Ultrastrong Fibers Assembled from Giant Graphene Oxide Sheets

Z hen X u ,H aiyan S un ,X iaoli Z hao ,a nd C hao G ao * U

ltrastrong Fibers Assembled from Giant Graphene Oxide Sheets

D r. Z. Xu, Dr. H. Sun, Dr. X. Zhao, Prof. C. Gao MO

E Key Laboratory of Macromolecular Synthesis and Functionalization

Department of Polymer Science and Engineering Zhejiang University

38 Zheda Road, Hangzhou 310027, P. R. China E -mail: c haogao@https://www.360docs.net/doc/6c8824066.html, D OI: 10.1002/adma.201203448

H igh-performance synthetic ? bers are an integral part of modern industry and life. For example, Kevlar and carbon ? bers possess ultrahigh strength and stiffness owing to regularly aligned linear

polymer chains and planar graphitic layers. [ 1 ]High-performance

? bers have also been made from highly-aligned carbon nano-tubes (CNTs). [ 2–5 ] It is believed that graphene, the newly dis-covered two-dimensional (2D) carbon allotrope, can rival CNTs in advanced materials and electronic devices, because of its

fascinating mechanical and electrical properties. [ 6–13 ]Interest-ingly, the large-scale availability of highly soluble graphene

oxide (GO), [ 10–13 ] a precursor of chemically converted graphene,

together with its liquid crystalline behavior [ 14–20 ]has promoted

preliminary efforts to create graphene-based ? bers. [ 14 , 15 ]Com-posite ? bers consisting of graphene, CNTs, and poly(ethylene

alcohol) (PVA) demonstrated superior strength because of the synergistic effect between CNTs and graphene, but limited electrical conductivity (1.3 ×10 ? 2 S/m) due to the presence

of polymer. [ 21 ] Recently, our group has raised a wet-spinning

methodology to fabricate neat graphene ? bers from concen-trated GO liquid crystal (LC). [ 14 ] Compared with the composite

? bers, [ 21 , 22 ] the neat graphene ? bers demonstrated several orders of magnitude higher electrical conductivity but relatively lower strength. Therefore, achieving graphene ? bers with both excellent mechanical properties and high electrical conductivity is becoming an urgent challenge to be addressed.

T o meet such a challenge, we focus on improving the mechan-ical strength of neat graphene ? bers in this work. Generally, greater aspect ratio of building blocks and their better alignment in the ? ber axis are the main factors to improve the mechanical per-formance of polymer and CNT ? bers. [ 1 , 3 , 23 ] Inspired by this under-standing, we employed giant graphene oxide (GGO) [ 17 ] sheets with extremely high aspect ratios as building blocks to reduce defective edges (Figure S1 in the Supporting Information), and achieved highly ordered alignment of GGO sheets by wet-drawing liquid crystalline gel ? bers to make high-performance graphene ? bers with both high strength and conductivity. After introduction of divalent ionic cross-linking, the graphene ? bers exhibited a record strength for macroscopic, neat graphene materials. G GO sheets were made from natural graphite ? akes on a

large-scale, [ 18 , 24 ] and had good solubility in water because of their

plentiful oxygen-containing functional groups ( F igure 1a , and

Figure S2 and S3) [ 13–15 , 25 , 26 ] . The thickness of GGO sheet ( t )was

measured as ≈ 0.8 nm (Figure 1 b ), indicating the single-layered

attribute of dispersed GGO sheets in water. [ 17 , 18 ] The average

lateral size ( w ) of the GGO sheets was 18.5 μ m with a relatively wide distribution ranging from dozens of micrometers (up to 42 μ m ) to several micrometers (down to 4 μm ) (Figure 1c and

Figure S4). According to the Onsager LC theory, [ 14 , 15 , 27 ]such

high aspect ratio for GGO sheets ( w /t ≈ 2.3 ×10 4)determines very low concentration to form stable LCs. From polarized optical microscope (POM) observations, as the concentration increased to ≈ 1–4 mg/mL, stable GGO nematic LCs formed, evidenced by the spreading Schlieren textures (Figure 1 d and Figure S5). Upon further increasing the GGO concentra-tion to 8 mg/mL, vivid colors in the textures appear (Figure 1 e ), which could be attributed to the local strong birefringence. Additionally, synchrotron small angle X-ray scattering (SAXS) tests of GGO dispersions revealed the enhancing orientational orderings with increasing GGO concentrations, and the local layered structure in the GGO LCs at a concentration of 8 mg/mL answered for the emerged vivid birefringence in Figure 1 e and Figure S6. Due to the giant size of GGO sheets and the low concentration of LCs, the stable liquid crystalline spinning dopes (down to 4 mg/mL) have relatively low viscosity, which favors the lowering of ? owing friction and the continuity of the spun ? bers.

A lthough as-prepared GGO LCs hold pre-aligned orienta-tional orderings at rest, the different linear disclinations dis-tributed in three dimensional (3D) spaces (see black brushes in Figure 1 c ) could disrupt the regular alignment of GGO sheets

in the ? nal assembled materials. [ 16 ] In the case of shear ? ow, the

asymmetrical colloids (e.g., CNTs and GO) in their LCs could easily form regular alignments under the directional ? owing

? eld. [ 15 , 28–30 ] Hence, GGO sheets in LCs are able to form regular alignments along the ? owing direction, as con? rmed by POM and SEM characterization (Figure S7). This alignment of GGO sheets is helpful to get continuously assembled GGO ? bers with highly ordered micrometer-sized structures. We obtained GGO ? bers with compact and regularly aligned structures by two optimized processes, including post-drying and stretching (or wet-drawing) in the gel state. The in situ POM inspection showed that GGO ? bers without water washing are ? at ribbons and those with water washing are compact ? bers with round sections (Figure S8). We stretched the GGO gel ? bers in the coagulation baths under a ? owing ? eld (Figure 1 e ), and the cor-responding elliptical 2D SAXS pattern distinctly indicated the interior orientational ordering (Figure S6e). The resultant solid ? bers after stretching had approximately round sections and compact structures, which are advantageous to improving their mechanical properties ( F igure 2 a ,b and Figure S9–11), similar

to the spinning of polymer-contained ? bers by wet or hot-drawing. [ 1 , 21 , 30 ] The diameter of the stretched GGO ? bers could be lowered to ≈6μ m (Figure 2 a ), close to that of commercial carbon ? bers. As a comparison, the sections of the collapsed ribbons dried from the GGO gel ? bers without stretching showed crumbled morphologies (Figure 2 c ,d and Figure S12). T he section morphology of GGO ? bers in Figure 2 e,f shows a compact origami-? ower-like structure with dentate bends, which resembles the section structures of the pitch-based

carbon ? bers. [ 1 ] In GGO ? bers, GGO sheets stacked densely

with local alignments and the layered GGO sheets regions exhibited different orientational orderings around these bends. In view of the aligned structures of the ? owing GGO dopes (Figure S7), these dentate bends into the solid ? bers could orig-inate from the orientational disclinations (points or lines) in the spinning dopes. The surface of GGO ? bers showed spreading ridges and their heights were measured as ≈ 300 nm (Figure 2 g ,h). During the drying process, GGO sheets on the gel ? ber surface underwent Eulerian buckling to form the dentate folds caused

by dehydration of the gel ? bers. [ 20 ] This dehydration-responsive

folding was also observed in the drying process of deposited

GO liquid crystalline drops on glass surfaces. [ 20 ] After chem-ical reduction by hydroiodic acid (HI) aqueous solution, the

resultant RGG ? ber possessed the similar section morphology as the case of GGO ? bers. As shown in Figure 2 i –k, RGG ? ber exhibited compact folding structure as seen from its section, and the dense layered stacking can be identi? ed in the magni-? ed images.

F ollowing the optimized spinning process with stretching, GGO ? bers spun in the KOH coagulation bath had a tensile strength of 184.6 MPa at 7.5% ultimate elongation with a 3.2 GPa Young’s modulus. This value of strength is about two time greater that of our previously reported GO ? bers spun from

small GO sheets with an average size of 0.84 μ m (102 MPa), [ 14 ]

showing the mechanically enhancing effect of larger size GO sheets. Such a size-induced enhancing effect has been previ-ously reported by Shi and co-workers in the case of ? ltrated GO papers. [ 33 ] Additionally, the KOH coagulation bath can remove

the oxidation debris on GGO sheets, [ 31 , 32 ] which is helpful to

effectively assemble the clean graphene sheets to compact struc-tures and to promote the mechanical strength of the ? bers.

I n the pursuit of high-performance graphene-based mate-rials, previous experimental and theoretical investigations pro-posed that divalent ions offered interlayer and intralayer cross-linking bridges between the oxygen containing groups and thus brought the enhancement in the mechanical properties to GO

papers/? lms. [ 34 , 35 ] We chose CaCl 2 and CuSO 4 solutions as the

coagulation bath solutions to improve the strength of graphene

? bers. As presented in F igure 3a and T able 1,Ca 2 + -cross-linked GGO ? bers had a tensile strength of 364.6 MPa at 6.8% ulti-mate elongation, and the ? bers cross-linked by Cu

2 + exhibited a tensile strength of 274.

3 MPa at 5.9% elongation. Compared

with the GGO ? bers spun in KOH solution, Ca 2 + -cross-linked

GGO ? bers doubled the tensile strength (364.6 compared with 184.6 MPa) and the Young’s modulus (6.3 compared with 3.2 GPa). The elemental mapping analyses manifested that Ca was homogenously distributed through the whole ? ber (Figure S13), con? rming that the enhancement in mechanical properties is mainly attributed to the cross-linking of divalent ions between GGO sheets. Additionally, all GGO ? bers without wet-drawing showed lower tensile strengths (125.0–258.6 MPa),

about 33% lower in the case of Ca 2 + -cross-linked GGO ? bers,

verifying our effective control over the performance of spun GGO ? bers (Figure S15 and Table S1).

T hrough chemical reduction by hydroiodic acid, the as-prepared GGO ? bers were turned to RGG ? bers by partly removing the original oxygen containing groups on GGO and

restoration of conjugated nets. [ 35 ]Signi? cantly, the resultant

RGG ? bers showed further enhanced strengths and Young’ s moduli (Figure 3 a and Table 1 ). For example, the reduced coun-terpart (Ca 2 + -cross-linked RGG ? ber) had a tensile strength of 501.5 MPa and Young’s modulus of 11.2 GPa, which are 41% and 70% higher than those values of original GGO ? bers. More-over, the enhancements in both strength and stiffness took no

F igure 1. a ) The photograph of 0.6 L GGO liquid crystalline aqueous dispersion at a concentration of 4 mg/mL. b) AFM image of GGO sheet deposited

on mica. c) SEM image of GGO sheets deposited on silicon. POM images of GGO liquid crystalline gels loaded in the planar cells, at concentrations of d) 4 mg/mL and e) 8 mg/mL. f) Schematic apparatus for spinning GGO ? bers. “o” indicates the rotating center, “St 1.0” and “St 1.3” indicate the locations of nozzle with distances of 2 cm and 1.5 cm to the rotating center, respectively. g) A ? ve-meter long GGO ? ber wound on a ceramic reel.

50 μm

0.8 nm

2 μm 200 μm 200 μm

b d e

Coagulation bath

and the remaining oxygen-metal cross-linking bridges. This

explanation of the cross-linking mechanism is supported by the residual oxygen functional groups (ca. 12.5 at% O) in the X-ray photoelectron spectroscopy (XPS) spectrum (Figure S16) and the homogeneous distribution of Ca (ca. 0.6 at%) in the EDS Ca-element mapping image (Figure 2 i and Figure S17).

negative effect on the ultimate elongation (6.7% compared to 6.8%), indicating a signi? cantly enhanced toughness of the RGG ? bers. We suggest that these enhancements in mechan-ical properties by chemical reduction could originate from the decreasing interlayer space between RGG sheets ( ≈0.37 nm),

as demonstrated by XRD measurements (Figure S15), [ 35 ] F igure 2. S EM images of ? ber sections spun in a) 5 wt% CaCl 2 and b) 5 wt% NaOH ethanol/water solutions with 1.3-fold stretching, and c,d) their corresponding ? ber section SEM images without stretching. SEM images of e,f) typical fracture and g,h) surface morphology of GGO ? bers, and

i–k) section morphology of RGG ? bers. l) SEM images of Ca 2 +

-cross-linked RGG ? ber (top), C-element mapping (middle) and Ca-element mapping (bottom). F igure 3. a ) Typical mechanical measurements under tension for GGO ? bers and RGG ? bers. b) Diagram of mechanical performance data for graphene-based neat papers and ? bers in previous reports and in this study. The hollow symbols indicate the neat GO-based ? bers/papers and the solid symbols indicate the reduced graphene-based ? bers and papers. The red pentagrams denote GGO and RGG ? bers in this work. SEM images of

fracture surfaces of c,d) the GGO-KOH ? ber, e–h) theGGO-Ca 2 + ? ber and –RGG-Ca 2 +

? ber. The dashed lines in (g) indicate either the boundary lines or the cracking lines of the pulled-out graphene sheets. k) The deformation mechanism model of GGO/RGG ? bers under tensile stress. The dashed lines indicate the hydrogen bonds and coordinative cross-linking bridges.

01234567891011

100200300400500RGGF KOH GGOF-Ca

RGGF-Ca 2+

RGGF-Cu

GGOF-Cu

T e n s i l e s t r e n g t h (M P a )

Strain (%)

GGOF KOH

T e n s i l e s t r e n g t h (M P a )

Breakage elongation (%)

10 μm 4 μm 500 nm

5 μm 500 nm

200 nm

c d

e f

i j

5 μm

graphene ? lms/papers with smoothly layered structures, more wrinkles of layered graphene sheets in graphene ? bers can contribute additional forces to their mechanical performance,

which could explain the fact that graphene ? bers generally pos-sess stronger tensile strength and higher break elongation than

? lms/papers.

A s shown in Figure 3 c –j, the tensile fractures of GGO and RGG ? bers display a certain degree of elongation in the ? ber axes, indicating their elastic breakage nature. Upon

closer analysis, the typical stretched ? brils can be identi? ed

at the fracture tips, and such a fracture feature resembles the “pull-out” characteristic of neat CNT ? bers under tensile

stress. [ 5 ] We further identi? ed the distinct boundary lines of the pulled-out graphene sheets at the fracture tips, which may originate from either the edges of stacked graphene or the

cracking lines of graphene sheets. In the frame of the tension-shear model of graphene ?

bers (Figure 3 k ), the constituent graphene sheets endure a pulling force to slide from the

stacked graphene blocks under the tensile force in the ?

ber axial direction. Therefore, these observed boundary lines of

graphene sheets support a tension-shear deformation mecha-nism under tension in graphene ? bers. Because of their good ?

exibility, graphene ? bers can be twisted together to obtain continuous threads and yarns (

F igure 4 a –c), which could be useful for making ?

exible and wearable sensors and superca-pacitors devices.

A fter chemical reduction, RGG ? bers showed excellent elec-trical conductivities in the range of 3.8–4.1

×10 4 S/m, about 4 times higher than that of reduced graphene papers (7.2

× 10 3S/m) [ 38 ] and about 40% higher than that of graphene ?

bers assembled from small graphene sheets (2.5

×10 4S/m), [ 14 ] because of the giant size of RGG sheets and their regular align-ment in the ?

bers. To investigate the electrical resistance of RGG ?

bers under bending, the ? ber was ? xed on an insulating tweezer by silver conducting paste and cycled bending was

operated by hand (Figure

4 d inset). During 1000 times bending, the conductivities of RGG ?

ber were stable and showed a negli-gible ? uctuation (Figure 4 d ). The high conductivity of the RGG ? bers together with their good ? exibility makes them useful as ?

exible and lightweight cables in wearable electronic devices. To further study the electrical transport in RGG ?

bers, we inves-tigated the temperature ( T )-dependent conductivity ( σ ) in the range of 5–300 K. The RGG ?

ber behaved as a typical semicon-ductor (i.e., a positive d

σ /d T ) in the investigated temperature range [ 41 ] and the conductivity decreased from 4.1

×10 4 S/m at 299 K to 1.8

×10 4 S/m at 5 K (Figure 4 e ). Simulation indicates that the electrical conduction in the RGG ? ber system is con-trolled by the hopping mechanism, which is in accordance with

the observed conductive mechanism in other neat graphene materials (Figure S18).

[ 42 ] I n conclusion, we designed and fabricated ultrastrong graphene ? bers with GGO sheets by wet-drawing and ion-cross-linking. The giant size of the constituent graphene sheets

together with their good alignment resulted in a considerable improvement in the mechanical performance of the graphene

? bers. The graphene ? bers cross-linked by divalent ions pos-sessed a record tensile strength (up to 0.5 GPa) among neat

graphene materials, with excellent electrical conductivity. Such

multifunctional graphene ?

bers have promise in versatile

T he giant sizes of GGO sheets and their regular alignments by the optimized spinning process were designed to translate the outstanding mechanical properties of graphene sheets into macroscopically assembled ? bers. Our GGO and RGG ? bers possessed higher strengths compared with the graphene-based neat papers and ? bers reported previously (Figure 3 b ). [ 34 , 36 ? 41 ] The GGO ? ber coagulated in KOH solution without coordi-nated cross-linking (GGOF-KOH) possessed about 1.5 times (184.5 MPa) greater tensile strength than the ? ltrated GO papers (125 MPa), [ 37 ] and its reduced ? ber (RGGF-KOH) even showed higher strength (303.5 MPa) than the annealed graphene papers (293 MPa). [ 38 ] The Ca 2 + -cross-linked GGO ? ber had 2.8 times

(364.4 MPa) tensile strength of the Ca 2 + -modi? ed GO papers (125.8 MPa) [ 34 ] and the borate-modi? ed GO ? lms (127 MPa). [ 39 ] In fact, the tensile strength of Ca 2 + -cross-linked RGG ? ber (501.5 MPa) sets a new record for the strongest neat graphene materials, which is about 20% higher than the highest ever-reported value of annealed graphene ? ber with possible covalent cross-linking (420 MPa). [ 40 ] Notably, our spun GGO and RGG ? bers had much higher breakage elongation ( ≈ 2–10 times) than the reported neat graphene ? lms/papers, which implies good ? exibility and higher toughness against fracture under tensile stress. This also suggests that the mechanical strength and stiffness of graphene ? bers could be highly improved by fur-ther modi? cation of the spinning process and introduction of stronger interactions such as covalent bonds between graphene sheets.

T he deformation mechanism of these GGO and RGG ? bers can be described by the tension-shear model, a prevalent theory for nanocomposites. [ 30 , 31 ] In the neat graphene system, there are three kinds of dominant interactions between graphene sheets, including van der Waals interaction, hydrogen bonds and coor-dinative cross-linking. In GGO ? bers, the hydrogen bonds dominate and contribute to their mechanical strength. [ 37 ]After

chemical reduction, the increasing van der Waals interaction along the decreasing interlayer space and the hydrogen bonds

between residual oxygen functional groups caused the enhance-ment in the mechanical properties. Additionally, the divalent ions bridged oxygen-containing groups on the GGO sheets and at the boundaries by formation of coordination bonds. For neat graphene materials, the interlocking effect between the out-of-plane structures can promote their mechanical strengths and these interlocking structures could have resulted from twisting and wrinkling of the graphene layers. [ 37 ] Compared with the T able 1. M echanical Properties and Electrical Conductivities of stretched GGO and RGG Fibers.

Sample Tensile strength [MPa]Break elongation [%]Young’s Modulus [GPa]Conductivity [S/m]

GGOF-KOH 184.67.5 3.2–GGOF-Cu

2 + 274.

3 5.9 6.4–GGOF-Ca

2 + 364.4 6.8 6.3–RGGF-KOH 303.5 6.4 6.1 3.9 ×10 4

RGGF-Cu 2 + 408.6 6.08.6 3.8 ×10 4 RGGF-Ca 2 + 501.5 6.711.2 4.1 ×10 4

[ 1]J . W. S. H earle ,H igh-Performance Fibers , W oodhead Publishing Ltd.,

C ambridge 2000.[ 2]B . V igolo ,A . P énicaud ,C . C oulon ,R . P ailler ,C . J ournet ,P . B ernier ,P . P oulin ,S cience 2000,290,1331 .[ 3]K . K oziol ,J . V ilatela ,A . M oisala ,M . M otta ,P . C unniff ,M . S ennett ,A . W indle ,S cience 2007,318,1892 .[ 4]K . L. J iang ,Q . Q. L i ,S . S. F an ,N ature 2002,419,801 .[ 5]W . B. L u ,M . Z u ,J . H. B yun ,B . S. K im ,T . W. C hou ,A dv. Mater. 2012,

24,1805 .[ 6]K . S. N ovoselov ,A . K. G eim ,S . V. M orozov ,D . J iang ,Y . Z hang ,S . V. D ubonos ,I . V. G rigorieva ,A . A. F irsov ,S cience 2004,306,666 .S upporting Information

upporting information is available from the Wiley Online Library or from the author.

A cknowledgements W

e thank Prof. F. G. Bian, Dr. X. H. Li, Dr. F. Tian, Dr. W. Q. Hua in SSRF (BL16B1 experimental station) for SAXS characterization. This work was supported by the National Natural Science Foundation o

f China (No. 20974093 and No. 51173162), Qianjian

g Talent Foundation of Zhejiang Province (No. 2010R10021), Fundamental Researc

h Funds for the Central Universities (No. 2011QNA4029), Research Fund for the Doctoral Program of Higher Education of China (No. 20100101110049), and Zhejiang Provincial Natural Science Foundation of China (No. R4110175).

eceived: August 20, 2012

R evised: September 10, 2012Published online: October 9, 2012

applications such as functional textiles, ? exible and wearable sensors, and supercapacitors devices. The realization of strong ? bers composed of giant graphene sheets opens the door for the next-generation of high-performance ? bers with superb strength, excellent toughness, and rich functionalities fabricated by a room-temperature supramolecular assembly strategy.

E xperimental Section

W et-Spinning for GGO Fibers : For a detailed description of the

materials used, preparation of GGO and the instrumentation see the Supporting Information. Following the wet-spinning protocol depicted in Figure 1 f , GGO spinning dopes (5.0 mg/mL) were injected into the rotating coagulation baths (20 rpm/min) at a rate of 100 μL/min (about 2 m/min). The chosen coagulation baths were ethanol/water (1:3 v/v) solutions of 5 wt% KOH, 1 wt% CaCl 2 , 5 wt% CaCl 2,and 5 wt% CuSO 4 . After 30 min immersion in coagulation baths, the GGO gel ? bers were transferred into the water bath to wash away the residual coagulation solution, and the washed GGO ? bers were collected onto the bracket. The wet-drawing of gel ? bers in the spinning process was realized by adjusting the location of the nozzle. For example, spinning at the location S-1.3 (about 2 cm from the rotation center) yielded 1.3-fold stretched GGO ? bers; spinning at the location S-1 (about 1.5 cm from the rotation center) yielded the GGO ? ber without stretching. The wet ? bers were dried at 60 °C in air for 1 h and then dried at 60 ° C under vacuum for 12 h, to give the ? nal GGO ? bers (Figure 1 f ).

C hemical Reduction of GGO Fibers : Dried GGO ? bers were immersed into the hydriodic acid solution (30%) and kept at 80 ° C for 12 h. After cooling to room temperature, the ? bers were washed by water and ethanol in succession and dried at 100 ° C under vacuum for 12 h. Through this chemical reduction process, GGO ? bers were turned to reduced giant graphene (RGG) ? bers.

F igure 4. a ) Photograph of the process to twist graphene ? bers to yarns. b,c) SEM images of the twisted yarn at different magnitudes. d) Electrical

resistance of RGGF under cycled bending. The insets are photos of straight (left) and bending (right) RGGF during the test. e) Plot of temperature-dependent conductivity of RGGF in the temperature range of 5–300 K. The inset is the sketched four-probe apparatus of the conduction test.

1x10

2x10

3x10

4x10

C o n d u c t i v i t y (S /m )

Temperature (K)

R /R 0

Bending Times

RGGF

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