SC I ENCE ADVANCES | R E S EARCH ART I C L E

MATER IALS SC I ENCE

1Department of Mechanical and Aerospace Engineering, University of Virginia,122 Engineer’s Way, Charlottesville, VA 22904-4746, USA. 2Department ofMaterials Science and Engineering, Northwestern University, 2220 Campus Drive,Evanston, IL 60208-3108, USA. 3Northwestern University Center for Atom-ProbeTomography, Northwestern University, Evanston, IL 60208, USA.*Corresponding author. Email: xl3p@virginia.edu

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Bioinspired, graphene-enabled Ni composites with highstrength and toughnessYunya Zhang1, Frederick M. Heim1, Jamison L. Bartlett1, Ningning Song1,Dieter Isheim2,3, Xiaodong Li1*

Nature’s wisdom resides in achieving a joint enhancement of strength and toughness by constructing intelligent, hi-erarchical architectures from extremely limited resources. A representative example is nacre, in which a brick-and-mortar structure enables a confluence of toughening mechanisms on multiple length scales. The result is anoutstanding combination of strength and toughness which is hardly achieved by engineering materials. Here, abioinspired Ni/Ni3C composite with nacre-like, brick-and-mortar structure was constructed from Ni powders andgraphene sheets. This composite achieved a 73% increase in strength with only a 28% compromise on ductility,leading to a notable improvement in toughness. The graphene-derivedNi-Ti-Al/Ni3C composite retained high hard-ness up to 1000°C. The present study unveiled a method to smartly use 2D materials to fabricate high-performancemetalmatrix compositeswith brick-and-mortar structure through interfacial reactions and, furthermore, created anopportunity of developing advanced Ni-C–based alloys for high-temperature environments.

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INTRODUCTIONA joint enhancement of strength and toughness is a vital requirementfor next-generation structural materials. Unfortunately, this pursuit of-ten ends with a compromise between hardness and ductility (1). Such adilemma originates from the fact that the size of the plastic deformationzone in front of the crack tip, which works to dissipate local stress, isinversely proportional to the yield strength (2). Moreover, in most en-gineering materials, once fracture is initiated, cracks propagate rapidlywithout any shielding behind the crack tip (3). The wisdom of natureaddresses this conflict by constructing materials with a hierarchicalarchitecture, all thewhile using only limitedmaterials andnontoxic pro-cesses (4). A common example is nacre, or mother-of-pearl. As oneof the most well-known natural armors, nacre is endowed by a brick-and-mortar structure composed of aragonite (amineral form of CaCO3)platelets and biopolymer (5). Acting as themajor load bearers, aragoniteplatelets (the bricks) that are 5 to 10mmin length and 0.5mmin thicknessare constructed by nanocrystals (5). The biopolymer (themortar), whichis only several nanometers in thickness, closely binds the aragonite pla-telets together (6). Such a complex structure enables multiple extrinsictoughening mechanisms on different length scales, leading to anoutstanding combination of strength and toughness that is hardly seenin engineered materials. The layer-by-layer architecture redirects thecrack growth into a tortuous path, effectively consuming fracture energyvia extension of the crack length and reduction in stress concentration(7). In addition, mineral bridges shield the crack opening (8), while bio-polymer layers dissipate fracture energy (9). In the meantime, surfacenanoasperities interlock the aragonite platelets, preventing large-scaledelamination (10). While entirely mimicking the reinforcing multitudeof these scale levels is difficult, the hierarchical architecturemay hold thekey to suppressing the dilemma between strength and toughness. There-fore many researchers have gained inspiration from these biodesigns forhigh-performance composite materials. Similar to nacre, ceramic com-posites with additional soft phases such as polymers or metals have

shown that brittle ceramic materials can be converted into tough ma-terials via architecture design (11–18). Bioinspired polymer compositeswith added hard ceramic platelets exhibited high strength, outperform-ingmost engineering polymers, while retaining ductility (19–23). Thesestudies are nothing short of remarkable; the toughness of such compo-sites was magnitudes higher than the simple mixture of constituents.However, the intrinsically low ductility of ceramics and the low strengthof polymers limit the overall potential mechanical performance. More-over, weak bonding between hard phases and soft phasesmay also lead tointerface delamination. Therefore, it can be expected that cloning nacre’sarchitecture with stronger constituents such as metals in engineeredcomposites is a more promising, as well as a more challenging, task.

Previously, ceramics and intermetallic compounds have been usedas hard phases in constructing metal-based composites with brick-and-mortar structure, which exhibited notablemechanical properties (24–27).Graphene, a single layer of carbon atoms with sp2 bonds, is consideredan ideal reinforcing agent for metal matrix composites because of itstwo-dimensional (2D) morphology, ultrahigh elastic modulus of 1 TPa,and high strength of 140 GPa (28). So far, graphene has been compositedwith metals such as Al, Cu, and Ni for the construction of laminatedstructures (29–34). However, the agglomeration and degradation of gra-phene sheets, as well as the poor bonding between graphene and themetal matrix, or the occurance of unexpected reactions, resulted in amuch lower than expected reinforcement efficiency.Therefore, unlockingthe true potential of graphene in metal composites is still unresolved. Analternative approach is to use the unique morphological features of gra-phene, inducing interfacial reactions to form nacre-like, brick-and-mortar architecture in composites (29). Ni andNi alloys are widely usedin different applications, especially in high-temperature, extreme en-vironments such as combustion engines or turbine blades because oftheir outstanding mechanical performance and stability (35). Con-sidering the substantial significance of developing advanced Ni alloyswith superior properties, several studies have attempted to compositegraphene with Ni, but the high solubility of carbon in Ni and the ten-dency of forming coarse Ni3C particles (36) required electrochemicaldeposition (32), spark plasma sintering (33), and/or laser sintering(34) as the only methods to sinter Ni/graphene composites. The ob-tained composites, although exhibiting superior hardness, are unlikelyto bemass-produced. Therefore, the question remains: Canwe fabricate

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graphene-enabled, high-performanceNimatrix composites with nacre-like, brick-and-mortar structure via feasible and scalable procedures?

Here, a graphene-derived Ni/Ni3C composite with a characteristicbioinspired, brick-and-mortar architecture was fabricated by con-ventional powder metallurgy. Ni powders were homogeneously coatedwith graphene by shear mixing and freeze drying. At high temperature,carbon dissolved into Ni, facilitating the sintering process. Subsequently,part of the carbon atoms reacted with Ni, forming Ni3C second-phaseparticles along grain boundaries. The Ni3C second-phase particleswere deformable and aligned into thin, long stripes, forming a brick-and-mortar structure via cold rolling. Another portion of carbon re-mained in the Ni matrix as interstitial solid solution atoms. The Ni3Cplatelets served as major load bearers and strengthened the composite,while the Ni matrix ensured ductility. Because of the confluence ofstrengthening and toughening mechanisms, the fabricated compositeexhibited a 73% improvement in strength and only a 28% reductionin ductility, leading to a notable enhancement of toughness. By add-ing 2 weight % (wt %) of Ti and Al, the graphene-derivedNi-Ti-Al/Ni3C composite exhibited a high hardness up to 1000°C, indicativeof possible advancedNi-based superalloys. The 2Dmaterial–enabledpowder processing can be applied to different material combinations,creating unlimited possibilities for new metal matrix composites.

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RESULTS AND DISCUSSIONFabrication of graphene/Ni powdersShear mixing was used to produce few-layered graphene and coat thegraphene sheets on Ni particles simultaneously, largely improving theproduction efficiency. Specifically, 1.5 g of graphite was dispersed in200 ml of dilute water and shear-mixed at 3000 rpm for 1 hour. The re-lationships between the shear mixing rate, duration, and the amount ofgraphene productionwere thoroughly studied by Paton et al. (37). Aftershear mixing, the coarse, undefoliated graphite powders were subsidedand removed, leaving few-layered graphene in the liquid. Under the cur-rent conditions, the weight of the fabricated graphene decreased from1.5-g graphite to about 0.09 to 0.12 g. The obtained graphene sheetsexhibited thin, flexible morphology (fig. S1A) with an almost intactcrystal structure (fig. S1, B and C). Without further treatment, thegraphene-containing suspension was shear-mixed with 5-g Ni powderswith irregular shape for 2 hours, corresponding to about 2 wt % andabout 10 atomic % (at %) of carbon. Subsequently, the Ni/graphenepowders were freeze-dried for 6 hours. After these procedures, scanningelectron microscopy (SEM) observation showed that the graphenesheets were closely coated over the Ni powders without noticeableaggregation (Fig. 1A). Transmission electron microscopy (TEM) in-spection unveiled that the graphene had closely coated the Ni powderswithout free space (Fig. 1B). The thickness of the graphene coating layerranged from 10 to 15 nm, equaling to 20 to 30 atomic layers. An in situheating observation revealed that the graphene sheets gradually dis-solved into the Ni particle even without any compression (Fig. 1C),which is indicative of a very intimate bonding between graphene andNi. Ni powders with 4 and 6 wt % graphene were also prepared as re-ferences. Both shear mixing and freeze drying are critical for homoge-neously coating graphene on Ni powders. Without shear mixing,graphene sheets were loosely attached on the Ni powders and couldnot permeate into the small gaps or crevices on the particle surface(fig. S2A). Without freeze drying, graphene sheets agglomerated intocoarse graphite particles (fig. S2B). The method can be feasibly adaptedto other materials and may find more applications in various fields.

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Microstructure and mechanical performance ofgraphene-enabled Ni/Ni3C compositeThe Ni particles with 2 wt % graphene produced by shear mixing andfreeze drying were compressed in a round mold that was 25 mm in di-ameter to a pressure of 80 MPa and then sintered at 1450°C under theprotection of argon, slightly lower than themelting point of Ni (1455°C).After sintering, the Ni/graphene powders were melted together, showingno evidence of pulverization or fracture (fig. S3A, inset). After chemicaletching, the Ni/graphene powder–derived ingot exhibited clean, well-defined grain boundaries, indicative of the formation of grain boundaryprecipitates (fig. S3A). Ni powders without graphene and with 4 and6 wt % of graphene were also compressed and sintered under the samecondition as references. The Ni powder–derived ingot displayed no ob-vious grain boundaries after chemical etching (fig. S3B). Sinteringenabled the Ni powders with 4 wt % graphene to form a coin, whichwas broken into pieces after minor deformation, indicative of weakbonding between particles (fig. S4A). The Ni powders with 6 wt % ofgraphene could not be sintered together (fig. S4B). Close-up inspectionof the sintered sample with 4 wt % graphene showed a discontinuousmicrostructure with isolated particles (fig. S4C). Energy-dispersivex-ray spectroscopy (EDS) carbon map uncovered that carbon seg-regated on the particle surfaces (fig. S4D). According to the Ni-C phasediagram, the highest solubility of carbon in nickel is 2.7 at %, and theeutectic reaction point is located at 10 at % of carbon (36), which cor-responds to about 2 wt %. When 4 wt % graphene (or 20 at %) wascomposited with Ni, some carbon remained on the Ni particle surfaceseven though part of theNi powdersmelted. The excessive carbon left onthe Ni particle surfaces prohibited the sintering of Ni powders.

Cold rolling with a 40% deformation reduction in thickness was firstapplied to the sintered coins. The grain boundary precipitates werebroken into long, thin stripes, which gradually aligned along the rollingdirection during deformation. After the deformation reduction in thick-ness was increased to 80%, the boundary precipitates were aligned in a

Fig. 1. Ni/graphenepowders after shearmixingand freezedrying. (A) SEM imageof Ni/graphene powders, showing no noticeable aggregation of graphene sheets.(B) TEM image of the surface of a Ni/graphene powder, showing that few-layeredgraphene closely coated around the Ni particle. (C) In situ heating observation of a Ni/graphene powder. Graphene gradually dissolved into Ni with increasing temperature.

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parallel fashion, forming a brick-and-mortar architecture (Fig. 2A) [RD(rolling direction), TD (transverse direction), and ND (normal direc-tion)]. The fracture surface also exhibited a laminated feature withelongated dimples (Fig. 2B), which had the same shape as the second-phase particles in Fig. 2A. TEM inspection unveiled large second-phaseparticles embedded within thematrix (Fig. 2C). The dislocation densitywithin the second-phase particles was lower than that of the Ni matrix.In addition, the grain size and shape of theNimatrixwere different nearand beyond the second-phase particles. Close-up inspection of the Nimatrix exhibited stripe-like grains with a thickness ranging from 100 to

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300 nm, a typical cold-rolledmicrostructure (Fig. 2D). Two-beamdark-field imaging revealed a high concentration of dislocations, and theirmigration was prohibited (Fig. 2E), indicative of the existence of preci-pitates or solid solution atoms. In contrast, Ni grains near the largesecond-phase particle were smaller in size with an equiaxed morphol-ogy. The difference in grain size and shape may derive from higherdeformation energy, which may stimulate dynamic recrystallization.High-resolution TEM (HRTEM) imaging of the boundary showed nonoticeable defects such as voids or cracks (Fig. 2G), indicating that thecold deformation did not break the bonding between the precipitates

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Fig. 2. Microstructure of the graphene-enabled Ni/Ni3C composite. (A) SEM image of cold-rolled Ni/Ni3C composite, showing brick-and-mortar structures. (B) Fracturesurface of Ni/Ni3C composite, showing laminated structure constructed by elongated dimples. (C) Low-magnification TEM image, showing a large second-phase particleembedded in the Ni matrix. (D) After cold rolling, Ni grains were deformed into long stripes with the thickness ranging from 100 to 300 nm. (E) Two-beam diffractiondark-field image of the Ni matrix, showing a high concentration of dislocations. (F) Close-up observation of the Ni/Ni3C boundary. (G) HRTEM image of the interface betweenNi and a second-phase particle, revealing a transition zone. (H) Ni3C crystal on [−110] plane. (I) HRTEM image of the [−110] plane of Ni3C particle, showing identical atomicarrangement as in the Fig. 2H. (J) Schematic illustration of the formation of Ni/Ni3C composite with a brick-and-mortar structure.

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and the matrix. Instead, there was an amorphous transition zone be-tween the Ni matrix and the second-phase particles, suggesting thatthe second-phase particles precipitated out from the Ni matrix.

A critical question then arises:What is the second-phase particle? Thex-ray diffraction (XRD) spectrumof theNi/graphene-derived compositeexhibited a weak peak at about 43° (fig. S5A), which was similar to theNi3C (002) peak illustrated in (36). To determine the composition of thesecond-phase particles, EDS analysis and HRTEM observation wereused. EDS spectra showed the existence of Ni and carbon with a Ni/Cratio of about 3:1 (fig. S5B). Therefore, we can likely assume that thesecond-phase particle is Ni3C. Ni3C is a stable nickel/carbon compoundbelow 300°C. It has a hexagonal close-packed structure with latticeparameters of a = 0.26 nm and c = 0.43 nm. The unit cell of Ni3Cplotted by Materials Studios is shown in fig. S6A, in which Ni andC atoms are arranged layer by layer. According to the crystal structure,on the [−110] plane of the Ni3C particle, Ni atoms should exhibit visualdistances of 0.23 and 0.26 nm (fig. S6B and Fig. 2H) under HRTEM.The HRTEM image (Fig. 2I) and corresponding fast Fourier transform(FFT) pattern (fig. S6C) exactlymatched the calculation. Therefore, bycoupling the results from XRD, EDS, and HRTEM, the large second-phase particles are conclusively determined to beNi3C. A close [01–1]//[−110] orientation relation between Ni and Ni3C can be derived fromFig. 2 (G and I), which may lead to the weak XRD peaks. Armed withsolid experimental results, we can conclude the formation mechanismof the bioinspired Ni/Ni3C composite with brick-and-mortar struc-ture. As illustrated in Fig. 2J, Ni powders were closely wrapped bygraphene after shearmixing and freeze drying. The graphene sheets dis-solved into the Ni matrix and precipitated out as Ni3C in the sinteringand cooling. TheNi3C second-phase particles were broken and elongatedinto long strips during cold deformation, forming the brick-and-mortarstructure.

Encouragingly, this graphene-derived Ni/Ni3C composite withbrick-and-mortar structure exhibited outstanding mechanical per-formance. The tensile specimens were fabricated into a dog bone shapewith a sample size demonstrated in Fig. 3A. Tensile tests were carriedout at a strain rate of 0.06 mm/min. The strain was measured by Cor-related Solutions’s Vic-2D digital image correlation system with onephoto every 2 s. Three tests were conducted on each group of samples.In a typical tensile test, the Ni/Ni3C composite showed a yield strengthof 780 GPa and an ultimate tensile strength of 1095 GPa (Fig. 3A). Thisstrength is comparable to the strongest Ni-based alloys (38). Although

Zhang et al., Sci. Adv. 2019;5 : eaav5577 31 May 2019

the strength had an obvious improvement, the ductility only exhibited aminor decrease. The Ni/Ni3C composite has a joint enhancement ofmechanical properties: A 73% improvement in strength, a 6% increasein Young’s modulus, and an 82.3% enhancement of hardness with onlya 28% reduction in ductility (Fig. 3B). Such outstanding combinationbetween strength and ductility resulted in a 44% increase in toughness(the area under the strength/strain curve) than the pure Ni referencesample, indicating that the bioinspired, brick-and-mortar architectureeffectivelymitigated the conflict between strength and toughness. So far,a large number of bioinspired, nacre-like composites including ceramic-based (diamond marks in Fig. 3C) (11–18), polymer-based (roundmarks in Fig. 3C) (19–23), and metal-based (square marks in Fig. 3C)(24–27, 29–34) composites have been developed. Freeze casting hasbeen widely used to construct laminated ceramic composites (11–13);polymer cross-linking, on the other hand, is the primary route to fab-ricate layer-by-layer, polymer-based composites (20, 21); compressingsintering, powder processing, electrochemical deposition, and lasersintering have been used to synthesize nacre-like, metal-based compo-sites (29–34). The soft phases (polymers andmetals) with a volume frac-tion ranging from 5 to 40% in the nacre-like, ceramic compositeseffectively improved the ductility of the composites, with a trade-offof reduction in the characteristically high strength of ceramics (dia-mond marks in Fig. 3C). The hard phase (mainly ceramic flakes) innacre-like, polymer-based composites prohibited the decoiling of poly-mer chains, which notably enhanced the strength of the polymermatrices (round marks in Fig. 3C). In the nacre-like, metal-based com-posites, the volume fraction of metal matrix, which is normally con-sidered as the soft phase, is often more than 70%. Hard phases withplatelet-like morphology, such as graphene, ceramics, and intermetalliccompounds, were homogeneously dispersed in the metal matrices, facil-itating a joint enhancement of strength and toughness (square marks inFig. 3C). Because of the intrinsically high strength of Ni and the con-structed brick-and-mortar architecture, the graphene-enabled Ni/Ni3Ccomposite outperformed most other nacre-like composites in termsof the combination of yield strength and ductility (starmark in Fig. 3C).

Stiffening, strengthening, and toughening mechanismsIt is important to understand the stiffening, strengthening, and tough-ening mechanisms of the graphene-derived Ni/Ni3C composite. TheYoung’s modulus of the composite was slightly higher than that of pureNi (Fig. 3B). Apparently, the enhancement of Young’s modulus derives

Fig. 3. Mechanical properties of graphene-enabled Ni/Ni3C composite with a brick-and-mortar structure. (A) Tensile stress-strain curves of Ni, Ni produced bypowder metallurgy, and Ni/Ni3C composite (inset shows the size of tensile specimen). (B) Comparative bar chart of mechanical properties of Ni and Ni/Ni3C composite.(C) Elongation versus yield strength plot showing that the as-fabricated Ni/Ni3C composite had an outstanding combination of strength and ductility (mechanicalproperties of nacre-like composites were derived from (11–27, 29–34).

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from theNi3Cplatelets. A volume fraction of 13.3%of theNi3Cplateletswas calculated from five low-magnification SEM images on ND/TDand ND/RD planes, respectively. Low-load nanoindentation tests werecarried out to identify the mechanical properties of the Ni3C platelets.As shown in a typical nanoindentation displacement-load curve, theNi3C platelets exhibited higher hardness and reduced modulus

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(Fig. 4A). On average, the Ni3C platelet has a hardness of 6.5 GPa(3.4 GPa higher than that of Ni matrix) and an elastic modulus of364 GPa (154 GPa higher thanNi). The hardness (Fig. 4B) and reducedmodulus (Fig. 4C) maps derived from nanoindentations exhibitedan alternating hard-soft-hard structure with the hard part as theNi3C platelets. Because of the linear nature of elastic deformation

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Fig. 4. Strengthening and toughening mechanisms of graphene-derived Ni/Ni3C composite with brick-and-mortar structure. (A) Nanoindentation load-displacement curves of Ni and Ni3C platelet. (B) Hardness map derived from nanoindentation tests. (C) Reduced modulus map derived from nanoindentation tests.(D) Finite element simulation of the Ni/Ni3C composite under tension. (E) APT map of Ni and C atom distribution. (F) APT map of C atom distribution. (G) In situ tensiletest with strain map. (H) In situ three-point bending test under SEM.

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and the intimate bonding betweenNimatrix andNi3C platelets, the en-hancement of Young’s modulus can be estimated by the rule of mixtures

E ¼ E1n1 þ E2n2 ð1Þ

where E is the modulus of the composite, E1 is the modulus of Ni matrix(210 GPa), n1 is the volume fraction of Ni (86.7%), E2 is the modulus ofNi3C platelets (364 GPa), and n2 is the volume fraction of Ni3C platelets(13.3%). The theoretical Young’s modulus of the composite was cal-culated to be 230.5 GPa, which is close to the average value of Young’smodulus obtained from tensile tests (222 GPa).

Subsequently, we attempted to unveil the strengthening mecha-nisms of the graphene-enabled Ni/Ni3C composite. According to thetensile stress-strain curves in Fig. 3A, the yield strength of Ni/Ni3Ccomposite had a 330-MPa improvement comparing with that of theNi sample produced by powdermetallurgy. Apparently, theNi3C plate-lets contributed to the improvement of yield strength. On the basis ofthe experimental results, a finite elementmodel (FEM) was constructed(Fig. 4D and fig. S7A). When a tensile stress of 600 MPa was exertedalong the x axis, the FEM simulation showed a high concentration ofstress on the platelets, indicating that the platelets acted as the loadbearers and effectively strengthened the composite. The strengtheningeffect from Ni3C platelets can be validated by the rule of mixtures be-cause we can assume that the Ni3C platelets deformed synchronouslywith theNimatrix during elastic deformation.On the basis of the tensiletest results, the Ni/Ni3C composite started yielding at a strain of about0.3%. The improved yield strength was then calculated by the equation

Dssec ¼ eE2n2 ð2Þ

where e is the strain at yield point (0.3%). The improvement of yieldstrength from Ni3C platelets was calculated to be 145.2 MPa.

The hardness derived from nanoindentation on the Ni sampleproduced by powder metallurgy was about 2.2 GPa, which was0.85 GPa lower than that of the Ni matrix in the Ni/Ni3C composite(fig. S7B). Therefore, the Ni matrix should be strengthened by othermechanisms. The second source of strengthening may derive fromthe grain boundaries, i.e., the grain size. The grain boundaries act aspinning points to impede dislocation propagation. The relationshipbetween grain size and strength can be demonstrated by the Hall-Patch equation (39, 40)

Dsy ¼ s0 þkyffiffiffi

dp ð3Þ

where sy is the yield strength, s0 is a materials constant for the startingstress for dislocation movement, ky is the strengthening coefficient, andd is the grain size. Therefore, the change of the yield strength due to thereduction of grain size should be

Dsgb ¼kyffiffiffiffiffi

d1p � ky

ffiffiffiffiffi

d2p ð4Þ

On the basis of TEM inspections (Fig. 2D and fig. S7C), the av-erage grain size along the ND of the Ni/Ni3C composite was 198 nm

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(d1), and that of the Ni sample produced by powdermetallurgy was 543nm (d2). The finer grain size after adding graphene may originate fromthe formation of second-phase particles, which can prohibit recrystal-lization and grain growth. The kywasmeasured to be 4.9MPamm1/2

(40). Thus, the strength enhancement derived from grain boundarieswas 138.3 MPa.

The overall enhancement from the Ni3C platelets and grain bound-aries was 283.5 MPa, at least 66.5 MPa less than the experimentalresults. Because the deformation was the same for all the samples, theremaining mechanisms were precipitate strengthening and solutionatom strengthening. A rational hypothesis is that carbon dissolved intotheNimatrix and then precipitated out as atom clusters and/or second-phase particles, which pinned the migration of dislocations andstrengthened the Ni matrix. A direct evidence of the precipitates shouldbe theweak secondary patterns that appear in selected-area electron dif-fraction (SAED) patterns. However, the SAED pattern of the Ni matrixalong both [011] and [112] directions showed no other patterns exceptfor the Ni (fig. S8), eliminating the possibility of nanosized precipitatesor atom clusters. Thus, solid-solution strengthening had the highestlikelihood for improving the yield strength. Figure 4 (E and F) showsatom probe tomography (APT) maps, showing homogeneously dis-persed carbon atoms in theNimatrix. The atomic percentage of carbonin Ni matrix was about 1 at %. The strengthening effect of interstitialsolution atoms originates from the pinning of dislocation due to latticedistortion. The strength contribution can be expressed as (2)

Dsss ¼ kc ð5Þ

where k is a parameter related to shear modulus and lattice distortionand c is the concentration of interstitial solution atoms. For body-centered cubic metals, such as Fe, carbon atoms can generate a non-symmetrical stress field, which strongly interacts with dislocations,leading to a large k of 5G (G is the shearmodulus).However, the smallerlattice distortion and symmetric stress field stemmed from carbon in-terstitial atoms in face-centered cubic Ni have much weaker pinningeffects on dislocations, making the k only G/10 (2). The shear modulusofNi is 72GPa, and the concentration of carbonwas 1 at%based on theAPT result. The theoretical yield strength enhancement contributed bycarbon interstitial atomswas calculated to be 72MPa. Therefore, theNi/Ni3C composite is triply strengthened by Ni3C platelets, grain bound-aries, and carbon interstitial solution atoms. The theoretical improve-ment of yield strength was calculated as follows

Ds ¼ Dssec þ Dsgb þ Dsss ¼ 355:5 MPa ð6Þ

which is close to the experimentally derived improvement of yieldstrength.

In engineering alloys, the increase of strength derived from intersti-tial atoms and second-phase particles usually trades off with a reductionin ductility. Especially large, brittle carbides often introduce defects,which may become the source of cracks. However, the graphene-derived Ni/Ni3C composite exhibited an obvious plastic deformationstage and a higher static toughness than the pure Ni (Fig. 3, A and B).An in situ tensile test was carried out under SEM to determine the influ-ence of Ni3C platelets on crack formation and propagation during defor-mation. A speckled coating created by amix of conductive silver glue andcarbon black was coated on the sample to trace the evolution of strain via

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digital image correlation. A small notch was also made on the sampleedge. As shown in Fig. 4G, no apparent strain concentration was foundat a low displacement of 0.05mm.With an increase of the exerted load, acrack started at the artificial notch. The crack propagated along 45o

against the loading direction, a typical fracture of ductile metals. There-fore, theNi3Cplatelets likely did not act as the fracture source to stimulatecracking. Close-up observation of the crack initiation near the artificialnotch showed thatmultiple small cracks appeared before themajor crackpropagated (fig. S9), indicating that, in addition to the intrinsictoughening mechanisms, such as fine grain size and parallelly alignedgrain orientation, the brick-and-mortar structuremay introduce extrinsictoughening mechanisms to further improve the toughness. To reveal thepossible extrinsic tougheningmechanisms, an in situ three-point bendingtest was carried out (Fig. 4H). Although an artificial notch was made, thecrack initiated at a large deflection, indicative of outstanding ductility. Af-ter the crack initiated, instead of propagating perpendicular toward an-other side of the three-point bending plate, the crack was graduallydeviated to be parallel to the three-point bending sample length directionby the parallelly aligned Ni3C platelets (fig. S10), which shifted the crackmode, leading to lower effective stress around the crack tip and higherdifficulty for crack opening. Moreover, the interlacing Ni3C platelets re-sulted in a zigzag morphology of crack edges and formation of smallcracks near the primary crack both in-plane and out-of-plane, which in-evitably elongated the crack length. The energy required to propagate thecrack,Ws, is related to the crack length by (29)

Ws ¼ 2abg ð7Þ

where a is the crack length, b is the out-of-plane thickness of the solidmaterial, and g is the sum of surface energy (gs) and energy related toplastic deformation (gp). Apparently, the longer the crack length is, the

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higher the toughness. In addition to the crack deflection, metal bridgesappeared behind the crack tip, and the layer-by-layer structure bluntedthe crack tip,which further prohibited the crack opening andpropagating(fig. S10A). Thus, the Ni3C/Ni brick-and-mortar structure contributed tothe improvement of toughness. Worth mentioning is that the inspectionof the chemically etched sample showed that the Ni3C platelets in factdeformed with the Ni matrix (fig. S10B). This result demonstrated thattheNi3C platelets are truly ductile; they can deformwith thematrix with-out inducing notable cracks. This answered two essential questions: (i)why the Ni3C grain boundary precipitates formed the brick-and-mortarstructure after cold working without introducing large defects and cracksand (ii) why the coarse Ni3C platelets did not induce fracture duringdeformation.

Microstructure and high-temperature hardness ofNi-Ti-Al/Ni3C compositeNi alloys are widely used in high-temperature environments because oftheir outstanding stability and creep resistance. To verify the impact ofthe Ni3C platelets and carbon interstitial solution atoms on the high-temperature performance, 2 wt % Ti and 2 wt % Al were added tothe Ni/graphene powders and sintered together. The obtained Ni-Ti-Al/Ni3C composite also exhibited a brick-and-mortar structure(Fig. 5A) and stripe-like grains (Fig. 5B) after cold rolling. The atomic-resolution EDS analysis of Ni, Ti, and Al (Fig. 5, C to E) showed that TiandAl had concentrated in some grains, forming a laminated structure atthe submicrometer scale. The distribution of carbon was homogeneous(Fig. 5F), which was consistent with the APT map of the Ni/Ni3Ccomposite (Fig. 4F). High-temperature Vickers hardness tests werecarried on pure Ni, Ni/Ni3C composite, Ni-Ti-Al/Ni3C composite, andcommercial HR-224 superalloy. At room temperature, the Ni/Ni3Ccomposite and Ni-Ti-Al/Ni3C composite exhibited a high hardness of

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Fig. 5. Microstructure of Ni-Ti-Al/Ni3C composite and high-temperature Vickers hardness of Ni, graphene-derived Ni/Ni3C composite, Ni-Ti-Al/Ni3Ccomposite, and HR-224 superalloy. (A) SEM image of Ni-Ti-Al/Ni3C composite after chemical etching. (B) High-angle annular dark-field (HAADF) image of the Ni-Ti-Al/Ni3C composite. (C to F) High-resolution EDS of Ni, Ti, Al, and C maps. (G) Hardness values from high-temperature Vickers hardness tests. (H) Room temperatureVickers hardness indentation impression on Ni-Ti-Al/Ni3C composite (the edge length of the inset image is 180 mm). (I) High-temperature (1000°C) Vickers hardnessindentation impression on Ni-Ti-Al/Ni3C composite (the edge length of the inset image is 180 mm).

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3.7 and 4.6GPa, respectively. Alloying theNi/Ni3CwithTi andAl furtherincreased the strength. The Ni/Ni3C composite maintained a high hard-ness from room temperature to 300°C. After 300°C, the hardness de-creased rapidly. The rapid decrease of hardness may be mainly due tothe failure of the pinning effect of interstitial solution carbon atoms onthe dislocations. In comparison, the Ni sample showed a constant de-crease in hardness with increasing temperature; the Ni-Ti-Al/Ni3Ccomposite andHR-224 superalloy showed almost no hardness reductionup to 500°C.When the temperature further increased, the hardness ofNi-Ti-Al/Ni3C composite decreased gradually and remained 1.8 GPaat 1000°C, which is higher than that of HR-224 superalloy, whichdroppeddown to 1.11GPa.Close-up inspections of the centers ofVickersindentations on theNi-Ti-Al/Ni3C composite at room temperature and1000°C are shown in Fig. 5 (H and I, respectively). At room tempera-ture, the indentation surface was relatively smooth (Fig. 5H). High-temperature (1000°C) indentation impression showed an oxidizedsurface with large, irregular particles (Fig. 5I). Therefore, althoughthe alloy recipes and heat treatments require further studies, the pro-posed Ni/Ni3C composite is promising as a basis for next-generationsuperalloys.

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CONCLUSIONSIn summary, a prototypical graphene-derived Ni/Ni3C composite withbioinspired, brick-and-mortar structure was developed. Grapheneclosely wrapped the Ni powders via shear mixing and freeze drying.The Ni/graphene powders were compressed and sintered at 1450°C,forming Ni3C at the grain boundary area. The Ni3C platelets weredeformable. Theywere rolled into long stripes during the cold deforma-tion, leading to the formation of a brick-and-mortar structure. Addi-tional carbon atoms were dissolved into the Ni matrix and existed asinterstitial solution atoms. The Ni3C platelets not only acted as the loadbearers but also redirected crack propagation. The small grains and in-terstitial solution atoms prohibited the dislocation propagation andenhanced theNimatrix. In total, the confluence ofmultiple strengtheningand toughening mechanisms enabled a 73% increase on strength and a6% increase on Young’s modulus, with only 28% reduction in ductility,leading to a 44% improvement in toughness. The Ni-Ti-Al/Ni3Ccomposite exhibited superior strength over commercial superalloys upto 1000°C. This strategy presents a new promise for the design and syn-thesis of advanced bioinspired materials to achieve exceptionally highmechanical robustness for applications in an extensive range of fields.

MATERIALS AND METHODSGraphite and Ni powders were purchased from Sigma-Aldrich Com-pany without further purification. Graphene was prepared using theshear mixing method. Specifically, 1.5 g of graphite powders with pu-rity of 99.9% was added into 200 ml of H2O and then shear-mixed at3000 rpm at room temperature for 1 hour by a Silverson L5M-A shearmixer. Subsequently, the suspension stood for 2 hours, and the largeparticles deposited at the bottom were filtrated for reuse. The upper,transparent liquid contained about 0.09 to 0.12 g of graphene. Nipowders (5 g) with 99% purity were then shear-mixed with the gra-phene sheets for 2 hours. After shear mixing, the Ni/graphene powderswere collected and lyophilized for 6 hours. The obtained powders werecompressed in a round mold with a diameter of 25 mm at 80 MPa andthen sintered at 1450°C for 1 hour with the protection of Ar gas. Aftersintering, the coin was rolled at room temperature for a reduction in

Zhang et al., Sci. Adv. 2019;5 : eaav5577 31 May 2019

thickness of 80%. For comparison, five reference samples were fabricated.For the first reference sample, a pure nickel plate (99% purity) waspurchased from ESPIMetals and cold-rolled with the same thicknessreduction as used for the Ni/Ni3C composite plates. For the secondreference sample, Ni powders were shear-mixed, freeze-dried, com-pressed, and sintered without adding graphene. The sample was cold-rolled with the same thickness reduction as used for the Ni/Ni3Ccomposite plates. For the third and fourth reference samples, Ni powderswere composited with 4 and 6 wt % graphene, respectively, and hybridpowders were then compressed and sintered under the same conditions.The last control sample was a commercial HR-224 superalloy.

Tensile testing was carried out on anAdmet eXpert 2600 tensile uni-versal testing machine with an extension speed of 0.06 mm/min. Strainwas measured using the Vic-2D digital image correlation system fromCorrelated Solutions with one photo every 2 s. Specimens for tensiletests were prepared in a dog bone shape, and the sample size is demon-strated in Fig. 3A. In situ tensile tests and three-point bending tests werecarried out on anMTI Instruments SEMTester 1000 tensile stage.High-temperature hardness testing was performed using a Bruker UMT-3tribometer with a peak indentation load of 30 N for 30-s holding.Nanoindentation tests were carried out using aMicroMaterials Vantagenanoindenter with a load of 200 and 10 mN. The nanoindenter was adiamond Berkovich tip whose shape function was carefully calibrated.The elasticmodulus of the diamond tipwas 1140GPa, and the Poisson’sratio was 0.07. XRDpatterns were obtained using PANalytical X’Pert ProMPD equipped with Cu Ka radiation (l = 0.15406 nm). The micro-structure of the specimens was characterized with a FEI Quanta650 SEMwith anEDSdetector and a FEI TitanG2 aberration-correctedTEM.TEMspecimens for cross-sectional imagingwere cut using aHeliosdual-beam focused ion beam. APT experiments were carried on aLEAP5000XS.

Finite element simulation was performed on ANSYS student edi-tion. The model size was 25 mm by 25 mm by 125 mm (W × H × L).The constituent particle size was determined on the basis of the SEMimages, ranging from 20 to 70 mm in length and approximately 2 to 3 mmin thickness, with a standard width of 5 mm. The volume fraction of theconstituent particles in the model was 10%. Boundary conditions wereassigned as fixed support at one end and distributed face load appliedat the other. Loading was a ramped load uniformly over 30 s to max-imum of 0.375 N (600MPa) along a positive x direction. The unit cell ofNi3C (R�3c space group, a=4.553 Å, and c=12.92 Å) was constructed viadensity functional theory using the Cambridge Serial Total EnergyPackage. The generalized gradient approximation–Perdew-Burke-Ernzerhof was selected to describe the exchange and correlation energy.The Broyden-Fletcher-Goldfarb-Shanno minimizer was used to per-form cell optimization, and the convergence tolerance of total energywas set to be 1 × 10−6 eV/atom.

SUPPLEMENTARY MATERIALSSupplementary material for this article is available at http://advances.sciencemag.org/cgi/content/full/5/5/eaav5577/DC1Fig. S1. Graphene sheets produced by shear mixing.Fig. S2. Microstructure of Ni/graphene particles without shear mixing or without freeze drying.Fig. S3. SEM images of sintered Ni/graphene and Ni sample produced by powder metallurgyafter chemical etching.Fig. S4. Ni with 4 and 6 wt % graphene after sintering.Fig. S5. XRD and EDS spectra of Ni/Ni3C composite and Ni produced by powder metallurgy.Fig. S6. Crystal structure and FFT pattern of Ni3C particle.Fig. S7. Strengthening of Ni/Ni3C composite.Fig. S8. SAED patterns of the Ni matrix.

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Fig. S9. Close-up observation of the crack initiation at the artificial notch.Fig. S10. Close-up inspection of the crack propagation of the three-point bending Ni/Ni3C sample.

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Acknowledgments: We thank the staff members at the University of Virginia NMCF and NorthCarolina State University AIF for electron microscopy technical support. Funding: Financialsupport for this study was provided by the U.S. National Science Foundation (CMMI-1537021).Author contributions: Y.Z. and X.L. designed the experiment. Y.Z. fabricated the composite andcarried out microstructural/mechanical tests. Y.Z. and F.M.H. conducted the in situ tensile andbending tests. F.M.H. performed digital image correlation analysis. J.L.B. constructed finiteelement analysis models. N.S. analyzed crystal structure of Ni3C. D.I. carried out APT tests. Y.Z.and X.L. wrote the article. All the authors proofread the manuscript. Competing interests: X.L.,Y.Z., F.M.H., J.L.B., and N.S. are inventors on a U.S. provisional patent application related tothis worked filed by the University of Virginia (no. 62/817,142, filed 12 March 2019). The authorsdeclare no other competing interests. Data and materials availability: All data needed toevaluate the conclusions in the paper are present in the paper and/or the SupplementaryMaterials. Additional data related to this paper may be requested from the authors.

Submitted 27 September 2018Accepted 17 April 2019Published 31 May 201910.1126/sciadv.aav5577

Citation: Y. Zhang, F. M. Heim, J. L. Bartlett, N. Song, D. Isheim, X. Li, Bioinspired, graphene-enabled Ni composites with high strength and toughness. Sci. Adv. 5, eaav5577 (2019).

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Bioinspired, graphene-enabled Ni composites with high strength and toughnessYunya Zhang, Frederick M. Heim, Jamison L. Bartlett, Ningning Song, Dieter Isheim and Xiaodong Li

DOI: 10.1126/sciadv.aav5577 (5), eaav5577.5Sci Adv 

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