the GN to Cu and thus leads to improved strengthening and toughening effect of the GN/Cu composites, which was validated by both in situ TEM tensile experiments and DFT calculation. Our results indicate that the in situ synthesized graphene with fewer defect and oxygen content is more effective for interface control and achieving robust interface bonding in the

41



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graphene-Cu composites.



ASSOCIATED CONTENT

Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI: Microstructure of the bulk composites characterized by TEM, interfacial wettability experiments, EBSD result of the grain microstructure, composites dislocation density calculation, grain refinement strengthening calculation, TGA-DSC plots of

C6H12O6 and

Cu(NO3)2•3H2O+NaCl freeze-dryied precursors, TEM images of 3D CN@Cu-400 °C and 3D GN@Cu-600 °C, XRD profiles and

Raman spectra of CN@Cu-400 °C,

GN@Cu-600 °C and GN@Cu-750 °C, TEM and HRTEM images of 3D CN -750 °C without copper, SEM and TEM images of the products of the calcinated Cu/C precursor at 750 °C, SEM and TEM images of Cu/RGO nanocomposites fabricated by changing the solid carbon source to GO nanosheets, the Raman spectra of GO, 0.5RGO-SPS and 0.5GN-SPS, the screen shot of the in situ TEM tensile test of 0.5RGO-SPS, screenshots of the in situ TEM tensile test of 0.5GN-SPS composites, energy minimized structure of Cu-1O and Cu-2O configurations, energy minimized structure of the topfcc-1O, tophcp-1O, hcpfcc-1O, topfcc-2O, tophcp-2O and hcpfcc-2O configurations and the illustration of the geometric shape and size of the tensile dog-bone specimen (PDF)



AUTHOR INFORMATION

Corresponding Authors

42


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*E-mail: [email protected] (C.N. He). *E-mail: [email protected] (N.Q. Zhao) ORCID Xiang Zhang: 0000-0002-5779-234X Chunsheng Shi: 0000-0003-3374-1449 Enzuo Liu: 0000-0002-3331-2532 Naiqin Zhao: 0000-0002-3725-0279 Chunnian He: 0000-0002-5768-6598

Notes The authors declare no competing financial interest.



ACKNOWLEDGMENT The authors gratefully acknowledge the financial support by the National Natural

Science Funds for Excellent Young Scholar (Grant No. 51422104), the National Natural Science Foundation of China (Grant No. 51531004, 51771130 and 51472177), the Tianjin youth talent support program, the Tianjin Natural Science Funds for Distinguished Young (Grant No. 17JCJQJC44300) and the Tianjin science and technology support project (Grant No. 17ZXCLGX00060).



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Enhanced

Mechanical

Properties

of

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Manocomposites. Acta Mater. 2008, 4, 1936-1940.

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