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reduced graphene oxide reinforced copper matrix composites through spark plasma sintering: An investigation of microstructure and mechanical properties
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Preparation of Ag/reduced graphene oxide reinforced copper matrix composites through spark plasma sintering: An investigation of microstructure and mechanical properties Reza sayyada, Mohammad Ghambaria,∗, Touradj Ebadzadehb, Amir Hossein Paksereshtc,d, Ehsan Ghasalib a
School of Metallurgy and Materials Eng., University College of Eng., University of Tehran, Tehran, Iran Ceramic Dept, Materials and Energy Research Center, Alborz, Iran Department of Mining and Metallurgical Engineering, Amirkabir University of Technology, Tehran, Iran d Dpt. of Coating Processes, FunGlass, Alexander Dubček University of Trenčín, Študentská 2, 911 50, Trenčín, Slovakia b c
A R T I C LE I N FO
A B S T R A C T
Keywords: Copper matrix composite Reduced graphene oxide Chemical synthesis Spark plasma sintering Mechanical properties
The reduced graphene oxide (rGO) decorated with Ag nanoparticles was synthesized by the chemical reduction of graphene oxide in an aqueous solution containing AgNO3, in the presence of hydrazine hydrate as a reducing agent. The reduction of graphene oxide was confirmed by FT-IR and raman spectroscopy analyses. The x-ray diffraction pattern and UV–visible investigations demonstrated the formation of Ag particles on the surface of rGO sheets. After successful decoration, the Ag/rGO nano-composite was used as the reinforcement in the copper matrix composite. Cu–Ag/rGO composites with different percentages of Ag/rGO (0.4, 0.8, 1.6 and 3.2 vol%.) were prepared by mechanical milling and spark plasma sintering (SPS). The effects of the Ag/rGO content on the consolidation process, micro-hardness, bending strength and also, fracture surface of the prepared samples were then investigated. The three-point bending strength of the sintered samples was increased from 285 to 472 MPa by the addition 0.8 vol%. of Ag/rGO, as compared to the pure Cu. Moreover, increasing the reinforcement content to the 3.2 vol%. Ag/rGO led to decreasing the bending strength to 433 MPa. The highest micro-hardness (81 Hv) was obtained for the composite sample containing the 1.6 vol%. Ag/rGO. By increasing Ag/r-GO as the reinforcement (3.2 vol%.), the Vickers hardness was decreased to 69 Hv. Also, investigation of the fracture surface morphology showed transformation of fracture mechanism from plastic changes to brittle ones by raising the Ag/rGO content volume from 0.8 to 1.6 vol %.
1. Introduction Copper and its alloys have been widely used in many branches of engineering applications due to such unique properties as high electrical and thermal conductivity and also, good corrosion resistance. These properties make copper and its alloys suitable candidates for use in electrical sliding contacts, resistance welding electrodes, heat sinks, etc, [1–3]. In spite of the above-mentioned advantages, copper has the main problem of low strength in comparison to other structural engineering metal alloys and composites such as Fe and Ti [4–7]. Therefore, numerous investigations have been devoted to enhancing the mechanical properties of copper with the minimum decrease in electrical conductivity, as the main properties of this metal, through making new alloys or composites [8–10].
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Graphene and CNTs have gained more attention among other reinforcements in copper metal matrix composites, such as SiC [11], Si3N4 [12], TiC [13], TiB2 [14], B4C [15], etc, due to the low loss of electrical conductivity and proper mechanical properties of these carbon-based components [16–18]. Graphene is a carbon allotrope with sp2 hybrid. It consists of a layerto- layer structure from carbon atoms. In the recent years, graphene has been considerably noted by researchers due to such unique properties as the Young's modulus of 1 TPa, and remarkable electrical and thermal conductivity (≈5000 W mk−1). Graphene can be considered as a suitable replacement for carbon nanotubes (CNTs) as the reinforcement in metal matrix composites due to its surface area (2630 m2g-1) and 2-D structure, as compared to CNTs (1315 m2g-1 surface area and 1-D structure). Moreover, using graphene instead of ceramic
Corresponding author. E-mail address: [email protected] (M. Ghambari).
https://doi.org/10.1016/j.ceramint.2020.02.142 Received 24 January 2020; Received in revised form 13 February 2020; Accepted 15 February 2020 0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Reza sayyad, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2020.02.142
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Fig. 1. Experimental procedure of GO reduction to rGO.
hydrate as the reducer factor; (2) faster reaction and time-saving, as compared to the two-post synthesis methods. Moreover, the effect of reinforcement contents on the mechanical properties and microstructure of Cu– Ag/rGO composites prepared via spark plasma sintering could be investigated.
reinforcements, such as Al2O3, SiC, etc., does not seriously reduce electrical and thermal conductivity of the composites, as compared to the pure copper [19–21]. Although extensive studies have been carried out to improve the mechanical properties of metallic materials with addition of graphene, the homogenous distribution of graphene in the metal matrix is still a major challenge in the preparation of metal-graphene composites. This is since graphene nanosheets (GNPs) tend to agglomerate due to their high surface area and van der Waals forces between GNPs. On the other hand, the weak interfacial bond between copper and graphene, which is due to their different thermal expansion coefficients, causes a significant reduction in the mechanical properties of Cu-based composites. To tackle the difficulty of GNPs agglomeration in metal matrix composites, several different methods have been proposed; these include combined liquid state ultrasonic processing and solid state stirring, flake powder metallurgy and molecular-level mixing. The uniform distribution of GNPs as a result of applying the mentioned method leads to improved mechanical properties [22–24]. The molecular-level mixing is a new method preventing the agglomeration of GNPs synthesized on metallic nanoparticles in metal matrix composites (MMCs). One of the most important benefits of metallic nanoparticles attached to GNPs is that a space obstacle is crated, thereby preventing GNPs from joining. Additionally, physical and mechanical properties of MMCs can be improved by correctly selecting metallic nanoparticles located on GNPs, helping to establish strong bonding in the metal-graphene interface. In the present work, Ag- rGO (reduced graphene oxide) nano-composite was synthesized by some simple chemical reduction. The employed method offered two benefits: (1) Ag-rGO nanocomposite could be easily provided in one post by the reaction between AgNO3 and graphene oxide (GO) under an aqueous solution containing hydrazine
2. Experimental procedures 2.1. Synthesis of Ag-rGO nanocomposite Ag- rGO nanocomposite was synthesized by the chemical reduction of silver nitrate and graphene oxide in an aqueous solution, in the presence of hydrazine hydrate as the reducing agent. For this purpose, 100 mg of GO was dispersed in 200 ml of deionized water by ultrasonication for 40 min in order to achieve a stable colloid of GNPs. The obtained colloid of GNPs was stirred for 1h; then 0.1 g of the AgNO3 powder was added under continuous stirring for 1h to form a homogeneous mixture. The hydrazine hydrate aqueous solution (1 μl for 2 mg) was added dropwise into the solution obtained from the previous step under stirring at the temperature of 85 °C for 3 h. The solution was washed 5 times with deionized water using a centrifuge (6000 rpm) for 10 min; then the formed sediment was dispersed in 30 ml of deionized water and stored at the temperature of −40 °C. The final product was dried at −40 °C in the freeze dryer for 12 h. All the above steps were used to rGO synthesis from GO and only the AgNO3 salt addition step was removed from the process. Fig. 1 exhibits the conversion process of GO to rGO. 2.2. Characterization of Ag-rGO composite Field-emission scanning electron microscope (FESEM, MIRA 3 2
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Fig. 2. Scheme of powder mixture preparation and the sintering process by spark plasma.
Fig. 3. XRD patterns of GO, rGO and Ag-rGO samples.
rGO nanocomposite were recorded by a UV–visible spectrophotometer (PerkinElmer Lambda 25, USA). The functional groups of GO and AgrGO samples were characterized by Fourier transform infrared spectroscopy (Bruker Vector 33). Raman spectra of GO and Ag-RGO samples
TESCAN, Czech Republic) was used to observe the morphology of the synthesized sample. The phase analysis of the sample was performed by X-ray diffraction (Philips-PW3710) in the 2θ range of 5–80° with Cu Kα radiation (λ = 1.54 A°) at 40 kV. The UV–visible spectra of GO and Ag3
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rGO nanocomposite in the previous stage were used as the starting materials. Different contents of Ag-rGO (0.4, 0.8, 1.6, 3.2 vol%) were selected for mixing with the copper powder. The initial powders were mixed in a planetary ball mill for 5h at 300 rpm with the ball/powder ratio of 10:1 by using different sizes of balls under Ar atmosphere and 0.5% stearic acid in order to prevent the agglomeration of the powder particles and adhesion of them to the inner walls of the cup. The milled powders were inserted into a graphite mold with the diameter of 32 mm. The graphite foil with 1 mm thickness was also used to prevent the attachment of the final part to the mold walls and upper and lower punches. The samples were sintered by SPS with the initial and final pressure of 10 and 35 MPa, respectively, under vacuum, at 800 °C (600 °C for the external wall of the graphite die). After the sintering process, the samples were extracted from the mold and the remaining graphite foil was removed from the surface of the samples by grinding. Fig. 2 shows different steps related to the preparation and sintering of the samples. 2.4. Characterization of the sintered samples Fig. 4. UV–Vis spectra of GO, rGO and Ag- rGO.
The density of the sintered samples was measured using the Archimedes’ principle. The phase composition of the samples was determined using X-ray diffraction (Philips-PW3710) in the 2θ range of 5–80° with Cu Kα radiation (λ = 1.54 A°) at 40 kV. Micro-hardness test was carried out on the polished surfaces of the samples using the Vickers microhardness tester (MVK-H21, Akashi Co) with a load of 100 g for 10 s. The polished samples were cut in 25 × 5 × 5 dimensions; then the 3-point bending flexural test was conducted using a Santam-STm 20 machine. The morphology and fracture surface of the samples were investigated by SEM (Stereoscan 360, Leica Cambridge) equipped with MAP elemental analysis. 3. Results and discussion Fig. 3 shows the X-ray diffraction (XRD) patterns of GO, rGO and Ag- rGO nanocomposites. As can be observed in Fig. 3a, a peak that appeared at 2θ = 11° that belonged to the (002) diffraction plane of GO. As shown in Fig. 3b, a broad peak that appeared at 2θ = 24.2° could be attributed to the (002) diffraction plane of rGO, implying the removal of oxy-containing functional group and reduction of GO [25,26]. Fig. 3c represents the diffraction peaks obtained at 38.1°, 44.3° and 64.5°, which determined (111), (200) and (220) crystallographic planes of Ag nanoparticles with a face center cubic (fcc) crystal structure in the Ag-rGO nanocomposite, respectively [JCPDS no.00-0040783]. By considering the low weight content of Ag nanoparticles in the Ag-rGO nanocomposite, it was expected that at 2θ = 24.2°, the (002) plane diffraction of rGO would appear; however, no diffraction pattern of rGO was found in Fig. 3c. These results, therefore, suggested that with the addition of hydrazine hydrate in the Ag- rGO synthesis process, the uniform distribution of Ag nanoparticles occurred on the surface of rGO and the rGO nanosheets exited from a stacking state to an exfoliation one. Some recent studies have shown that the diffraction peaks of rGO or GO become weak or even disappear when a layer-to-layer and regular structure of rGO nano-sheets is destroyed through exfoliation [27,28]. It has also been approved that Ag nano-particles attached to rGO nanosheets can prevent the agglomeration of rGO nanosheets and then remove the characteristic peaks of rGO from the XRD pattern. Fig. 4 shows the UV–Vis spectra of GO, rGO and Ag-rGO. In the UV–Vis spectrum of the GO sample, the two characteristic absorption peaks at 236 and 307 nm could be assigned to the π→π* transition of aromatic C–C bond in the GO plane and the n→π* transition of the C═O bond, respectively. The absorption spectrum of the rGO sample demonstrated only a peak at 277 nm, corresponding to the π→π* transition of the aromatic C–C bond [29–31]. The reduction of GO to rGO was accompanied by the retrieval of the electronic conjugation of graphene sheets through the removal of oxygen-containing functional
Fig. 5. FT-IR spectra of GO and Ag- rGO.
Fig. 6. Raman spectra of GO and Ag/rGO.
were obtained by a Raman spectrometer (Takram P50C0R10) with the Nd:YAG laser excitation of 532 nm.
2.3. Consolidation of the Cu/Ag-rGO composite powder The powder samples were sintered by spark plasma sintering (SPS, 20 T-10, Easy Fashion Metal Products Co. China). Copper powder with 99.7% purity and a particle size less than 45 μm and the synthesized Ag4
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Fig. 7. FE-SEM images: (a) rGO, and (b, c, d) Ag-rGO.
(1065 cm−1) disappeared in the FT-IR spectrum of the rGO sample and those at 3443 (−OH) and 1619 cm−1 (C–C skeleton vibration) became visible with weaker intensity, as compared to the GO sample [35,36]. The obtained results, therefore, proposed that GO had been successfully converted to rGO. Besides, a peak loss at 3443 in Ag-rGO sample implied the reduction of oxygen-containing groups, demonstrating the strong interaction between Ag nanoparticles and hydroxyl O atoms. It is noteworthy that the presence of Ag nanoparticles attached to graphene sheets after several steps of centrifugation and ultrasonication confirmed the results of the FT-IR analysis [37,38]. Raman spectroscopy is a beneficial tool to characterize carbon structures, especially for the investigation of their order/disorder structure. Fig. 6 shows the Raman spectrum of GO and Ag-rGO. In the GO sample, there were two main peaks of D and G bands at 1350 and at 1600 cm−1, respectively [39,40]. The D band corresponded to disordered sp3-hybrid carbon atoms and G band arose from ordered sp2hybrid carbon atoms. The D and G band (ID/IG) intensity ratio determined the average size sp2-hybride domain and the disorder degree. The ID/IG value was slightly larger for the Ag-rGO sample, as compared to the GO sample (0.91 and 1.14 for GO and Ag-rGO samples, respectively). The GO reduction could be associated with exfoliation of the graphitic plane, producing the graphene sheets with a smaller average size but more numerous ones, as compared to GO. The strong intensity of D band originated from the edge effects of the Ag-rGO sample, producing more exfoliation, as compared to the GO sample [41,42]. FE-SEM images of rGO and the Ag/rGO nanocomposite can be observed in Fig. 7a and b, c, d, respectively. The exfoliation structure and the layer-by-layer structure of the rGO sample could be observed in Fig. 7a. In the next step, Ag nanoparticles were reduced on graphene nanosheets (Fig. 7 b, c, and d). Fig. 7b shows that Ag nanoparticles attached to graphene sheets acted as spacers to separate the graphene sheets and also hindered the agglomeration and stacking of the rGO nanosheets. Fig. 7c and d shows the homogenous distribution of spherical Ag nanoparticles on the graphene sheets. Ag nano-particles had varied size changes between 40 and 55 nm. Fig. 8a–e shows the XRD patterns of copper composite samples with different contents of Ag/rGO after the sintering process. As can be seen, the peaks appeared at 2θ = 43.3, 50.4 and 74, belonging to (111), (200) and (220) crystallographic planes of copper with the fcc crystal structure, respectively. The important point that should be mentioned
Fig. 8. XRD patterns of the prepared composites: a) pure Cu, b) Cu-0.4 vol% Ag/rGO, c) Cu-0.8 vol% Ag/rGO, d) Cu-1.6 vol% Ag/rGO, and e) Cu-3.2 vol% Ag/rGO.
groups. This removed the absorption peak at 307 nm, leading to a redshift in the absorption peak of rGO from 236 to 277 nm. In the Ag-rGO sample, the strong absorption peak at 421 nm could be attributed to the surface plasmon resonance (SPR) of Ag nanoparticles, indicating the formation of Ag nanoparticles. The surface plasmon resonance is the interaction between the incident light and the valence band electron of Ag, leading to the oscillation of electrons proportional to the frequency of the incident light. Furthermore, with the formation of Ag nanoparticles on the rGO sheets surface, the absorption peak corresponding to the π→π* transition of the aromatic C–C bond was gradually blueshifted from 277 to 261 nm. This phenomenon could be attributed to the charge transfer interactions between Ag nanoparticles and graphene sheets [32–34]. Fig. 5 shows the results of FT-IR spectroscopy analysis on the GO and Ag-rGO nanocomposite. For the GO sample, a peak at 3403 cm−1 confirmed the –OH stretching vibration; also, the peaks at 1731, 1619, 1233 and 1065 cm−1 corresponded to the C═O stretching of the COOH groups, C–C skeletal vibration, epoxy C–O stretching vibration and alkoxy C–O groups vibration, respectively. The peaks corresponding to carboxyl (1731 cm−1), epoxy (1233 cm−1) and alkoxy groups 5
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Fig. 9. Displacement-temperature-time curves during spark plasma sintering of: a) Cu–Ag/rGO composite powder, and b-f) Cu, Cu-0.4 vol% Ag/rGO, Cu-0.8 vol% Ag/rGO, Cu-1.6 vol% Ag/rGO and Cu-3.2 vol% Ag/rGO sintered samples.
temperature changes for Cu– Ag/rGO composite samples. For all samples, the substantial changes of punch displacement began around 23–27 min from the start of the sintering process. The three distinct zones were recognizable in punch displacement versus time curves. Zone I included one-step sintering of the sample in which the middle shrinkage occurred by the heating process at the low pressure (10 MPa) of the hydraulic press. The maximum pressure of 35 MPa was applied in zone II, in which a sudden increase of displacement could be seen. The shrinkage was reduced in zone II with the increase of the reinforcement content, which was an effective factor in the reduction of the final obtained density of the sintered samples. In zone III, the samples were
regarding the XRD patterns of the sintered samples was the absence of copper oxide peaks according to the XRD detection limits. It is known that copper has a strong tendency for oxidation; nevertheless, in the final sintered product, the copper oxide formation can be prevented by making the necessary arrangements in the milling process using Ar atmosphere and the suitable vacuum during the SPS process. The absence of the reduced graphene oxide in the XRD patterns of the Cu–Ag/ rGO sintered samples could be related to large differences of mass absorption coefficients of graphene and copper as well as low amounts of reinforcements. Fig. 9 represents the relationship between punch displacement and 6
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Fig. 10. Relative density (a), micro-hardness (b) and bending strength and extension (c) dependence on the changes of the reinforcement content.
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Fig. 11. SEM images of the fracture surface of the prepared composites containing different contents of Ag/rGO: a) pure Cu, b) Cu-0.4 vol% Ag/rGO, c) Cu-0.8 vol% Ag/rGO, d) Cu-1.6 vol% Ag/rGO and e)Cu-3.2 vol% Ag/rGO.
between copper and graphene sheets was reduced by decreasing the amount of Ag nanoparticles on the graphene nano-sheets due to the poor wettability of the graphene nanosheets by copper. Moreover, the agglomeration of graphene nano-sheets was increased with the rise of the graphene content as a result of their large surface area. Difficulties in the diffusion of copper through agglomerated graphene nanosheets could be regarded as the main reason for the decrease the relative density, as confirmed experimentally, by reducing the punch displacement with an increase of the graphene content (Fig. 9). The changes of micro-hardness with the reinforcement amount in the sintered composites (Fig. 10b) showed that the micro-hardness value was raised from 52 to 81 Hv with the increase of the graphene content from 0 to 1.6 vol%, while further increase in the graphene content to 3.2 vol% decreased the micro-hardness to 69 Hv. As revealed in Fig. 10, the use of graphene with the excellent young's modulus of ~1 TPa as the secondary phase in copper matrix composites improved the strength and hardness of the composite. However, further increase in the graphene content caused the reduction of strength and hardness of copper matrix
heated to the maximum pressure of 35 MPa and the temperature of 600 °C (the external wall temperature of the die, as determined by thermocouple), resulting in the low increase of displacement and final densification. After zone III, the samples were held at above-mentioned conditions for 5 min. It is worth noting that the increase in the amount of graphene from 0 to 3.2 vol% was accompanied by a decrease in shrinkage from 9 to 7 mm during spark plasma sintering. This behavior could be explained by the fact that the increase of the Ag/rGO contents led to the agglomeration of graphene nanosheets along matrix grain boundaries, given that the same volume of powders mixture was inserted to the graphite mold. Therefore, the diffusion of copper into agglomerated areas was restricted due to the lack of wettability of the graphene sheets by copper. The relative density changes showed the decrease of relative density by increasing the reinforcement content, according to Fig. 10a. The minimum and maximum relative density was related to the pure copper and Cu-3.2 vol% Ag/rGO composites, which were 99.4 and 95.8%, respectively (Fig. 10a). During the milling process, the bonding 8
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Fig. 12. SEM images and EDS elemental mapping of a) Cu-0.8 vol% Ag/rGO, b) Cu-1.6 vol% Ag/rGO, and c) Cu-3.2 vol% Ag/rGO samples.
the grain boundaries, which can act as an obstacle against the deformation of the composite during loading. One of the important mechanisms of strengthening in composites is the load transfer from the matrix to the reinforcement phase. In copper-graphene composites, the proper load transfer does not occur due to the poor bonding between the copper matrix and graphene. However, previous research [43] shows that the use of graphene oxide instead of graphene can reduce the deboning between copper and graphene by the formation of the Cu–O bond (bonding Cu with oxygen-containing functional groups on the surface of graphene oxide). In the present work, the use of Ag/rGO reinforcement could help to establish the bond between the reinforcement and the matrix by the formation of Cu–Ag bonds, subsequently improving the composite strength by increasing the load transfer from copper to Ag/rGO. Finally, the Orowan's strengthening is another mechanism effective on the strengthening of the metal matrix composites [44–46].
composites due to an increase in the agglomeration probability of graphene, which was accompanied by the deboning of graphene and copper as a result of the low wettability of copper. Fig. 10c exhibits the maximum bending strength of 472 MPa at the graphene volume percentage of 0.8, which was 1.65 times more than that of pure copper. In addition, the elongation of fracture was decreased gradually with increasing the graphene content beyond 0.8 %vol. There are several different mechanisms for strengthening the metal matrix composites by adding graphene. The addition of graphene influences the microstructure of metal matrix composites; the high density of dislocations is created in the vicinity of graphene-metal intersection due to the high thermal expansion mismatch between graphene and the matrix. The increase of dislocations could improve the strength of the metal matrix composites. The reduction of the grain size is another effect of adding graphene on the microstructure of the metal matrix composite. The reduction of grain size is equal to an increase in
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● The change of the fracture mechanism from a plastic one to a brittle one occurred in the Cu-0.8 vol% Ag/rGO sample. The change of the fracture mechanism originated from the increase of the graphene content in the samples and this as accompanied with the aggregation of the graphene sheets.
Fig. 11 shows the fracture surface of the sintered samples containing different volume fractions of graphene decorated with Ag nano-particles. The fracture surface morphology of pure copper and Cu-0.4 vol% Ag/rGO samples revealed the grains with a rounded corner as a typical characteristic of the plastic fracture. The fracture surface of the specimen containing 0.8 vol% Ag/rGO demonstrated a sharper corner of grains. Furthermore, the pulled-out and separated grains from boundaries in the Cu-0.8 vol% Ag/rGO sample (Fig. 11 c) showed the change of fracture mechanism from a plastic one to a brittle one. The fracture surface morphology of Cu- 1.6 vol% Ag/rGO and Cu3.2 vol% Ag/rGO samples exhibited a layer-by-layer structure indicating a brittle fracture. The porosity and layer-by-layer structure marked in Fig. 11 d and e could be attributed to the debonding of the copper-graphene interface as a result of the weak wettability of copper on the agglomerated graphene. The bending strength of Cu- 1.6 vol% Ag/rGO and Cu- 3.2 vol% Ag/rGO samples was low because of the agglomeration of graphene and formation of porosities between the graphene layers, which could be considered as the stress concentration centers and the preferred sites of crack nucleation. Therefore, accumulation of the initial shallow cracks formed larger cracks, ultimately leading to the brittle fracture of Cu- 1.6 vol% Ag/rGO and Cu- 3.2 vol% Ag/rGO samples, as compared to the Cu- 0.8 vol% Ag/rGO specimen. From the above results, it may be concluded that the graphene nanosheets in 0.8 vol% of the Ag/rGO sample did not act as stress concentration centers and even served as a barrier for dislocation motion, thereby helping to improve the bending strength of the composites. Fig. 12 exhibits SEM images along with their elemental map analysis of the sintered composites containing 0.8, 1.6 and 3.2 vol% Ag/rGO, confirming the presence of graphene sheets. Additionally, Fig. 12a shows the rupture area of the graphene sheet in the fracture surface of the copper matrix composite. The tear and string structure of graphene sheets in the fracture cross-section could be attributed to overcoming the bending forces on the bond between Ag nanoparticles on the graphene plate and the copper matrix. This is because, if there were no bond between graphene plate and copper matrix, then the graphene sheets could not be torn off. Moreover, the transparency of the graphene sheet can be observed in Fig. 12b, which is one of the characteristics of graphene plates due to the very small thickness of these plates. The agglomeration of graphene plates in the Cu- 3.2 Vol.%Ag/ rGO sample is clearly observable in Fig. 12c, showing a non-uniform distribution of graphene and confirming the decrease of the bending strength values for this composite.
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4. Conclusions ● Ag/rGO nanocomposite was synthesized by a chemical reduction of graphene oxide in an aqueous solution under hydrazine hydrate as the reducing agent and silver nitrate. The morphology of the synthesized nanocomposite showed the homogenous distribution of Ag nanoparticles with a size in the range of 40–55 nm on the rGO sheets. ● An Ag/rGO nanocomposite with different volume contents was used as the second phase in the copper matrix composite. Cu– Ag/rGO composites were prepared by mechanical milling and spark plasma sintering. ● The samples of the sintered composites had high relative density. The relative density of the composites samples was reduced from 99.4% to 95.8% for the samples of pure copper and Cu- 3.2 vol% Ag/rGO, respectively. ● The micro-hardness and bending strength of the composite samples showed a significant increase, as compared to the pure copper sample. The highest values micro-hardness (81 Hv) and bending strength (472 MPa) belonged to Cu-1.6 vol% Ag/rGO and Cu-0.8 vol % Ag/rGO samples, respectively, while the values of micro-hardness and bending strength of pure copper were obtained to be 51 Hv and 285 MPa, respectively. 10
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graphene nanocomposite, Nano Res. 3 (5) (2010) 339–349. [39] A.C. Ferrari, D.M. Basko, Raman spectroscopy as a versatile tool for studying the properties of graphene, Nat. Nanotechnol. 8 (4) (2013) 235. [40] T.A. Saleh, M.M. Al-Shalalfeh, A.A. Al-Saadi, Silver loaded graphene as a substrate for sensing 2-thiouracil using surface-enhanced Raman scattering, Sensor. Actuator. B Chem. 254 (2018) 1110–1117. [41] E. Dervishi, et al., Large-scale graphene production by RF-cCVD method, Chem. Commun. (27) (2009) 4061–4063. [42] Z. Zhang, et al., A facile one-pot method to high-quality Ag-graphene composite nanosheets for efficient surface-enhanced Raman scattering, Chem. Commun. 47 (22) (2011) 6440–6442. [43] D. Zhang, Z. Zhan, Strengthening effect of graphene derivatives in copper matrix composites, J. Alloys Compd. 654 (2016) 226–233. [44] A. Bisht, et al., Strengthening mechanism in graphene nanoplatelets reinforced aluminum composite fabricated through spark plasma sintering, Mater. Sci. Eng., A 695 (2017) 20–28. [45] W. Kim, T. Lee, S. Han, Multi-layer graphene/copper composites: preparation using high-ratio differential speed rolling, microstructure and mechanical properties, Carbon 69 (2014) 55–65. [46] S. Xiang, et al., Graphene nanoplatelets induced heterogeneous bimodal structural magnesium matrix composites with enhanced mechanical properties, Sci. Rep. 6 (2016) 38824.
[28] C. Xu, X. Wang, J. Zhu, Graphene− metal particle nanocomposites, J. Phys. Chem. C 112 (50) (2008) 19841–19845. [29] R. Pasricha, S. Gupta, A.K. Srivastava, A facile and novel synthesis of Ag–graphene‐based nanocomposites, Small 5 (20) (2009) 2253–2259. [30] K. Yang, et al., Intercalating oleylamines in graphite oxide, Langmuir 28 (5) (2012) 2904–2908. [31] C. Zhu, et al., Reducing sugar: new functional molecules for the green synthesis of graphene nanosheets, ACS Nano 4 (4) (2010) 2429–2437. [32] M.R. Das, et al., Synthesis of silver nanoparticles in an aqueous suspension of graphene oxide sheets and its antimicrobial activity, Colloids Surf. B Biointerfaces 83 (1) (2011) 16–22. [33] S. Kundu, et al., Photochemical formation of electrically conductive silver nanowires on polymer scaffolds, J. Colloid Interface Sci. 344 (2) (2010) 334–342. [34] V.K. Sharma, R.A. Yngard, Y. Lin, Silver nanoparticles: green synthesis and their antimicrobial activities, Adv. Colloid Interface Sci. 145 (1–2) (2009) 83–96. [35] W. Yuan, Y. Gu, L. Li, Green synthesis of graphene/Ag nanocomposites, Appl. Surf. Sci. 261 (2012) 753–758. [36] L. Zheng, et al., Preparation and capacitance performance of Ag–graphene based nanocomposite, J. Power Sources 201 (2012) 376–381. [37] R.V. Hull, et al., Pt nanoparticle binding on functionalized multiwalled carbon nanotubes, Chem. Mater. 18 (7) (2006) 1780–1788. [38] J. Shen, et al., Facile synthesis and application of Ag-chemically converted
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Contents lists available at ScienceDirect
Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Preparation of Ag/reduced graphene oxide reinforced copper matrix composites through spark plasma sintering: An investigation of microstructure and mechanical properties Reza sayyada, Mohammad Ghambaria,∗, Touradj Ebadzadehb, Amir Hossein Paksereshtc,d, Ehsan Ghasalib a
School of Metallurgy and Materials Eng., University College of Eng., University of Tehran, Tehran, Iran Ceramic Dept, Materials and Energy Research Center, Alborz, Iran Department of Mining and Metallurgical Engineering, Amirkabir University of Technology, Tehran, Iran d Dpt. of Coating Processes, FunGlass, Alexander Dubček University of Trenčín, Študentská 2, 911 50, Trenčín, Slovakia b c
A R T I C LE I N FO
A B S T R A C T
Keywords: Copper matrix composite Reduced graphene oxide Chemical synthesis Spark plasma sintering Mechanical properties
The reduced graphene oxide (rGO) decorated with Ag nanoparticles was synthesized by the chemical reduction of graphene oxide in an aqueous solution containing AgNO3, in the presence of hydrazine hydrate as a reducing agent. The reduction of graphene oxide was confirmed by FT-IR and raman spectroscopy analyses. The x-ray diffraction pattern and UV–visible investigations demonstrated the formation of Ag particles on the surface of rGO sheets. After successful decoration, the Ag/rGO nano-composite was used as the reinforcement in the copper matrix composite. Cu–Ag/rGO composites with different percentages of Ag/rGO (0.4, 0.8, 1.6 and 3.2 vol%.) were prepared by mechanical milling and spark plasma sintering (SPS). The effects of the Ag/rGO content on the consolidation process, micro-hardness, bending strength and also, fracture surface of the prepared samples were then investigated. The three-point bending strength of the sintered samples was increased from 285 to 472 MPa by the addition 0.8 vol%. of Ag/rGO, as compared to the pure Cu. Moreover, increasing the reinforcement content to the 3.2 vol%. Ag/rGO led to decreasing the bending strength to 433 MPa. The highest micro-hardness (81 Hv) was obtained for the composite sample containing the 1.6 vol%. Ag/rGO. By increasing Ag/r-GO as the reinforcement (3.2 vol%.), the Vickers hardness was decreased to 69 Hv. Also, investigation of the fracture surface morphology showed transformation of fracture mechanism from plastic changes to brittle ones by raising the Ag/rGO content volume from 0.8 to 1.6 vol %.
1. Introduction Copper and its alloys have been widely used in many branches of engineering applications due to such unique properties as high electrical and thermal conductivity and also, good corrosion resistance. These properties make copper and its alloys suitable candidates for use in electrical sliding contacts, resistance welding electrodes, heat sinks, etc, [1–3]. In spite of the above-mentioned advantages, copper has the main problem of low strength in comparison to other structural engineering metal alloys and composites such as Fe and Ti [4–7]. Therefore, numerous investigations have been devoted to enhancing the mechanical properties of copper with the minimum decrease in electrical conductivity, as the main properties of this metal, through making new alloys or composites [8–10].
∗
Graphene and CNTs have gained more attention among other reinforcements in copper metal matrix composites, such as SiC [11], Si3N4 [12], TiC [13], TiB2 [14], B4C [15], etc, due to the low loss of electrical conductivity and proper mechanical properties of these carbon-based components [16–18]. Graphene is a carbon allotrope with sp2 hybrid. It consists of a layerto- layer structure from carbon atoms. In the recent years, graphene has been considerably noted by researchers due to such unique properties as the Young's modulus of 1 TPa, and remarkable electrical and thermal conductivity (≈5000 W mk−1). Graphene can be considered as a suitable replacement for carbon nanotubes (CNTs) as the reinforcement in metal matrix composites due to its surface area (2630 m2g-1) and 2-D structure, as compared to CNTs (1315 m2g-1 surface area and 1-D structure). Moreover, using graphene instead of ceramic
Corresponding author. E-mail address: [email protected] (M. Ghambari).
https://doi.org/10.1016/j.ceramint.2020.02.142 Received 24 January 2020; Received in revised form 13 February 2020; Accepted 15 February 2020 0272-8842/ © 2020 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: Reza sayyad, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2020.02.142
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Fig. 1. Experimental procedure of GO reduction to rGO.
hydrate as the reducer factor; (2) faster reaction and time-saving, as compared to the two-post synthesis methods. Moreover, the effect of reinforcement contents on the mechanical properties and microstructure of Cu– Ag/rGO composites prepared via spark plasma sintering could be investigated.
reinforcements, such as Al2O3, SiC, etc., does not seriously reduce electrical and thermal conductivity of the composites, as compared to the pure copper [19–21]. Although extensive studies have been carried out to improve the mechanical properties of metallic materials with addition of graphene, the homogenous distribution of graphene in the metal matrix is still a major challenge in the preparation of metal-graphene composites. This is since graphene nanosheets (GNPs) tend to agglomerate due to their high surface area and van der Waals forces between GNPs. On the other hand, the weak interfacial bond between copper and graphene, which is due to their different thermal expansion coefficients, causes a significant reduction in the mechanical properties of Cu-based composites. To tackle the difficulty of GNPs agglomeration in metal matrix composites, several different methods have been proposed; these include combined liquid state ultrasonic processing and solid state stirring, flake powder metallurgy and molecular-level mixing. The uniform distribution of GNPs as a result of applying the mentioned method leads to improved mechanical properties [22–24]. The molecular-level mixing is a new method preventing the agglomeration of GNPs synthesized on metallic nanoparticles in metal matrix composites (MMCs). One of the most important benefits of metallic nanoparticles attached to GNPs is that a space obstacle is crated, thereby preventing GNPs from joining. Additionally, physical and mechanical properties of MMCs can be improved by correctly selecting metallic nanoparticles located on GNPs, helping to establish strong bonding in the metal-graphene interface. In the present work, Ag- rGO (reduced graphene oxide) nano-composite was synthesized by some simple chemical reduction. The employed method offered two benefits: (1) Ag-rGO nanocomposite could be easily provided in one post by the reaction between AgNO3 and graphene oxide (GO) under an aqueous solution containing hydrazine
2. Experimental procedures 2.1. Synthesis of Ag-rGO nanocomposite Ag- rGO nanocomposite was synthesized by the chemical reduction of silver nitrate and graphene oxide in an aqueous solution, in the presence of hydrazine hydrate as the reducing agent. For this purpose, 100 mg of GO was dispersed in 200 ml of deionized water by ultrasonication for 40 min in order to achieve a stable colloid of GNPs. The obtained colloid of GNPs was stirred for 1h; then 0.1 g of the AgNO3 powder was added under continuous stirring for 1h to form a homogeneous mixture. The hydrazine hydrate aqueous solution (1 μl for 2 mg) was added dropwise into the solution obtained from the previous step under stirring at the temperature of 85 °C for 3 h. The solution was washed 5 times with deionized water using a centrifuge (6000 rpm) for 10 min; then the formed sediment was dispersed in 30 ml of deionized water and stored at the temperature of −40 °C. The final product was dried at −40 °C in the freeze dryer for 12 h. All the above steps were used to rGO synthesis from GO and only the AgNO3 salt addition step was removed from the process. Fig. 1 exhibits the conversion process of GO to rGO. 2.2. Characterization of Ag-rGO composite Field-emission scanning electron microscope (FESEM, MIRA 3 2
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Fig. 2. Scheme of powder mixture preparation and the sintering process by spark plasma.
Fig. 3. XRD patterns of GO, rGO and Ag-rGO samples.
rGO nanocomposite were recorded by a UV–visible spectrophotometer (PerkinElmer Lambda 25, USA). The functional groups of GO and AgrGO samples were characterized by Fourier transform infrared spectroscopy (Bruker Vector 33). Raman spectra of GO and Ag-RGO samples
TESCAN, Czech Republic) was used to observe the morphology of the synthesized sample. The phase analysis of the sample was performed by X-ray diffraction (Philips-PW3710) in the 2θ range of 5–80° with Cu Kα radiation (λ = 1.54 A°) at 40 kV. The UV–visible spectra of GO and Ag3
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rGO nanocomposite in the previous stage were used as the starting materials. Different contents of Ag-rGO (0.4, 0.8, 1.6, 3.2 vol%) were selected for mixing with the copper powder. The initial powders were mixed in a planetary ball mill for 5h at 300 rpm with the ball/powder ratio of 10:1 by using different sizes of balls under Ar atmosphere and 0.5% stearic acid in order to prevent the agglomeration of the powder particles and adhesion of them to the inner walls of the cup. The milled powders were inserted into a graphite mold with the diameter of 32 mm. The graphite foil with 1 mm thickness was also used to prevent the attachment of the final part to the mold walls and upper and lower punches. The samples were sintered by SPS with the initial and final pressure of 10 and 35 MPa, respectively, under vacuum, at 800 °C (600 °C for the external wall of the graphite die). After the sintering process, the samples were extracted from the mold and the remaining graphite foil was removed from the surface of the samples by grinding. Fig. 2 shows different steps related to the preparation and sintering of the samples. 2.4. Characterization of the sintered samples Fig. 4. UV–Vis spectra of GO, rGO and Ag- rGO.
The density of the sintered samples was measured using the Archimedes’ principle. The phase composition of the samples was determined using X-ray diffraction (Philips-PW3710) in the 2θ range of 5–80° with Cu Kα radiation (λ = 1.54 A°) at 40 kV. Micro-hardness test was carried out on the polished surfaces of the samples using the Vickers microhardness tester (MVK-H21, Akashi Co) with a load of 100 g for 10 s. The polished samples were cut in 25 × 5 × 5 dimensions; then the 3-point bending flexural test was conducted using a Santam-STm 20 machine. The morphology and fracture surface of the samples were investigated by SEM (Stereoscan 360, Leica Cambridge) equipped with MAP elemental analysis. 3. Results and discussion Fig. 3 shows the X-ray diffraction (XRD) patterns of GO, rGO and Ag- rGO nanocomposites. As can be observed in Fig. 3a, a peak that appeared at 2θ = 11° that belonged to the (002) diffraction plane of GO. As shown in Fig. 3b, a broad peak that appeared at 2θ = 24.2° could be attributed to the (002) diffraction plane of rGO, implying the removal of oxy-containing functional group and reduction of GO [25,26]. Fig. 3c represents the diffraction peaks obtained at 38.1°, 44.3° and 64.5°, which determined (111), (200) and (220) crystallographic planes of Ag nanoparticles with a face center cubic (fcc) crystal structure in the Ag-rGO nanocomposite, respectively [JCPDS no.00-0040783]. By considering the low weight content of Ag nanoparticles in the Ag-rGO nanocomposite, it was expected that at 2θ = 24.2°, the (002) plane diffraction of rGO would appear; however, no diffraction pattern of rGO was found in Fig. 3c. These results, therefore, suggested that with the addition of hydrazine hydrate in the Ag- rGO synthesis process, the uniform distribution of Ag nanoparticles occurred on the surface of rGO and the rGO nanosheets exited from a stacking state to an exfoliation one. Some recent studies have shown that the diffraction peaks of rGO or GO become weak or even disappear when a layer-to-layer and regular structure of rGO nano-sheets is destroyed through exfoliation [27,28]. It has also been approved that Ag nano-particles attached to rGO nanosheets can prevent the agglomeration of rGO nanosheets and then remove the characteristic peaks of rGO from the XRD pattern. Fig. 4 shows the UV–Vis spectra of GO, rGO and Ag-rGO. In the UV–Vis spectrum of the GO sample, the two characteristic absorption peaks at 236 and 307 nm could be assigned to the π→π* transition of aromatic C–C bond in the GO plane and the n→π* transition of the C═O bond, respectively. The absorption spectrum of the rGO sample demonstrated only a peak at 277 nm, corresponding to the π→π* transition of the aromatic C–C bond [29–31]. The reduction of GO to rGO was accompanied by the retrieval of the electronic conjugation of graphene sheets through the removal of oxygen-containing functional
Fig. 5. FT-IR spectra of GO and Ag- rGO.
Fig. 6. Raman spectra of GO and Ag/rGO.
were obtained by a Raman spectrometer (Takram P50C0R10) with the Nd:YAG laser excitation of 532 nm.
2.3. Consolidation of the Cu/Ag-rGO composite powder The powder samples were sintered by spark plasma sintering (SPS, 20 T-10, Easy Fashion Metal Products Co. China). Copper powder with 99.7% purity and a particle size less than 45 μm and the synthesized Ag4
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Fig. 7. FE-SEM images: (a) rGO, and (b, c, d) Ag-rGO.
(1065 cm−1) disappeared in the FT-IR spectrum of the rGO sample and those at 3443 (−OH) and 1619 cm−1 (C–C skeleton vibration) became visible with weaker intensity, as compared to the GO sample [35,36]. The obtained results, therefore, proposed that GO had been successfully converted to rGO. Besides, a peak loss at 3443 in Ag-rGO sample implied the reduction of oxygen-containing groups, demonstrating the strong interaction between Ag nanoparticles and hydroxyl O atoms. It is noteworthy that the presence of Ag nanoparticles attached to graphene sheets after several steps of centrifugation and ultrasonication confirmed the results of the FT-IR analysis [37,38]. Raman spectroscopy is a beneficial tool to characterize carbon structures, especially for the investigation of their order/disorder structure. Fig. 6 shows the Raman spectrum of GO and Ag-rGO. In the GO sample, there were two main peaks of D and G bands at 1350 and at 1600 cm−1, respectively [39,40]. The D band corresponded to disordered sp3-hybrid carbon atoms and G band arose from ordered sp2hybrid carbon atoms. The D and G band (ID/IG) intensity ratio determined the average size sp2-hybride domain and the disorder degree. The ID/IG value was slightly larger for the Ag-rGO sample, as compared to the GO sample (0.91 and 1.14 for GO and Ag-rGO samples, respectively). The GO reduction could be associated with exfoliation of the graphitic plane, producing the graphene sheets with a smaller average size but more numerous ones, as compared to GO. The strong intensity of D band originated from the edge effects of the Ag-rGO sample, producing more exfoliation, as compared to the GO sample [41,42]. FE-SEM images of rGO and the Ag/rGO nanocomposite can be observed in Fig. 7a and b, c, d, respectively. The exfoliation structure and the layer-by-layer structure of the rGO sample could be observed in Fig. 7a. In the next step, Ag nanoparticles were reduced on graphene nanosheets (Fig. 7 b, c, and d). Fig. 7b shows that Ag nanoparticles attached to graphene sheets acted as spacers to separate the graphene sheets and also hindered the agglomeration and stacking of the rGO nanosheets. Fig. 7c and d shows the homogenous distribution of spherical Ag nanoparticles on the graphene sheets. Ag nano-particles had varied size changes between 40 and 55 nm. Fig. 8a–e shows the XRD patterns of copper composite samples with different contents of Ag/rGO after the sintering process. As can be seen, the peaks appeared at 2θ = 43.3, 50.4 and 74, belonging to (111), (200) and (220) crystallographic planes of copper with the fcc crystal structure, respectively. The important point that should be mentioned
Fig. 8. XRD patterns of the prepared composites: a) pure Cu, b) Cu-0.4 vol% Ag/rGO, c) Cu-0.8 vol% Ag/rGO, d) Cu-1.6 vol% Ag/rGO, and e) Cu-3.2 vol% Ag/rGO.
groups. This removed the absorption peak at 307 nm, leading to a redshift in the absorption peak of rGO from 236 to 277 nm. In the Ag-rGO sample, the strong absorption peak at 421 nm could be attributed to the surface plasmon resonance (SPR) of Ag nanoparticles, indicating the formation of Ag nanoparticles. The surface plasmon resonance is the interaction between the incident light and the valence band electron of Ag, leading to the oscillation of electrons proportional to the frequency of the incident light. Furthermore, with the formation of Ag nanoparticles on the rGO sheets surface, the absorption peak corresponding to the π→π* transition of the aromatic C–C bond was gradually blueshifted from 277 to 261 nm. This phenomenon could be attributed to the charge transfer interactions between Ag nanoparticles and graphene sheets [32–34]. Fig. 5 shows the results of FT-IR spectroscopy analysis on the GO and Ag-rGO nanocomposite. For the GO sample, a peak at 3403 cm−1 confirmed the –OH stretching vibration; also, the peaks at 1731, 1619, 1233 and 1065 cm−1 corresponded to the C═O stretching of the COOH groups, C–C skeletal vibration, epoxy C–O stretching vibration and alkoxy C–O groups vibration, respectively. The peaks corresponding to carboxyl (1731 cm−1), epoxy (1233 cm−1) and alkoxy groups 5
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Fig. 9. Displacement-temperature-time curves during spark plasma sintering of: a) Cu–Ag/rGO composite powder, and b-f) Cu, Cu-0.4 vol% Ag/rGO, Cu-0.8 vol% Ag/rGO, Cu-1.6 vol% Ag/rGO and Cu-3.2 vol% Ag/rGO sintered samples.
temperature changes for Cu– Ag/rGO composite samples. For all samples, the substantial changes of punch displacement began around 23–27 min from the start of the sintering process. The three distinct zones were recognizable in punch displacement versus time curves. Zone I included one-step sintering of the sample in which the middle shrinkage occurred by the heating process at the low pressure (10 MPa) of the hydraulic press. The maximum pressure of 35 MPa was applied in zone II, in which a sudden increase of displacement could be seen. The shrinkage was reduced in zone II with the increase of the reinforcement content, which was an effective factor in the reduction of the final obtained density of the sintered samples. In zone III, the samples were
regarding the XRD patterns of the sintered samples was the absence of copper oxide peaks according to the XRD detection limits. It is known that copper has a strong tendency for oxidation; nevertheless, in the final sintered product, the copper oxide formation can be prevented by making the necessary arrangements in the milling process using Ar atmosphere and the suitable vacuum during the SPS process. The absence of the reduced graphene oxide in the XRD patterns of the Cu–Ag/ rGO sintered samples could be related to large differences of mass absorption coefficients of graphene and copper as well as low amounts of reinforcements. Fig. 9 represents the relationship between punch displacement and 6
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Fig. 10. Relative density (a), micro-hardness (b) and bending strength and extension (c) dependence on the changes of the reinforcement content.
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Fig. 11. SEM images of the fracture surface of the prepared composites containing different contents of Ag/rGO: a) pure Cu, b) Cu-0.4 vol% Ag/rGO, c) Cu-0.8 vol% Ag/rGO, d) Cu-1.6 vol% Ag/rGO and e)Cu-3.2 vol% Ag/rGO.
between copper and graphene sheets was reduced by decreasing the amount of Ag nanoparticles on the graphene nano-sheets due to the poor wettability of the graphene nanosheets by copper. Moreover, the agglomeration of graphene nano-sheets was increased with the rise of the graphene content as a result of their large surface area. Difficulties in the diffusion of copper through agglomerated graphene nanosheets could be regarded as the main reason for the decrease the relative density, as confirmed experimentally, by reducing the punch displacement with an increase of the graphene content (Fig. 9). The changes of micro-hardness with the reinforcement amount in the sintered composites (Fig. 10b) showed that the micro-hardness value was raised from 52 to 81 Hv with the increase of the graphene content from 0 to 1.6 vol%, while further increase in the graphene content to 3.2 vol% decreased the micro-hardness to 69 Hv. As revealed in Fig. 10, the use of graphene with the excellent young's modulus of ~1 TPa as the secondary phase in copper matrix composites improved the strength and hardness of the composite. However, further increase in the graphene content caused the reduction of strength and hardness of copper matrix
heated to the maximum pressure of 35 MPa and the temperature of 600 °C (the external wall temperature of the die, as determined by thermocouple), resulting in the low increase of displacement and final densification. After zone III, the samples were held at above-mentioned conditions for 5 min. It is worth noting that the increase in the amount of graphene from 0 to 3.2 vol% was accompanied by a decrease in shrinkage from 9 to 7 mm during spark plasma sintering. This behavior could be explained by the fact that the increase of the Ag/rGO contents led to the agglomeration of graphene nanosheets along matrix grain boundaries, given that the same volume of powders mixture was inserted to the graphite mold. Therefore, the diffusion of copper into agglomerated areas was restricted due to the lack of wettability of the graphene sheets by copper. The relative density changes showed the decrease of relative density by increasing the reinforcement content, according to Fig. 10a. The minimum and maximum relative density was related to the pure copper and Cu-3.2 vol% Ag/rGO composites, which were 99.4 and 95.8%, respectively (Fig. 10a). During the milling process, the bonding 8
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Fig. 12. SEM images and EDS elemental mapping of a) Cu-0.8 vol% Ag/rGO, b) Cu-1.6 vol% Ag/rGO, and c) Cu-3.2 vol% Ag/rGO samples.
the grain boundaries, which can act as an obstacle against the deformation of the composite during loading. One of the important mechanisms of strengthening in composites is the load transfer from the matrix to the reinforcement phase. In copper-graphene composites, the proper load transfer does not occur due to the poor bonding between the copper matrix and graphene. However, previous research [43] shows that the use of graphene oxide instead of graphene can reduce the deboning between copper and graphene by the formation of the Cu–O bond (bonding Cu with oxygen-containing functional groups on the surface of graphene oxide). In the present work, the use of Ag/rGO reinforcement could help to establish the bond between the reinforcement and the matrix by the formation of Cu–Ag bonds, subsequently improving the composite strength by increasing the load transfer from copper to Ag/rGO. Finally, the Orowan's strengthening is another mechanism effective on the strengthening of the metal matrix composites [44–46].
composites due to an increase in the agglomeration probability of graphene, which was accompanied by the deboning of graphene and copper as a result of the low wettability of copper. Fig. 10c exhibits the maximum bending strength of 472 MPa at the graphene volume percentage of 0.8, which was 1.65 times more than that of pure copper. In addition, the elongation of fracture was decreased gradually with increasing the graphene content beyond 0.8 %vol. There are several different mechanisms for strengthening the metal matrix composites by adding graphene. The addition of graphene influences the microstructure of metal matrix composites; the high density of dislocations is created in the vicinity of graphene-metal intersection due to the high thermal expansion mismatch between graphene and the matrix. The increase of dislocations could improve the strength of the metal matrix composites. The reduction of the grain size is another effect of adding graphene on the microstructure of the metal matrix composite. The reduction of grain size is equal to an increase in
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● The change of the fracture mechanism from a plastic one to a brittle one occurred in the Cu-0.8 vol% Ag/rGO sample. The change of the fracture mechanism originated from the increase of the graphene content in the samples and this as accompanied with the aggregation of the graphene sheets.
Fig. 11 shows the fracture surface of the sintered samples containing different volume fractions of graphene decorated with Ag nano-particles. The fracture surface morphology of pure copper and Cu-0.4 vol% Ag/rGO samples revealed the grains with a rounded corner as a typical characteristic of the plastic fracture. The fracture surface of the specimen containing 0.8 vol% Ag/rGO demonstrated a sharper corner of grains. Furthermore, the pulled-out and separated grains from boundaries in the Cu-0.8 vol% Ag/rGO sample (Fig. 11 c) showed the change of fracture mechanism from a plastic one to a brittle one. The fracture surface morphology of Cu- 1.6 vol% Ag/rGO and Cu3.2 vol% Ag/rGO samples exhibited a layer-by-layer structure indicating a brittle fracture. The porosity and layer-by-layer structure marked in Fig. 11 d and e could be attributed to the debonding of the copper-graphene interface as a result of the weak wettability of copper on the agglomerated graphene. The bending strength of Cu- 1.6 vol% Ag/rGO and Cu- 3.2 vol% Ag/rGO samples was low because of the agglomeration of graphene and formation of porosities between the graphene layers, which could be considered as the stress concentration centers and the preferred sites of crack nucleation. Therefore, accumulation of the initial shallow cracks formed larger cracks, ultimately leading to the brittle fracture of Cu- 1.6 vol% Ag/rGO and Cu- 3.2 vol% Ag/rGO samples, as compared to the Cu- 0.8 vol% Ag/rGO specimen. From the above results, it may be concluded that the graphene nanosheets in 0.8 vol% of the Ag/rGO sample did not act as stress concentration centers and even served as a barrier for dislocation motion, thereby helping to improve the bending strength of the composites. Fig. 12 exhibits SEM images along with their elemental map analysis of the sintered composites containing 0.8, 1.6 and 3.2 vol% Ag/rGO, confirming the presence of graphene sheets. Additionally, Fig. 12a shows the rupture area of the graphene sheet in the fracture surface of the copper matrix composite. The tear and string structure of graphene sheets in the fracture cross-section could be attributed to overcoming the bending forces on the bond between Ag nanoparticles on the graphene plate and the copper matrix. This is because, if there were no bond between graphene plate and copper matrix, then the graphene sheets could not be torn off. Moreover, the transparency of the graphene sheet can be observed in Fig. 12b, which is one of the characteristics of graphene plates due to the very small thickness of these plates. The agglomeration of graphene plates in the Cu- 3.2 Vol.%Ag/ rGO sample is clearly observable in Fig. 12c, showing a non-uniform distribution of graphene and confirming the decrease of the bending strength values for this composite.
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4. Conclusions ● Ag/rGO nanocomposite was synthesized by a chemical reduction of graphene oxide in an aqueous solution under hydrazine hydrate as the reducing agent and silver nitrate. The morphology of the synthesized nanocomposite showed the homogenous distribution of Ag nanoparticles with a size in the range of 40–55 nm on the rGO sheets. ● An Ag/rGO nanocomposite with different volume contents was used as the second phase in the copper matrix composite. Cu– Ag/rGO composites were prepared by mechanical milling and spark plasma sintering. ● The samples of the sintered composites had high relative density. The relative density of the composites samples was reduced from 99.4% to 95.8% for the samples of pure copper and Cu- 3.2 vol% Ag/rGO, respectively. ● The micro-hardness and bending strength of the composite samples showed a significant increase, as compared to the pure copper sample. The highest values micro-hardness (81 Hv) and bending strength (472 MPa) belonged to Cu-1.6 vol% Ag/rGO and Cu-0.8 vol % Ag/rGO samples, respectively, while the values of micro-hardness and bending strength of pure copper were obtained to be 51 Hv and 285 MPa, respectively. 10
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