Strain Rate Dependent Shear Plasticity in Graphite Oxide - Nano

Letter pubs.acs.org/NanoLett

Strain Rate Dependent Shear Plasticity in Graphite Oxide Soumya Vinod,† Chandra Sekhar Tiwary,*,† Leonardo D. Machado,‡ Sehmus Ozden,† Juny Cho,† Preston Shaw,† Robert Vajtai,† Douglas S. Galvaõ ,*,‡ and Pulickel M. Ajayan*,† †

Department of Materials Science and Nanoengineering, Rice University, Houston, Texas 77005, United States Instituto de Física “Gleb Wataghin”, Universidade Estadual de Campinas, CP 6165, 13083-970 Campinas, São Paulo, Brazil

‡

S Supporting Information *

ABSTRACT: Graphene oxide film is made of stacked graphene layers with chemical functionalities, and we report that plasticity in the film can be engineered by strain rate tuning. The deformation behavior and plasticity of such functionalized layered systems is dominated by shear slip between individual layers and interaction between functional groups. Stress−strain behavior and theoretical models suggest that the deformation is strongly strain rate dependent and undergoes brittle to ductile transition with decreasing strain rate.

KEYWORDS: Graphene oxide, strain rate, stick−slip, brittle-ductile transition, MD simulation The films were synthesized by vacuum filtration of GO colloidal dispersion in water. The shape and thickness of the film is easily tunable using this method. Figure 1 shows the morphology of the film at different magnifications (optical to SEM). A well-defined continuous stacked structure can be clearly seen in the figure. A detailed structural (XRD, TEM, and AFM) and spectroscopic (XPS, Raman, and FTIR) analysis are presented in Supporting Information (Figure S1−S3). The results confirm the obtained structure is similar to previous GO films reported. The tensile test of GO film was conducted using a dynamic mechanical analyzer (DMA). The specimens used were rectangular strips (L = 50 mm, W = 4 mm). The engineering stress vs strain plots of controlled strain rate experiments ((0.1%/min(ε1), 0.05%/min(ε2), and 0.01%/min(ε3)) are shown in Figure 2A. The tensile strength shows significant strain rate dependency, decreasing from 85 to 50 MPa. Meanwhile, the fracture strain values show drastic increase, from 0.1 to 9%. A low magnification representative image of brittle and ductile fracture of the specimen is shown in Figure 2B and C, respectively. The major difference in morphology can be observed at subsurface (near to fracture surface). Low to high magnification images shown in Figure 2D, E, and F reveal the steps (deformation/slip bands) near fracture surface. An AFM topography image (Figure 2G) shows that the surface is rather rough, with step-like features. As we lower the strain rate, the

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arge efforts have been made in the materials community to introduce plasticity into high strength structural materials. Reducing grain size, reinforcing with soft phases, and activating slip systems are few of the classical approaches of tailoring plasticity in brittle materials.1 Since the discovery of graphene, numerous theoretical studies have predicted it to be the strongest material.2−5 However, recent mechanical testing of polycrystalline graphene shows brittle behavior with a very limited plasticity that hinders their overall mechanical performance.6,7 Graphene oxide (GO) is constituted of graphene layers with copious oxygen functionalities on the basal planes and edges that leads to increased interlayer spacing which gives rise to different mechanical properties.8−14 Recent studies of 3D stacking of these 2D GO sheets (frequently termed as GO film) show that high stiffness is governed by various factors such as interlayer spacing, bonding between layers, and layer uniformity.15−17,16However, the fundamental understanding of the mechanical response of these complex structures, especially plasticity, is still missing. The origin of plasticity in such atomically thin layered materials can be of importance for future applications of several recently developed 2D materials that go beyond graphene. In this work, we demonstrate that plasticity in GO film (by controlling strain rate) originates from the sliding of layers in association with significant frictional stick−slip motions (lock-key type) produced by the interactions among the functional groups. Tensile testing of GO film at different strain rates was performed to quantify plasticity. Fracture patterns were investigated through SEM and AFM techniques. MD simulations were also carried out to gain insight into the plasticity mechanisms. © XXXX American Chemical Society

Received: October 26, 2015 Revised: December 23, 2015

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DOI: 10.1021/acs.nanolett.5b04346 Nano Lett. XXXX, XXX, XXX−XXX

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Figure 1. Morphology of GO paper. (A) Digital image of the GO paper (3 cm diameter), and SEM images at different magnifications showing uniform morphology of the film without any discontinuities, (B) side view schematic of GO paper showing the layered structure and corresponding cross section SEM image.

Figure 2. Mechanical response of GO paper. (A) Engineering stress vs strain at different strain rates 0.1% (ε1), 0.05% (ε2), and 0.01% (ε3) which shows increase in plasticity as strain rate decreases. Inset showing the setup and loading direction of the GO paper. (B−C) Lower magnification SEM images at two extreme strain rates (ε1) and (ε3), respectively. (D−F) Higher magnification images having step like edges indicating sliding of the GO platelets. (G) AFM scan of the fractured surface confirming the bands from the sliding of GO layers. (H) Schematic showing the steps. (I) The depth of the band from fracture surface for the three different strain rate tests.

stress required for yielding decreases while the fracture strain increases multiple-fold. Such large strain is accommodated by sliding the layers and interlocking. A detailed investigation using different magnification SEM imaging of top fracture surface as well as subsurface clearly shows the slip bands are results of sliding rather than delamination of layers or crack

propagation through the interface between the layers. This suggests that the deformation bands are caused by sliding of GO layers, as this would cause the appearance of step-like features, as schematically illustrated in Figure 2H. Finally, note that the length of total deformed region increases as we decrease strain rate, as shown in Figure 2I. B

DOI: 10.1021/acs.nanolett.5b04346 Nano Lett. XXXX, XXX, XXX−XXX

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By measuring the energy change (ΔE) for a given displacement, we were able to estimate the interlayer force, F ≈ −ΔE/Δx. Figure 4A shows the variation of the interlayer force as a

Figure 3. Strain rate dependent plasticity. (A) Raman spectra before and after the straining experiments. No visible changes are observed, suggesting no deformation of individual GO flakes. (B) number of bands (blue) and aspect ratio (h/L) of deformation band (red) as a function of the strain rate. (C) Schematic model summarizing experimental observations. (D) Representative snapshots from MD simulations for three strain rate values ε•3t = 4ε•2t = 10 and ε•1t = 20 Å/ps. The left side of the bottom flake was fixed, and the right side of the one on top was pulled.

Figure 4. Plasticity mechanism: Force and conformation analysis. (A) Variation of interlayer interaction force between top and adjacent layers as a function of displacement, for low (yellow) and high (blue) strain rates. Small displacements of the top sheet can separate interacting functionalities or bring them closer together. Positive and negative forces will arise accordingly, giving rise to oscillations. (B) Snapshots from MD simulations for high and low strain rate conditions. The color in the figures correspond to the local charge value, which can be read from the scale bar. (C) Zoom-in of a region in which the positive hydrogen of an OH functionality interacts strongly with the negative charge of two epoxide functionalities.

Raman spectra taken before and after the mechanical test did not show changes in the ID/IG ratio (Figure 3A), indicating that the tensile force did not lead to permanent stretching or plastic deformation of the individual GO sheets. The data presented in blue in Figure 3B shows that the number of the bands increases as the strain rate decreases. The aspect ratio height/length (h/L, see Figure 3C) of the deformation band closest to the fracture surface in turn increases with increasing strain rate. Figure 3C summarizes our experimental findings. At high strain rate, we observe lower number of bands with higher aspect ratio. As we decrease the strain rate, the length of band increases and the aspect ratio decreases. In order to elucidate the origin of deformation bands as a function of strain rate, we carried out fully atomistic molecular dynamics (MD) simulations. Our simplified model consisted of stacked GO layers, subjected to different strain rates mimicking the experimental loading conditions. More detailed information is provided in the Supporting Information. Representative MD results for different strain rates are presented in Figure 3D. For clarity we are displaying the case for three layers. Our results show that for the highest strain rate the load transfer from top layer to subsequent layers is not significant, which results in sliding of the top layer. This stems from the fact that the upper layer is moving too fast. For the intermediate rate we started observing layer coupling, which results in layers moving together, creating deformation bands. This is amplified for the lowest strain rate (see Movie S1, which compares different strain rates). A better understanding of the deformation band dynamics can be obtained from the analysis of interlayer force profiles between the top and middle layers. To obtain the interlayer force, we first calculated the interaction energy between the top and the adjacent sheet. We then applied a displacement (Δx) to the upper sheet and recalculated the interaction energy.

function of the top layer displacement, for cases representative of low and high strain rates. The corresponding energy profiles can be found in the Supporting Information. When strain rate is low/high (yellow/blue), interaction forces are higher/lower, and layer coupling is more/less effective. This behavior can be explained by analyzing GO film conformations during the MD simulations. In Figure 4B we show representative snapshots for the low (ε•t3) and high (ε•t1) strain rates. For ε•t1, the layers have weak interactions and keep their planar form, and just the top layer slides. On the other hand, for ε•t3, we observe a strong interaction (mainly Coulombic and van der Waals contributions) giving rise to local deformation where the functional groups are present. Figure 4C shows the zoomed OH group of the upper sheet interacting strongly with two epoxide groups present in the subsequent sheet. Such strong interaction causes stick−slip like movement (see also Movies S2 and S3). The stick−slip phenomenon is rather common. The macroscopic theory indicates there is a tendency for intermittent motion when the kinetic friction is less than the static friction.18,19 Other molecular dynamics investigations have been carried out to investigate its atomistic origin. Nucleation and motion of dislocations;20 elastic deformation of surface layers;18 and, for two passivated diamond surfaces, hydrogen−hydrogen repulsive interactions21 have been found to lead to stick−slip movement. The stick−slip mechanism in our simulations is similar to the latter. Rather than repulsive, however, analysis of our trajectories indicates that stick−slip stems from attractive C

DOI: 10.1021/acs.nanolett.5b04346 Nano Lett. XXXX, XXX, XXX−XXX

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not described by the standard force field were parametrized using the ParamChem Web site, which is led by one the force field cocreators.27 In all our simulations, a Nosé−Hoover thermostat was employed to keep the temperature at 300 K, and the integration of the resulting equations of motion was carried out using the LAMMPS MD package,23 with a time step of 0.25 fs.

interaction between the functional groups. Note also that the transition from smooth sliding to stick−slip motion by decreasing the velocity of the moving object has been previously reported.22 Although the real material is much more complex (bent regions, defects, disorder, etc.), our simplified model captures the essential experimental aspects. Theory and experiments support that GO film plasticity can be explained in terms of interaction of functional groups (structural locking, van der Waals and coulumbic contributions) between two layers. The classical plasticity in crystalline materials is explained with help of dislocation and twinning based models, but in case of nanoscale materials such as carbon nanotubes the mechanism is quite different.23,24 The plasticity generated in nanolayered structure of GO architecture is similar to that observed in natural layered materials, such as nacre shell.25 Our studies indicated that it is the presence of functional groups between the layers that is the predominant factor in creating the stick−slip movement that leads to strain controlled plasticity. Similar MD studies on graphene showed the layers tend to slide away easily as there is very limited interaction between the layers and the force and energy profiles showed much variation from that of GO layers (shown in Figure S5 of Supporting Information), and a comparison video of graphene and graphene oxide sliding is shown in Video S3. In conclusion, the GO film deformation behavior presents a unique plasticity mechanism depending on the strain rate. Contrary to the easy cleavage leading to brittle fracture in graphite, the functional groups present on the basal planes of GO avoid easy cleavage and prevent mobility between layers. This work is of broader importance, as we can exploit the functional chemistry to engineer the plasticity of other pure or hybrid 2D materials. Materials and Methods. Synthesis. GO was synthesized using Improved Hummer’s method. In brief, 3 g of SP1 graphite powder (Bay Carbon) was mixed with 18 g of potassium permanganate and stirred well in 360 mL of sulfuric acid + 40 mL of phosphoric acid. After 12 h stirring, the solution was poured onto ice from 500 mL of DI water, and 14 mL of hydrogen peroxide was added to it. The solution turns yellow in color and is filtered and washed repeatedly using DI water, 30% HCl, and ethanol to remove unwanted ions. The dried material obtained is GO, and for making the GO paper, 100 mg of GO powder is sonicated in 10 mL of water, and 20 μL of glutaraldehyde, and a very small amount of resorcinol is added to it. The GO solution is sonicated for 3 h and vacuum filtered through a 4.5 mm diameter, 0.22 μm pore sized PTFE membrane. After 24 h the solution dries and forms a dry, flexible film which can be peeled off from the membrane. Characterization. FTIR for determining the different chemical bonds present in GO paper was analyzed using Nicolet FTIR infrared microscope. Renishaw in Via Raman Microscope at laser excitation of 514.5 nm was used for studying the Raman spectrum of the sample. The XPS data was taken using PHI Quantera XPS. The sample morphology and cross-section surface was imaged using FEI Quanta 400 scanning electron microscope under high vacuum at 20 kV. TEM images were taken using JEOL 2100 field emission gun transmission electron microscope.The load/unload tensile test on rectangular strip of GO was done on Q800 dynamic mechanical analysis. MD Simulation Details. We used the well-known CHARMM force field to describe atomic interactions.26 Those that were

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.nanolett.5b04346. Detailed structural (XRD, TEM, and AFM) and spectroscopic (XPS, Raman, and FTIR) analysis (PDF) Comparison video of different strain rates (AVI) Stick−slip like movement (AVI) Comparison video of graphene and graphene oxide sliding (AVI)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. *E-mail: galvao@ifi.unicamp.br. Author Contributions

S.V., C.S.T., and L.D.M. contributed equally to the work. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Authors acknowledge the funding support from U.S. Department of Defense: U.S. Air Force Office of Scientific Research for the Project MURI: “Synthesis and Characterization of 3-D Carbon Nanotube Solid Networks” Award No. FA9550-12-1-0035 LDM and DSG acknowledge financial support from the Brazilian Agencies CNPq and FAPESP.

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REFERENCES

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DOI: 10.1021/acs.nanolett.5b04346 Nano Lett. XXXX, XXX, XXX−XXX