Strong and Tough Layered Nanocomposites with Buried Interfaces

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Strong and Tough Layered Nanocomposites with Buried Interfaces Ke Chen, Xuke Tang, Yonghai Yue, Hewei Zhao, and Lin Guo ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.6b01752 • Publication Date (Web): 12 Apr 2016 Downloaded from http://pubs.acs.org on April 12, 2016

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Strong and Tough Layered Nanocomposites with Buried Interfaces Ke Chen, Xuke Tang, Yonghai Yue*, Hewei Zhao, Lin Guo* Key Laboratory of Bio-Inspired Smart Interfacial Science and Technology of Ministry of Education, Beijing Key Laboratory of Bio-inspired Energy Materials and Devices, School of Chemistry and Environment, Beihang University, Beijing 100191, P R China Email: [email protected] & [email protected] KEYWORDS: bio-inspired, nacre, mechanical properties, ions, graphene oxide,

carboxymethyecellulose

ABSTRACT: In nacre, the excellent mechanical properties of materials are highly

dependent on their intricate hierarchical structures. However, strengthening and toughening effects induced by the buried inorganic-organic interfaces, actually originate from various minerals/ions with small amounts, and have not drawn enough attention yet. Herein, we present a typical class of artificial nacres, fabricated by graphene oxide (GO) nanosheets, carboxymethylcellulose (CMC) polymer and multivalent cationic (Mn+) ions, in which the Mn+ ions cross-linking with plenty of oxygen-containing groups serve as the reinforcing ‘evocator’, working together with other cooperative interactions (e.g., hydrogen (H)-bonding) to strengthen the GO/CMC interfaces. Comparing with the pristine GO/CMC paper, the cross-linking strategies dramatically reinforce the mechanical properties of our artificial nacres. This special reinforcing effect opens a promising route to strengthen and toughen materials to be applied in aerospace, tissue 1

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engineering, and wearable electronic devices, which also has implication for better understanding of the role of these minerals/ions in natural materials for the mechanical improvement.

The attainment of both strength and toughness is a vital requirement for engineering materials. Unfortunately, these properties are mutually exclusive in most cases. Researchers have to compromise high strength to avoid catastrophic failure in most safety-critical cases, which greatly limits materials’ applications.1,2 However, nature uses its genius and magical ways to produce lightweight, strong, and tough materials with complex, hierarchical architectures.3,4 Nacre, a great gift from nature, has been intensively studied recently due to its outstanding mechanical properties profiting from its staggered “brick-and-mortar” architectures.5-8 Based on the nacre’s structure, most of the bio-inspired layered materials were fabricated primarily by mimicking the ordered microstructures and simple interfaces to achieve combination of stiffness, strength and toughness.7-9 As previously reported, the relationship between the mechanical property and the hierarchical structures of natural nacre has been well demonstrated.10-16 The hard minerals (mostly indicated as CaCO3) account for the high strength;8 the intrinsic toughening mainly originates from crack deflection at the platelet interfaces, since crack deflection, such as propagating along the interfaces, circumventing the microscopic platelets and generating tortuous path, dissipates more energy.13,17-19 However, mechanical enhancement coming from other structural and component factors has been always ignored.3,4,8,15,17 For example, during the bio-mineralization process, besides the 2


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predominant mineral (CaCO3) and organic (polysaccharides and proteins) component, biological organisms from mollusks can guide some small amount of other minerals/ions (e.g., Mg2+, Al2O3, ZrO2, TiO2, CaO and so on) to form the layered composite materials,14,15 these minerals/ions are probably occluded during mineralization and interacted with organic matrix to form buried inorganic-organic complex interfaces at multiple hierarchical levels, referring to the mineralized platelets and organic matrix. Although it has been reported that some special minerals/ions play key roles in controlling tissue fiber formation, matrix-mineral interactions, and crystal polymorph and orientation,7,15 it is still very difficult to prove and even quantify the improvement of mechanical performance caused by various minerals/ions with small amounts due to the lack of ideal models. The obscure synergistic buried interfacial effects caused by these small amounts of minerals/ions in nacre or some other natural materials on the mechanical property have not actually been studied systematically and comprehensively. Therefore, it remains questions whether these minerals/ions have the ability to accommodate large strains by elastic uncoiling as the organic matrix molecules in nacre, and whether the introduction of these buried/complex interfaces in nacre’s platelets can really enhance the intrinsic toughness by crack deflection? The answers to these questions will play a vital role in the improvement of mechanical properties of nacre-inspired materials.

In this work, inspired by buried inorganic-organic interfaces derived from various minerals in nacre, a simple vacuum-assisted filtration (VAF) was used to fabricate a 3


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typical class of artificial nacres through assembling hybrid GO/CMC building blocks via cooperative interactions (e.g., H-bonding) at first and then followed by intercalation of three different types of Mn+ ions (M = A2+ (Mg2+, ZrO2+, Ni2+, Ca2+, Cu2+, Co2+), (TiO)n2n+ (TiO2+), and Al3+), which was different from the previous reports by only studying one/two ionic cross-linking strategy.3,20,21 Herein, GO nanosheets serve as “bricks” while CMC polymers are taken as “mortar” phase, and Mn+ ions cross-linking with plenty of oxygen-containing groups are invoked as the reinforcing ‘evocator’, in combination with other cooperative interactions, to enhance the buried inorganic-organic interfaces, providing us an ideal model to explore/understand its synergistic effect derived from various minerals on the mechanical properties in nacre or other natural materials. Experimental results show that, the mechanical properties of our artificial nacres have been improved significantly. The tensile strength and toughness have been measured to be approximately 198.2 ~ 304.5 MPa (A2+:198.2 ~ 225.2 MPa; TiO2+: 228.2 ± 11.0 MPa; Al3+: 286.4 ± 19.4 MPa) and 5.2 ~ 18.9 MJ·m-3 (A2+: 5.2 ~ 9.7 MJ·m-3; TiO2+: 15.7 ± 3.2 MJ·m-3 Al3+: 5.4 ± 1.1 MJ·m-3), respectively, which are 2 ~ 3 and 3 ~ 10 folds higher than those of some natural materials (e.g., nacre, bone, dentin) and superior to other mostly binary/ternary layered GO- and CMC-related nanocomposites. Besides, three types of cross-linking strategies are clearly proved by the variation in micro-mechanical behaviors. Considering the excellent mechanical properties and high flexibility combined with biocompatibility, this strategy will open a promising route to design outstanding engineering materials to be applied in aerospace, tissue engineering, and wearable electronic devices potentially. 4


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RESULTS AND DISCUSSION Nacre has characteristic structural features on multiple length-scales from molecular to near-macroscopic dimensions (Figure 1a).1-3 We consider that apart from aragonite crystals, β-chitin fibrils, and analogous silk protein, small amounts of inorganic minerals/granules are randomly dispersed in nacre, which can cause complex buried hybrid structural interface to enhance the capability of energy dissipation and crack deflection

under stretch/compress.15

Structure

characterization

by

transmission

electron microscopy (TEM) proved the existence of the complex grain and phase boundary, which caused the formation of complex interfacial polycrystalline system in nacre (aragonite) (Figure 1b-e). The amorphous phase and boundary could be attributed to some biopolymer or various minerals. Scanning electron microscope-energy dispersive spectra (SEM-EDS) element mapping, combined with inductively coupled plasma mass spectroscopy (ICP-MS) analysis verified the ingenious hierarchical structural feature, where a relative poor staggered platelets microstructure in bivalve pearl oyster shell (e.g., Clam), and a well-organized staggered “brick-and-mortar” microstructure in nacre from red abalone,8 and various mineral elements of nacre, as shown in Figure 1f and Figure S1a (see supplementary information (SI)). Obviously, a minority of other minerals (e.g., Al2O3, TiO2, ZrO2, and MgO) also directly participated in the growth of nacreous platelets to form complex hybrid mineral crystals, being in contact with the organic membrane/matrix between platelets to form buried inorganic-organic interfaces. The complex hybrid mineral crystal is analogous to the chemical composition of the chiton 5


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tooth in the nanoscale.15 In spite of various minerals segregation, a relative uniform distribution was demonstrated by the atomic force microscopy (AFM) phase map (Figure S1b-c) and via comparing secondary electron (SE) images with backscattered electron (BSE) images of the platelets in nacre (Figure S1d-g). Moreover, quantification of the micro-mechanical behavior of freshly cleaved Clam shells had been carried out via TI 950 nano-indentation (Figure 1g-h). The heterogeneous distribution ranging from approximately 32.3 GPa (Ec) and 0.63 GPa (Hc) to 138.2 Ga (Ec) and 6.07 GPa (Hc) were obtained by analyzing the contour map of Ec and Hc dispersion of nacreous platelets in the Clams within 10 µm × 10 µm, respectively. This result verifies the existence of the complex buried hybrid interface again, which is in agreement with experimental observations in the bivalve Placuna placenta.22 On one hand, these small amount minerals/granules existed in the organic membranes (interfaces) of nacre play a vital role in the mechanical response to tensile stress, as they can potentially change the mechanical deformation and energy dissipation of organic membranes.22 On the other hand, under compression, they can also help dissipate more energy through prompting microcracks in platelets to propagate along the probably buried interfaces, generating tortuous path, which dissipate more energy. Therefore, we can also acquire theoretical insights into the importance of the buried interface for mechanical properties of the staggered composite such as nacre (see these theoretical models in SI).23

Our hypothesis about the critical effect of mineral ions in nacre was further testified experimentally. A series of advanced artificial nacres with both the ingenious hierarchical 6


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structure and the different buried mineral components (i.e., different Mn+ ions) reinforced by inorganic-organic interfaces have been fabricated successfully. The designed fabrication process and proposed structural model for the artificial nacres are shown in Figure 2a. Hard “bricks” were hybridized by individual GO nanosheets, achieved by the attractive cooperating interactions. Individual GO nanosheet has been well known as a unique two-dimensional sheet structure containing various oxygen functional groups, with light weight, high aspect ratio, and outstanding mechanical properties. These extraordinary mechanical properties make GO as an ideal candidate of hard platelet-filler to strengthen the composites.9,24,25 Besides, the organic membranes of nacre are taken place by CMC, which is composed of β-(1-4)-linked glycanes, a water-soluble anionic linear polysaccharide produced from cellulose, monochloroacetic acid and sodium hydroxide.26 The adhesion properties and the numerous hydroxyl and carboxylic groups in CMC can be exploited to enhance the interfacial interactions between GO nanosheets and CMC, e.g., H-bonding network, van der Waals interaction. In terms of molecular design, molecular structure of CMC in nacre is similar to that of chitin as organic matrix. Thus, CMC can be utilized as a good soft/organic matrix to enhance the toughness in the artificial nacre composites. Then, nacre-like layered composites can be fabricated by VAF. Interactions on the GO/CMC interfaces are generally regarded as multiple H-bonding network and van der Waals interactions.27,28 Subsequently, different Mn+ ions are incorporated into the binary GO/CMC system to bond with oxygen-contained groups on the GO/CMC interface, a typical type of artificial nacres, named by Mn+-GO/CMC formed which can be sorted into three types of papers based on the hybrid interfacial 7


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reinforcing mechanism: A2+-, TiO2+-, and Al3+-GO/CMC, where A2+ represents divalent metal ions (e.g., Mg2+, ZrO2+, and so on), TiO2+ represents zigzag long chain (TiO)n2n+, A3+ represents trivalent metal ions (e.g., Al3+). It is worth noting that the zigzag long chain (TiO)n2n+ is formed by TiO2+ because of no simple TiO2+ ions in aqueous solution of Titanyl sulfate (TiOSO4).29

To fabricate the composite papers, GO was firstly oxidized from natural graphite flakes by a modified Hummers method.30-34 The optimal volume ratio of the colloidal GO suspension to aqueous CMC solution was determined by the optimization of mechanical properties (Figure S2-S3). The change of GO content could be easily monitored by UV-vis spectroscopy as shown in Figure S2c. The addition of Mn+ ions immediately resulted in agglomeration of the blends and uncontrollable precipitation to some extent, probably owing to the chemical cross-linking between Mn+ ions and reactive oxygen-contained functional groups located at the GO/CMC interface (Figure S2a). Different red shifts occurred for the absorption edge of the blends after the addition of different Mn+ ions, which proved existence of the cross-linking reactions (Figure S2c). Notably, the optimal amounts of Mn+ ions to improve the mechanical properties have also been determined based on a series of mechanical tests (see Figure S3 and Table S1) since the optimal amount of Mn+ ions could obtain the strongest chemically cross-linking density on the GO/CMC interfaces. Our artificial nacres (e.g., Al3+-GO/CMC) were brown flexible papers (Figure 2b) with uniform phase distributions as shown in Figure 2c. Besides, Figure 2d-e showed the topography and height profile of the composites with a 8


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relatively small root-mean square (RMS) roughness (134.5 ± 5.8 nm), indicating the flattening of nanoscopic asperities on the plane of artificial nacre. The direct observation of the closely-packed composite layers could be obtained by cross-sectional SEM analysis of the fracture surface (Figure 2f-g), indicating that the GO/CMC interface has undergone strong cross-linking reaction. Other samples also had analogous phenomenon. Besides, the thickness difference (1.0 ~ 2.5 µm) between GO/CMC and Mn+-GO/CMC also confirmed the existence of a certain amount of Mn+ ions (Figure S4a). SEM-EDS element mapping images showed that these eight different types of cations were dispersed uniformly in the samples (Figure S4b), which are in good consistence with our predicted results.

The X-ray diffraction (XRD) patterns of the Mn+-GO/CMC papers (Figure S5a-b) presented a typical diffraction peak, suggesting the ordered structures31 and showing a slight increased d-spacing (d = 0.84 ~ 0.88 nm) compared to that of the GO/CMC paper (d = 0.84 nm) (Figure S5a), which indicated the existence of different Mn+ ions in the gallery regions (their ionic radius ranges from 53.5 pm for Al3+ to 106 pm for Ca2+, see Table S2). Meanwhile, the peak intensities were slightly lower and the diffraction peak position shifted to lower θ values in various degrees than that of the GO/CMC paper, which was caused by the combination of the different bonding energies in the M-O bond and different intercalation contents of the ions.9,35 The Fourier transformed infrared (FTIR) spectra (Figure S5c) exhibited inordinately decreased C=O stretch intensities, increased carboxyl C-O, and increased aromatic C=C/carboxyl O-C=O stretch intensities 9


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whose peak positions shifted to lower wave numbers to some extent in Mn+-GO/CMC. Simultaneously, the peak position of the carboxyl O-C=O bending vibration also shifted to higher wave numbers. This can be interpreted as an evidence for the coordination of carboxylic acid to Mn+ ions, which can potentially cause chemical cross-linking reaction.35,36 Noticeably, the relative intensity of the epoxy/ether C-O stretch at 1222 cm-1 gradually decreased from bottom (GO) to top (Al3+-GO/CMC) (Figure S5c) and disappeared for the Mn+-GO/CMC, accompanying with a simultaneous slight increase of the relative intensity of the stretch at 1100 cm-1, corresponding to hydroxyl and alkoxide C-O stretches, which could be attributed to the ring-opening of the epoxides. In the C1s XPS spectra (Figure S5d), the C1s components of O-C=O groups shifted slightly from 289.2 in GO/CMC to 289.1/289.0 eV in Mn+-GO/CMC, indicating the coordination between Mn+ ions and O-C=O groups.27,37 Meanwhile, the C-OH percentage decreased dramatically from 9.3 % for GO/CMC to 2.3 % for Mg2+-reinforced, further down to 1.0 % for Al3+-reinforced composite, together with the great decrease of the C-O-C percentage at 286.0 eV for most of Mn+-GO/CMC (Table S3), suggesting that the chemical cross-linking reactions resulted in the reduction of these fitting C1s peaks.4,31 Apparently, the difference in the C1s XPS spectra, including the peak position shifts and the percentages decreased, was mainly due to the different chemical cross-linking capabilities (Table S2). Raman spectra (Figure S5e) showed that all of the ID/IG ratios (1.36 ~ 1.10) in Mn+-GO/CMC were clearly lower than that (1.46) of the GO/CMC paper, indicating that Mn+ ions could enhance the interaction behavior between GO and CMC significantly.38 10


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Additionally, in-situ Raman measurements (Figure S5g) revealed that the downshifted trend of the G- band for Mn+-GO/CMC was bigger than that of the GO/CMC paper, indicating that the load-bearing capability of the GO nanosheets was clearly enhanced after adding Mn+ ions.37,39 Notably, the different load-bearing status of artificial nacres could bring different final downshifts. Mg2+-GO/CMC showed the smallest downshifts, whereas Al3+-GO/CMC exhibited the biggest final downshifts. As we know, the downshift of the G- band represents the strain of the neighboring GO nanosheets in the GO-containing composite papers. Therefore, larger downshifts suggest increased shear strength of adjacent GO nanosheets.

The macro-mechanical performance were evaluated by tensile testing under natural environment (25 °C, 15 % RH), as shown in Figure 3a-d and Table S1. In comparison with pure GO, CMC, or binary GO/CMC papers, the mechanical properties of our artificial nacres have been improved significantly (their notable improvements have been summarized in Table 1 in comparison with that of the GO/CMC paper), in which, an unique combinations of high strength (approximately 198.2 ~ 286.4 MPa, in some cases up to 304.5 MPa for Al3+-reinforced) and high toughness (about 5.4 ~ 15.7 MJ·m-3, in some case up to 18.9 MJ·m-3 for TiO2+) could be achieved. For example, the ultimate stress and toughness of A2+-GO/CMC reached 198.2 ± 2.9 ~ 225.2 ± 25.4 MPa and 9.0 ± 1.3 ~ 5.2 ± 2.0 MJ·m-3 from Mg2+ to Co2+, respectively, which were approximately 47.6 ~ 76.2 % and 173.7 ~ 410.5 % higher than the GO/CMC paper, with a tensile strength of 129.5 ± 7.9 MPa and 1.9 ± 0.6 MJ·m-3. For TiO2+-reinforced composites, the toughness 11


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achieved a very high value of 15.7 ± 3.2 MJ·m-3 (the largest improvement of ~ 726.3 %) with a relatively high ultimate strength (228.2 ± 11.0 MPa), which can be attributed to the zigzag long chain (TiO)n2n+ structure. In contrast, the most significant enhancement in the ultimate strength of 286.4 ± 19.4 MPa (increase of ~ 121.2 %) and the toughness of 4.4 ± 1.6 MJ·m-3 (127.7 ± 12.3 %) could be achieved. However, in terms of Young’s modulus, compared with the GO/CMC paper (6.9 ± 0.8 GPa), our artificial nacres apparently showed lower values/negative increases, except for Al3+-reinforced. In particular, we also found a very interesting phenomenon about the conflicts between Young’s modulus and toughness in the these papers (e.g., Cu2+-reinforced, Table S1). Here, the higher toughness obtained by adjusting the content of Mn+ ions is usually at the expense of the loss of Young’s modulus. Moreover, the creep-recovery curves of the Mn+-GO/CMC papers revealed smaller maximum deformation under the stress of 30.0 MPa and residual deformation after recovery compared with those of GO, CMC, and GO/CMC paper (Figure 3d), which further indicates that the strong chemical cross-linking capabilities from Mn+ ions can effectively restrain the slip/deformation between GO nanosheets and CMC layers, as demonstrated by the in situ Raman measurements (Figure S5g). With respect to A2+-reinforced, the maximum strains and residue strains ranged from 2.7 % and 2.0 % for Mg2+ to 1.9 % and 1.4 % for Ni2+ to 0.9 % and 0.5 % for Cu2+, respectively. The difference mainly due to the different bonding energies among A2+ ions, which had been rationalized in characterization parts (Figure S5). The relatively small maximum strain (1.1 %)/residue strain (0.5 %) could be obtain in Al3+-GO/CMC. Owing to the

12


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structure feature of the zigzag long chain (TiO)n2n+, TiO2+-GO/CMC can keep a balance between mild maximum strain and moderate residue strain. The micro-mechanical behaviors of the Mn+-GO/CMC papers in natural environment (25 °C, 15 % of RH) have been conducted by the nano-indentation, to further elucidate the reinforcing effects of different types of Mn+ ions (Figures 3e-g and Figure S6a-d). Residual indent region has been investigated by AFM, providing the unique insight into the micro-mechanical response of our artificial nacres (see Figure S6a, b).34 As a whole, with the addition of Mn+, load curves of Mn+-GO/CMC paper were even steeper (Figure 3e) compared with binary GO/CMC paper. And the slope also increased clearly from A2+ to TiO2+ to Al3+-reinforced to varying degrees. The steeper load curves also confirm the strengthening effect indirectly. Upon unloading, as the load was released, the behavior of the depth recovery was in accordance with that of the loading. Variations in the cationic metals leaded to significant changes in Young’s modulus (E'g-c) and Hardness (H'g-c), corresponding to typical load-displacement curves, as shown in Figure 3f. The GO/CMC paper exhibited a typical E'g-c, 4.8 ± 1.3 GPa and H'g-c, ~ 470.0 GPa which were obviously higher than those of the pure GO or CMC paper, indicating the interfacial H-bonding effect. Furthermore, although Mn+-GO/CMC papers had even higher value than that of the binary GO/CMC paper, with respect to their E'g-c (A2+: approximately 5.1 ~ 7.7 GPa; TiO2+: 5.5 ± 1.1 GPa, Al3+: 9.6 ± 2.1 GPa), the value of their H'g-c, except for ~ 553.8 MPa for Al3+-reinforced, showed visibly lower than that of the GO/CMC paper, indicating the notable improvement of toughness originated from the addition of the optimized ions content. Likewise, we could speculate that the significant enhancement in 13


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hardness could be attributed to the strongest chemical bonding energy (512.0 ± 4.0 KJ·mol-1, Table S2) of the Al-O bond. Figure 3g further displays the typical distribution maps of E'g-c (top) and H'g-c (bottom) of A2+- (e.g., Mg2+), TiO2+-, and Al3+-GO/CMC plane, corresponding to the results in Figure 3e, suggesting a relatively stable mechanical properties and microstructure of our artificial nacres. Figure 4a-d present the comparison of the micro-scale deformation (Figure 4a1-d1) and multiple individual load-displacement curves (Figure 4a2-d2) for GO/CMC, Mg2+-, TiO2+-, and Al3+-GO/CMC (see detailed process in Movies S1-4). For GO/CMC, Figure 4a1-2 reveal a largely deformed region/depth (red dashed circle; zone radius, R ≈ 20 µm), the maximum displacement (hmax ≈ 1500 nm) and an indentation residue (collapse) close to the indentation site with increasing of load. In contrast, Figure 4b-d show an intact region (red dashed circle: Mg2+, R > 21µm; TiO2+, R ≈ 22 µm; Al3+, R ≈ 15 µm; Figure 4b1-d1), a relatively small quasi-elastic deformation (Mg2+, hmax ≈ 500 nm; TiO2+, hmax ≈ 350 nm; Al3+, hmax ≈ 300 nm; Figure 4b2-d2). Indeed, the reinforcing order (Mg2+ < TiO2+ < Al3+) could be obtained. These results indicate that the different reinforcing effects and resistance deformations in the composites truly exit again, which is in accordance with the macro-mechanical analysis mentioned above.

In addition, the fracture morphologies of typical artificial nacres were carefully examined to further reveal the different synergistic reinforced mechanisms, as displayed in Figure 4e-g. The more dense inter-cross layered structures (Figure 4e1-3) as compared with the pristine GO/CMC paper (Figure 2f and Figure S6c), had been created in 14


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A2+-GO/CMC papers by intercalating different A2+ ions to strengthen the GO/CMC interfaces. Obviously, distinct differences in cross-section morphologies directly reflect different mechanical properties. Long-range cracks propagation paths have been monitored on the surface of the TiO2+- and Al3+-reinforced papers (Figure 4f1-g1). They exhibited a kind of nonlinear/tortuous fracture mode on their surfaces, which indicated the improved fracture energy.10,34 Meanwhile, inside the crack expanding (Inset in Figure 4f1-g1), no soft tearing was found during the fracture processes. Furthermore, the adhesive layers (the alternate arrangement of GO nanosheets and CMC layer) in TiO2+-GO/CMC performed the incomplete fracture feature, and showed a different fracture mode in Al3+-GO/CMC, large pieces of dense adhesive layers were pulled out as shown in Figure 4f2-g2, indicating a ductile feature for TiO2+-GO/CMC which is different from the rigid feature for Al3+-GO/CMC. Besides, the tidy fracture feature in most dense cross section morphology (Figure 4g3), clearly different from the jagged fracture in Figure 4g2 further proved the disparity. Base on the above analysis, three crack extension models were suggested to depict the synergistic reinforcing effect (Figure 4ei-gi). (i) Initially subjected to stress: a deflected microcrack caused by a GO nanosheet encounters the chemically cross-linking sites (Figure 4e1i-f1i). For A2+, a relatively large slip with slight deflection happened; for TiO2+ and M3+ (e.g., Al3+), the relatively large deflected microcrack with a relatively small slip can be caused by the strongest load transfer capability of the Al-O bonding and the increase of the friction force originated from the complex chains (-Ti-O-Ti-), together with chemical cross-linking sites on the GO/CMC interfaces, respectively. (ii) Increasing stress: the chemical cross-linking sites bridges the 15


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microcrack. The cross-linking Mn+ ions can effectively restrict the relative slip of these two parts (strain hardening). For TiO2+, the complex chains could improve the sliding distance of adjacent GO nanosheets obviously. The enhanced stress was transferred rapidly to a vicinal one along the interfacial layer, which subsequently motivated the potential loosening/break of adjacent multiple adhesive layers (Figure 4e2i-g2i). (iii) Fracture: such crack deflection, strain hardening and motivation of the potential loosening/break of multiple adjacent adhesive layers were accumulated step by step until the failure (Figure 4e3i-g3i). Besides, from the nonlinear stress-strain curve (Figure 3a), the deformation process, possibly including viscoplastic and plastic deformation, was demonstrated by the strain hardening, which is similar to the deformation of hydrated nacre.10,40 In this case, energy dissipation by sacrificing cooperative bonding and sliding of GO nanosheets and CMC layers, resulting in the synergistic reinforced strategy.

The mechanical properties of the composites containing CMC polymer with different fillers or different polymeric materials with GO fillers, including natural materials (e.g., natural nacre, bone, and dentin) and our artificial nacres, have been summarized in Figure 5a. Our artificial nacres reinforced by the synergistic interactions of different Mn+ ions cross-linking and intermolecular bonding achieve the advantage of integrated strength and toughness, comparing with some natural materials, other GO-, and CMC-related nanocomposites with different types of the reinforcing interactions (e.g., H-bonding, ionic bonding, covalent bonding, and some synergistic interactions), regardless of their stiffness. The strength is approximately 2 ~ 3 folds higher than that of natural nacre (80.0 16


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~ 135.0 MPa), bone (150.0 MPa), and dentin (105.0 ± 16.4 MPa), respectively.7 Meanwhile, the toughness of Al3+- and TiO2+-GO/CMC are more than 3 to 10 times higher than that of the natural nacre (1.8 MJ·m-3),10 5 to 15 times higher than that of the bone (1.2 MJ·m-3),41 and over 3 to 9 times higher than that of the dentin (2.0 MJ·m-3),43 respectively. Furthermore, the combination of the strength and toughness of our optimized composites is distinctly higher than most of GO- and CMC-based nanocomposites, such as GO/PPA,48 GO/(PAH-PSS),49 CMC/MTM,26 by hydrogen bonding; GO/Mn+ (Mg2+, Ca2+, Fe3+)35 by ionic bonding; GO/GA,37 GO/Borate,4 GO/PCDO,51 GO/PVA,53 by covalent bonding; GO/MoS2-TPU,54 Fe3+-GO/TA56 by synergistic interactions. However, although these GO/SF, GO/MTM, and GO/RGO-Silk composites exhibit very high strength of 330, 320, and 300 MPa, respectively, their toughness are only 3.4, 4.0 and 2.8 MJ·m-3 which is clearly lower than that of our artificial nacre;9,26,60 GO/PU and GO/PCL show the high toughness of 58.0 and 9.8 MJ·m-3, respectively, their ultimate stresses are only 13.0 and 20.3 MPa due to the weak polyurethane matrix).9,33 Although the strength and toughness A2+-GO/CMC are comparable to those of GO/CMC, GO/PVP, GO/G4NH2, and GO/PEI,56 the toughness of TiO2+-GO/CMC is much higher than that of them. While the strength and toughness of Al3+-GO/CMC which are similar to those of GO/CS,56 its strength is slightly lower than that of GO/CS, which is mainly due to the stronger synergistic interactions of covalent and H-bonding.61 Generally, strength and toughness of the materials are deemed to be mutually exclusive.1 Intrinsic toughening mechanisms are linked to plasticity and thus strength, a compromise is often achieved in engineering materials where either one of the 17


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properties is sacrificed. However, for our specially designed nacres, the organization of microstructure based on Mn+ ions through chemical cross-linking reactions, together with H-bonding interactions, on the hard and soft (GO/CMC) interface, induces the emergence of enhanced properties both in strength and toughness without compromise. On the demand of light weight and energy saving materials, the density is a dominating consideration, particularly for structural application.43,44 In spite of exhibiting relatively low stiffness (about 5.0 ~ 12.0 GPa), the specific strength (σf/ρ) of our artificial nacres is close to that of engineering metallic alloys (e.g., Titanium, Magnesium, Nickel) and some ceramic materials (e.g., Carborundum, Alumina),1,17,45, but obviously higher than those of polymeric materials and cellulose.63 The unique combination of specific strength (σf/ρ) and specific toughness (Kc/ρ) in our artificial nacres (red shadow circle) is slightly better than those metallic alloys, ceramic and natural materials as summarized in Figure 5b. Actually, owing to the dominating polymer constituent in the composite, the strength, toughness, and operating temperature of our materials are at the usual level of engineering technical polymers, highlighting their potential applications in structural, tissue engineering, wearable electronics-related fields.

CONCLUSION Inspired by the role of various minerals formed buried interfaces in nacre, a typical kind of artificial nacres (Mn+-GO/CMC) with high strength and toughness have been successfully fabricated. The measurements of macro-micromechanical properties show that the chemical cross-linking reactions between Mn+ ions and oxygen-containing groups 18


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of the two components can dramatically strengthen/toughen the GO/CMC interfaces, which significantly restrict the relative slip of between GO nanosheets and CMC layers, resulting in a great improvement of strength and toughness. However, owing to the differences in the bonding styles and different bonding energies of the M-O bonds, the reinforcing strategies of different types of Mn+ ions reflect the diverse increases in mechanical properties. For example, A2+ ions can achieve the moderate improvement in tensile strength and toughness, trivalent Al3+ possesses the outstanding tensile strength and Young’s modulus, but the zigzag long chain (TiO)n2n+ results in the high toughness with high strength. Notably, for meeting practical needs, it is also very meaningful to adjust the relationship between toughness and Young’s modulus by optimizing the Mn+ content while keeping no change of the stress. It is also worth mentioning, although the single Mn+-reinforced effect/strategy is beneficial to better understand the role of these minerals/ions in natural materials for the mechanical improvement, the complex synergic reinforcing strategy matched with the various minerals in natural materials, which can really reflect the mechanical improvement of natural materials and will become more meaningful in near future. With perspective to engineering materials, our artificial nacres realize the optimal integration of high specific strength (σf/ρ) and high specific toughness (Kc/ρ). More importantly, this special reinforcing effect opens a promising route to strengthen and toughen materials, which also is widely implicated in exploring indirectly the role of these minerals/ions in natural materials for the mechanical improvement. We believe that this investigation provides new opportunities in developing the bio-inspired composite 19


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materials with superior mechanical properties for a wide range of applications, such as aerospace, artificial muscles, tissue engineering, and wearable electronic devices.

EXPERIMENTAL Fabrication of pure GO, CMC, GO/CMC, and Mn+-GO/CMC composite papers (Mn+: Mg2+, ZrO2+, Ni2+, Ca2+, Cu2+, Co2+, TiO2+, Al3+). GO, CMC, and GO/CMC composite papers were prepared by filtering the diluted colloidal GO dispersions, aqueous CMC solution, and composite dispersions, through an anodisc filter membrane, respectively, The wet GO/CMC composite paper with 6.56 vol. % of GO could be obtained by tuning the volume ratios of colloidal GO dispersions to aqueous CMC solution, followed by addition of aqueous Mn+ solution (20 ~ 100 mL of 0.05 mM solutions). After the filtration, specimens were air-dried until the paper could be peeled off for analysis. A water-circulation multi-function vacuum pump with vacuum filter holder was utilized for vacuum filtration. Characterization. Plane/cross-sectional nacre from Clam and samples were coated with ultra-thin gold to reduce charging effects prior to scanning electron microscopy (SEM) imaging. SEM images were gained by the FEI QuantaTM 250 FEG at acceleration valtages of 5 kV and a working distance of ~ 6 - 9 mm. Backscattered electron (BSE, right) images were generated using a field emission JEOL 7500F with accelerating voltage of 20 kV and a beam current of 2.7 nA and the working distance was 6 mm. Transmission electron microscopy (TEM) with typical dark-field, and selected area electron diffraction (SAED) techniques was carried out using a JEOL JSM-2100 operated 20


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at 120 kV. Higher magnification (HR) TEM was carried out using a field emission JEOL JSM-2100 at 200 kV. The atomic force microscopy (AFM) was characterized by a Bruker Dimension Icon. All of the X-ray photoelectron spectroscopy (XPS) measurements were taken in an ESCALab220i-XL (Thermo Scientific) using a monochromatic Al-Kα X-ray source. The Raman spectroscopy measurements were taken using a LabRAM HR800 (Horiba Jobin Yvon) with a 514 nm wave-length incident laser, combined with a HYLT NV 1302 in-situ tensile test instrument (Beijing, Hyltuo Technology Co. Ltd) X-ray diffraction (XRD) experiments were carried out with a Shimadzu Lab XRD-6000 X-ray diffractometer using Cu-Kα radiation. FTIR spectra were collected using a Thermo Nicolet nexus-470 FTIR instrument. Tensile test and Nanoindentation. The tensile mechanical properties were measured using a Shimadzu AGS-X Tester with a dynamic mechanical analyzer (DMA). Static tensile tests were evaluated at a load speed of 1 mm·min-1 with a gauge length of 5 mm. Short time creep tests were conducted in the tensile mode at 27 °C with an applied stress of 30 MPa and a ramp rate of 20 µm·min-1, and the creep strain was determined as a function of time (When tensile stress was up to 30 MPa, holding for 30 mins, and then deformation recovery, waiting for 60 mins.). All of the papers for the tensile test were cut into strips with the length of 20 mm and the width of 3 mm. In the case of fracture tests, for each ligament length, at least two samples were fractured. The absorbed energy until failure was calculated by the integral area of the load-displacement. Deeply Double Edge Notched Tension (DDENT) tensile samples with a width of 30 mm, a gauge length of 20 21


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mm, and a thickness of about 25 µm were used for the Essential Work of Fracture (EWF) measurements. The initial sharp cracks were notched with a fresh zazor blade. The ligament lengths between 5 to 20 mm with an interval 3 mm were measured with a microscopy after the tests. The specimen thickness was obtained from SEM imaging of the fracture edges. Micro-mechanical properties and surface features were measured in ambient conditions (25 °C, 15 % of RH) by a commercial TI 950 triboindenter (Hyitron) equipped with a Berkovich diamond tip (R = 100 nm) using a continuous depth-sensing indentation technique and a PI-85 SEM Picoindenter (Hyitron) (See detailed methods in SI). Supporting Information Available: Materials and Detailed experimental methods for the synthesis of GO, the preparation of the dispersions, the element component and micro-structural features of natural nacre, comparison of Macromechanical data, cross-sectional SEM images, AFM images, EDS element, XRD, FTIR, XPS, Raman spectra analysis, Micro-macromechanical properties, fracture morphologies, theoretical methods, and some Movies about the deformation behaviors of the samples under in situ nanoindentation test. This material is available free of charge via the Internet at http://pubs.acs.org. Acknowledgements. We thank Juanjuan Qi and Jie Lin for designing schematic drawing of the fabrication process. This work was supported by the National Basic Research Program of China (2014CB931802), Research Fund for the Doctoral Program of Higher Education of China (20131102120053), and the National 973 Program of China (2010CB9374701). 22


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AUTHOR INFORMATION Corresponding Author E-mail:

[email protected]

&

[email protected],

fax

&

Tel:

+86-010-82338162.

Notes The authors declare no competing financial interest.

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Wan, S.; Peng, J.; Li, Y.; Hu, H.; Jiang, L.; Cheng, Q. Use of Synergistic Interactions to Fabricate Strong, Tough, and Conductive Artificial Nacre Based on Graphene Oxide and Chitosan. ACS Nano 2015, 9, 9830-9836.

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Yadav, M.; Rhee, K. Y.; Park, S. J. Synthesis and Characterization of Graphene Oxide/Carboxymethylcellulose/Alginate Composite Blend Films. Carbohydr. Polym. 2014, 110, 18-25.

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Inagamov, S. Y.; Mukhamedov, G. I. Structure and Physical-Mechanical Properties

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Based

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Carboxymethylcellulose. J. Appl. Poly. Sci. 2011, 122, 1749-1757. 64

Wegst, U. G. K.; Ashby, M. F. The Mechanical Efficiency of Natural Materials. Phil. Mag. 2004, 84, 2167-2186.

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Figure legends

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Figure 1 Hierarchical structure feature, phase dispersion, and micro-mechanical behavior in natural nacre. (a) Schematic drawing of hierarchical structure (five levels) of natural nacre from nano, micro, to structural length scales. At the nanoscale level, showing the components including aragonite crystals, β-chitin fibrils, analogous silk 33


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protein, and other inorganic minerals. At coarser levels, showing the model of the organic membrane/matrix between platelets; at the microscale levels, showing the mechanism of growth of nacreous platelets and sheet nacre. On the largest length scales, the apparent shape of the bivalve. (b) TEM image of nacre nanograins. (c) The selected area diffraction pattern corresponding selected area [rectangular area G in (a)], exhibiting the polycrystalline characteristic pattern of the aragonite. (d) HRTEM image of nacre nanograins, showing the complex grain and phase boundary. (e) The fast Fourier transform (FFT) pattern corresponding selected area H in (d). Some nanocrystals are outlined with dark yellowed dashed lines and the inter-planar distances of (021), (012), (200), (112), and (113) planes of the aragonite are observed. (f) ICP-MS element analysis and SEM-EDS element mapping of nacre (bivalve pearl oyster shell), indicating various metal elements/minerals. (g-h) The contour maps of Ec and Hc dispersion on the nacre plane corresponding to the image (b) of an array, respectively, indicating the non-uniform mechanical dispersion (P'max = 800 µN).

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Figure 2 Design strategy depicting the control of structural self-organization, and microstructural features for our artificial nacres. (a) Design fabrication process describing the strengthening and toughening control at multiple scales of structural self-assembly based on individual GO nanosheets (①), renewable CMC (②), and different cationic metals/oxygen-contained metals (Mn+), and interfacial strengthening and toughening strategy. Self-assembling of the staggered layered structural features is obtained by VAF (③). The chemically strong cross-linking interactions between the oxygen-containing groups of GO nanosheets/CMC polymers and Mn+ ions (④ and ⑤). (b) Digital image of a brown free-standing paper (e.g., Al3+-GO/CMC), showing a good flexibility. Within 10.0 µm × 10.0 µm. (c) The map of the uniform phase distribution (e.g., Al3+-reinforced) under 100 µN force by using the nanoindentation within 5.00 µm × 5.00 36


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µm surface area. (d) Tapping mode AFM height images, showing the relative flattening of nanocopic asperities on the plane of the sample (b). (e) Height profile of the red line in part d. (f-g) Cross-section morphology of the layered GO/CMC composite (f) and a typical artificial nacre paper (Al3+-reinforced) (g), indicating a layered arrangement: bottom to top: low- and high- resolution SEM images, respectively.

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Figure 3 Macro-microscopic mechanical behaviors of artificial nacres. (a) Representative stress-strain data from tensile test. (b) Graphical comparison of the ultimate stress (σg-c) and Young’s modulus (Eg-c). (c) Graphical comparison of the toughness (Wg-c). (d) Typical creep and recovery curves. (e) Typical load-displacement curves with hmax = 130 ~ 450 nm performed in the load of 800 µN (Nanoindentation points, n = 5). (f) Graphical comparison of the calculated Young’s modulus (E'g-c) and hardness (H'g-c) based on Oliver-Pharr analysis as function of maximum load (P'max) from 150 µN to 1000 µN (n > 6). (Note: G = GO, C = CMC, G/C = GO/CMC). (g) The typical contour maps of E'g-c (top) and H'g-c (bottom) dispersion on the composite papers plane within 15 µm × 15 µm (e.g., Mg2+, TiO2+, Al3+, P'max = 800 µN ), respectively.

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Figure 4 Deformation behaviors, fracture morphologies, simplified structural schematics and proposed Mn+-reinforced mechanisms in artificial nacres. (a-d) In situ SEM images of the 15 °C view of the indentation zone and multiple individual load-displacement curves for GO/CMC (a) and artificial nacres (b-d) (A flat end 5 µm 60 °C conical tip, P'max = 5000 µN, see Movies 1-4). (a) The microscale deformations (a1-d1) and load-displacement curves (a2-d2) of GO/CMC, Mg2+- TiO2+-, and Al3+-GO/CMC (b1-b3), respectively. Green arrows represent the moving direction of the tip. Loading portions (color region) of multiple individual load-displacement curves (Indentation points, n ≥ 7). (e-g) Fracture morphologies, simplified structural schematics and proposed Mn+-reinforced mechanisms. (e1) Cross-section: Mg2+, inset: ZrO2+; (e2) Ni2+; inset: Ca2+; (e3) Cu2+; inset: Co2+. (f1-g1) Plane: fracture surfaces for TiO2+ and Al3+ exhibit

long-range

crack

propagations,

respectively;

inset:

amplifying

crack

tearing/extension. (f2-g2) Plane: Adhesive layers (GO nanosheets adhered to CMC layers) are pulled out; inset: amplifying incomplete fracture layers. (f3-g3) Cross-section: fracture features; inset in f3: amplifying topical fracture morphologies. (ei-gi) Simplified structural schematics, showing the alternate arrangement of CMC polymer layers and GO nanosheets, in which anionic GO and anionic CMC polymer are interconnected by different types of Mn+ ions. (e1i-g1i) Under tensile stress, the GO nanosheets effectively transfer loading to CMC matrix through various bonding interactions and deflect cracks. (e2i-g2i) Apart from the H-bonding/van der Waals force, the chemically cross-linking forces can strengthen the GO/SA interface to restrict GO nanosheets sliding, which spread rapidly to other GO/CMC interfacial layers, accompanied by further increasing 42


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tensile stress. (e3i-g3i) The composite papers finally break under GO nanosheets pulled out modes. Note that red arrows indicate the direction of the microcrack extension and red dotted lines stand for the crack extension paths.

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Figure 5 Comparison of these related materials performance. (a) Toughness versus tensile strength for several natural materials, GO- and CMC-related film/papers. The stars represent our artificial nacres. The references are shown as follow: (Nacre,8 Bone,41 Dentine,42 1: GO/PU (Song, 2012),33 2: GO/PCL (Inoue, 2008),46 3: GO/PS (Wu, 2015),47 4: GO/PPA (Ruoff, 2009),48 5: GO/GA (Zhang, 2011),37 6: GO/(PAH-PSS) (Tsukruk, 2010),49 7: GO/PCDO (Cheng, 2013),50 8: GO/PAH (Gun’ko, 2010),51 9: GO/CS (Yan, 2011),52 10: GO/PVA (Maser, 2013),53 11: GO/MoS2-TPU (Cheng, 2015),54 12: GO/NCCA (Namazi, 2014),55 13: Fe3+-GO/TA (Xu, 2014),56 14: GO/Borate (Nguyen, 2011),4 15: GO/PAD (Cheng, 2014),38 16: GO/PBI (Wang, 2013),57 17: GO/PAPBx (Shi, 2014),58 18: GO/PVP (Shi, 2015),59 19: GO/G4NH2 (Shi, 2015),59 20: GO/PEI (Shi, 2015),59 21: GO/CS (Shi, 2015),59 22: GO/RGO-Silk (Tsukruk, 2013),60 23: GO/SF (Tsukruk, 2013),9 24: GO/CS: (Cheng, 2015),61 25: Mg2+/GO (Ruoff, 2008),35 26: Fe3+/GO (Xu, 2014),56 27: Ca2+/GO (Ruoff, 2008),35 28: GO/CMC/SA (Park, 2014),62 29: CMC/UFOs (Inagamov, 2011),63 30: CMC/MTM (Walther, 2013),26 31: GO/CMC (Guo, 2015), 32: GO/CMC (Shi, 2015).59 (b) Ashby diagram of specific strength versus specific toughness of a range of engineering and natural materials.43 The red stars in red circle refer to our artificial nacres. Fracture toughness Kc ≈ (E'·Jg-c)1/2 (MPa·m1/2), where Kc is 44


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the fracture toughness, E'g-c is the Young’s modulus (GPa), and Jg-c (kJ·m-2) is the toughness of our artificial nacres, which is equal to the essential work (we) (according to Essential Work of Fracture (EWF) approach,64 we is obtained in SI).

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Table 1 Summary of macro-mechanical improvements for Mn+-GO/CMC in comparison with GO/CMC. Materials

% (σ)

% (W)

% (E)

GO/CMC

0

0

0

Mg2+

~ 53.1

~ 373.7

~ - 14.5

ZrO2+

~ 56.0

~ 310.5

~ - 8.7

Ni2+

~ 56.0

~ 373.7

~ - 7.2

Ca2+

~ 62.6

~ 300.0

~ - 23.2

Cu2+(40)a

~ 64.6

~ 36.8

~ 34.8

Cu2+(60)a

~ 47.6

~ 410.5

~ - 20.3

Co2+

~ 73.9

~ 173.7

~ - 8.7

TiO2+

~ 76.2

~ 726.3

~ - 26.1

Al3+

~ 121.2

~ 184.2

~ 44.9

a “40” and “60” represent the theoretical content of addition of Cu2+ ions (See SI).

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Strong and Tough Layered Nanocomposites with Buried Interfaces

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