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Carbon 158 (2020) 728e737

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Carbon

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Controllable fabrication of elastomeric and porous graphene filmswith superior foldable behavior and excellent electromagneticinterference shielding performance

Dengguo Lai 1, Xiaoxiao Chen 1, Yin Wang*

CAS Key Laboratory of Urban Pollutant Conversion, Institute of Urban Environment, Chinese Academy of Sciences, Xiamen, 361021, China

a r t i c l e i n f o

Article history:Received 24 September 2019Received in revised form3 November 2019Accepted 14 November 2019Available online 16 November 2019

* Corresponding author.E-mail address: yinwang@iue.ac.cn (Y. Wang).

1 These authors contributed equally to this work.

https://doi.org/10.1016/j.carbon.2019.11.0470008-6223/© 2019 Elsevier Ltd. All rights reserved.

a b s t r a c t

Materials possessing ultrathin, lightweight and foldable characteristics while simultaneously exhibitingsuperb electromagnetic interference shielding effectiveness (EMI SE) are urgently required to guaranteethe operation of portable electronics. Herein, a novel and facile approach is reported to controllablyfabricate porous graphene films (PGFs) with chosen thickness by reduction-induced foaming of grapheneoxide film with a spatial confinement strategy. In striking contrast to unconstrainedly expanded gra-phene films, the confined PGFs surprisingly exhibit excellent elastomeric behavior and strong repeatedfolding stability. Even after high and ultralow temperature treatment, the folding performance of PGFs isnot compromised. Intriguingly, considerably enhanced EMI SE and microwave-absorbing phenomenonare achieved for PGFs compared to compact graphene films due to the efficient wave attenuation in theporous structure and highly conductive network. Significantly, the annealed PGFs exhibit a remarkablyhigh EMI SE of 63.0 dB, enough to block and absorb 99.99995% of the incident radiation, which is farsuperior in performance compared to all graphene-based shielding materials. The largest specific SE/thickness of 49750 dB cm2 g�1 is among the highest values of all the shields known to date. Therefore,such thin, lightweight, elastic PGFs with outstanding EMI-shielding performance are highly promisingfor applications in foldable and wearable electronics.

© 2019 Elsevier Ltd. All rights reserved.

1. Introduction

Currently, high-performance electromagnetic interference(EMI) shielding materials are urgently needed to avoid theincreasingly serious radiation pollution created by the explosiveuse of electronic devices, which has detrimental impacts on humanhealth and sensitive electronic equipment and systems [1]. Inparticular, the explosive growth of portable devices and smartwearable electronics has led to more stringent requirements forlighter, thinner, more flexible and even foldable EMI shieldingmaterials with higher shielding efficiency [2,3]. However, most ofthe present materials cannot simultaneously integrate theseintriguing characteristics [4]. Macroscopic layered graphene films(GFs) tightly and orderly assembled by two-dimensional nano-meter-scale graphene sheets are highly promising for excellent EMI

shielding performance by taking full advantage of graphene’s su-perior electrical property, and they are considered a candidate toreplace heavy and easily-corroded metal-based shielding materials[5e7]. As pristine graphene can only be dispersed in a few choicesolvents at extremely low concentrations (<1 mg ml�1) [8], gra-phene oxide (GO) has been used as a reliable chemical function-alized graphene with wide utilization for constructing GFs due toits good solubility in water. However, the electrical conductivity ofGO is seriously reduced because of its disrupted sp2 structures. Toimprove the electrical conductivity of constructed GO film and inturn achieve higher EMI shielding effectiveness (SE) of the resultingGFs, ultrahigh-temperature annealing (over 2000 �C) for graphiti-zation is commonly used to repair the crystal defects of GO nano-sheets (e.g., an 8.4 mm-thick graphene film obtained 20 dB after2000 �C annealing [5], a 13 mm-thick film obtained 48.9 dB after2800 �C annealing [9], and 8.8 mm-thick GFs achieved 37 dB after3000 �C annealing [10]). Nevertheless, such harsh conditions notonly tremendously increase the processing cost, time and energyconsumption and make mass production infeasible but also inevi-tably reduce the mechanical properties of GFs [11e13].

D. Lai et al. / Carbon 158 (2020) 728e737 729

Elastic GFs, especially those that allow reversible large-straindeformation, are highly desirable for applications in foldable andwearable devices [14]. Note that the introduction of structural hi-erarchy and porous architecture can significantly improve materialutilization with further enlarged elastic strain range and ductility[15]. Coincidentally, the foaming strategy is undoubtedly an effec-tive way to decrease the weight or density of materials, and theresultant foam structure in GFs is beneficial to improve its shieldingefficiency and microwave absorption property [3,16,17]. Therefore,the foaming of compact layered GFs into porous graphene films(PGFs) is an efficient strategy to combine excellent elastic andlightweight properties with remarkable EMI shielding into onematerial. Recently, graphene aerogel film was developed by3000 �C self-expansion of reduced graphene film and exhibited ahigh EMI SE of over 60 dB [4]. However, the pores generated instructure were in a nonuniform distribution and uncontrollable,and the flexibility of these films was not satisfactory for the diversepurposes of foldable and wearable devices. Furthermore, suchextremely high temperatures are inevitably accompanied by flexi-bility deterioration, and this unpleasant ultrahigh temperaturetreatment is extremely energy-consuming, which limits the in-dustrial use of the films. To the best of our knowledge, few papershave focused on the fabrication of elastic PGFs and have consideredtheir EMI shielding performance. Thus far, although many greatimprovements in graphene-based shielding materials have beenachieved, the daunting challenge of the facile and controllablefabrication of fully flexible or foldable, ultrathin, lightweight, andporous graphene films with outstanding EMI shielding perfor-mance remains unresolved [18].

Herein, we demonstrate the controllable fabrication of thin,lightweight, elastic and porous graphene films by assembling GOnanosheets into films followed by a simple leavening-accompaniedreduction with a spatial confinement strategy. We used a confinedspace to constrain the expanding films to a desired and uniformthickness. In contrast to conventional unconstrained foamed GFs(brittle), our uniformly expanded PGFs with engineered porosityexhibit excellent elastomeric behavior and strong 1000-cyclerepeated folding stability. Such PGFs present unchanged flexi-bility even after high-temperature annealing (1000 �C) and ultra-low liquid nitrogen temperature (�196 �C) treatment. Moreimportantly, a greatly enhanced EMI-shielding performance isachieved for the porous GFs compared to their unfoamed compact

Fig. 1. Confined foaming strategy and resultant porous graphene films. (a) Schematic rep(b) Digital photos of porous graphene films with chosen thicknesses of 60, 100, 200, 400 andpicture represents its thickness. (c) A photograph of expanded paper-like porous film with afilm is further folded into a boat. (A colour version of this figure can be viewed online.)

film counterpart reduced by HI because of the favorable porousnetwork. Also, the EMI shielding behavior and its enhancementmechanism of the foam structure are elucidated.

2. Results and discussion

The methodology and mechanism of controllably fabricatingPGFs is schematically illustrated in Fig. 1a. The nacre-like GOlayered film is prepared by evaporating an aqueous suspension ofGO at room temperature, and then it is confined between twoparallel glass plates with a gap formed by placing spacers of definedthickness between the plates. A small amount of hydrazine as thereducing agent is precoated on the surface of the bottom glassplate. The two plates are then clamped together to prevent furtherexpansion of the film beyond the desired thickness during thereduction-accompanied foaming process. When heating to 90 �C,the hydrazine vapor initiates the chemical reduction of GO andyields gaseous species, such as CO2, CO and H2O, which cangenerate enough inner pressure to overcome the van der Waalsforces holding the GO layers together; accordingly, the film israpidly expanded to the target thickness with effective expansion-exfoliation of the GO layers [14,18,19]. This confined foaming of GOfilm is achieved by two competing effects: the pressure generatedby the rapid evolution of gas products, which has been suggested asthe main driving force for the graphene laminate exfoliation [19];and the confinement of the top and bottom plates, which applymechanical pressure to control the release of gas products andprevent excessive expansion from destroying the porous structure.By changing the gap above the film, the final thickness of theexpanded porous film could be strictly controlled to chosen valuessuch as 60, 100, 200, 400 and 1000 mm, which are tens or hundredsof times the original GO film thickness (~10 mm). The digital imagesof the corresponding PGFs are shown in Fig. 1b (surface) and Fig. S1(cross-sectional view) in the Supporting Information, presenting asmooth and flat surface without any bumps or cracks and uniformthickness. In addition, this confined foaming process could be easilyamplified to a large scale, and enlarged paper-like PGFs(100 mm � 100 mm) with a thickness of 100 mm are prepared(Fig. 1c). Such porous graphene films can be further folded intoboats and floats on flowers, exhibiting excellent flexibility andfoldable properties as well as light weight.

Scanning electronmicroscopy (SEM) observation shows that the

resentation illustrating the method and mechanism for confined foaming of GO films.1000 mm from the confined expansion of original 10 mm-thick GO film, the number indimension of 100 mm � 100 mm and target thickness of 100 mm, the flexible porous

D. Lai et al. / Carbon 158 (2020) 728e737730

brown GO film exhibits a compact layer-by-layer nanostructurewith a thickness of ~10 mm, as shown in Fig. S2a (Supporting In-formation). In the case of unconstrained foaming of GO film(without the confinement of plates), the surface of the final gra-phene film has many bumps and voids (Fig. S2b, Supporting In-formation). The cross-sectional SEM views indicate that suchbumps and large voids are a structure of over-expanded and over-exfoliated layers, and the linking of the graphene laminates appearsto be quite weak and random. In such an unconstrained foamingprocess, the originally generated microvoids would readily evolveto large voids and bumps due to the accumulation and randomrelease of gaseous products, thus inevitably damaging theconnection between graphene layers and the mechanical proper-ties of films. The unconfined foaming of a 10 mm-thick GO filmcauses the film to expand approximately 100-fold in thickness butwith an uneven pore structure and an extremely rugged surface.The variables of hydrazine amount and distribution, its contactform with the GO film, and processing temperature immenselyinfluence the final film thickness using the unconstrained foamingmethod, making the fabrication of uniform PGFs uncontrollableand difficult. By contrast, the rapid expansion of spatially confinedGO films is restricted because the top and bottom surfaces of theexpanding film cannot extend past the confining glass plates, andthe film is thus limited in the height reached. Accompanying thisconfinement, once the gas products are generated, they escape inthe horizontal direction and release rapidly without giving theevolved gas any chance to over-expand the film or destroy theuniform porous network.

The cross-sectional SEM views of confined foaming films, asdepicted in Fig. 2a-e, clearly reveal the formation of a cross-linked

Fig. 2. Morphological characterization of porous graphene films. (aee) Cross-sectional SE200, 400 and 1000 mm, respectively. (f) Surface SEM views of 200 mm-thick PGFs. (gei) Theviewed online.)

microcellular structure with cell sizes ranging from several to tensof microns in an oriented structural hierarchy. The porewalls in thisstructure are continuously interlinked, with intermittent separa-tion between different layers. The thickness of confined PGFs,determined by the thickness of the spacer used, is strictly 60, 100,200, 400 and 1000 mm, with parallel top and bottom surfaces. Inother words, the size of the engineered microvoids in the micro-cellular structure could be precisely controlled using the confinedfoaming strategy. The graphene layers within these PGFs are highlyordered in a roughly parallel fashion and are well-aligned, whichmay be due to the way that the evolved gas escapes in the hori-zontal direction. All the confined PGFs have a noticeable uniformpore structure across each of their cross sections relative to theunconstrained film with a disordered foam structure containingmany large voids. The layers become more undulating as theexpansion ratio increases from 6 to 100, and the average spacebetween the layers increases. Even at a low expansion ratio, theseparation of graphene layers happens, and microvoids are created.When the expansion ratio reaches 10, the uniformly porousstructure is totally formed, as shown in the inset of Fig. 2b. Sur-prisingly, evenwith a high expansion ratio of 100, the connection ofthe graphene laminates is still sinewy and robust (Fig. 2h) and is notdestroyed by the evolved gas. The graphene laminates overlap atthe edges and construct a continuous cross-linked microcellularstructure, forming a multilayer architecture with alternating gra-phene laminates and space layers. Such overlapped laminatescontain several loosely stacked graphene monolayers with manywrinkles on their surface (Fig. 2g). It is remarkable that the formedpores are relatively closed in a configuration, in which the middlelaminate bulges and the surrounding edge overlaps compactly with

M images of controllably expanded porous graphene films with thicknesses of 60, 100,magnified linking region of (def), respectively. (A colour version of this figure can be

D. Lai et al. / Carbon 158 (2020) 728e737 731

adjacent pore edges (Fig. 2h); this characteristic is totally differentfrom the completely open pores in graphene foams [20,21] oraerogels [22], carbon nanotube sponges [23] and/or their polymercomposites [24,25]. As an example, the content of the relativelyclosed cell in the 200 mm-thick PGF is measured to be ~84.2% usingthe pycnometer method, demonstrating that the pores in PGFs aremainly in the form of relatively closed cell structure. Such relativelyclosed pores and microcellular structures not only provide abun-dant interfaces for microwave internal reflections but also help toentrap the incident radiation until it is absorbed. Moreover, thesurface of these PGFs is quite flat and airtight due to the compactstacking of graphene sheets, with some wrinkles on their surface,as shown in Fig. 2fei. Consequently, the porous graphene filmswithengineered homogeneous porosity possess a configuration of“compact in plane and porous in cross-section”, and this perfectstructure would dramatically improve the EMI shieldingperformance.

The numerous uniform microvoids introduced into the gra-phene film are expected to greatly improve its flexibility. This isverified by the single and double folding and even crumpling ofthese PGFs. With the release of the force, the folded and crumpledfilm can rapidly spring back to its initial flat shape akin to anelastomer, without any damage or plastic deformation (Fig. 3a andMovie S1, Supporting Information), which has never been reportedfor paper-like films assembled using other 2D building blocks. Mostflexible GO and rGO films reported so far can be bent to 180� andfolded, but they exhibit plastic deformation, and creases remainafter folding, with no elastic or self-recovering property observed(Fig. S3, Supporting Information) [11,13]. Even for films composedof microfolded layers of graphene, they could not recover to theirinitial shape after folding because of irreversible plastic deforma-tion [26]. These results indicate an unprecedented elastomericfolding performance of our porous graphene films.

Supplementary video related to this article can be found athttps://doi.org/10.1016/j.carbon.2019.11.047.

Fig. 3. Mechanical flexibility of graphene films. (a) Video snapshots of single, double foldinfilm reveals the performance of the elastomer. Crumpling and releasing of (b) thermal annea(1000 mm-thick). Optical images of folding (d) unconfined foaming graphene film and (e) rGversion of this figure can be viewed online.)

We annealed a piece of the as-expanded 200 mm-thick PGF at1000 �C for 2 h in nitrogen, and the resulting annealed porousgraphene film (APGF) is almost unchanged in its geometry,although it is further deoxygenated, indicating that it retains a“memory” of its earlier constraint. Surprisingly, when this ther-mally annealed PGF is folded and crumpled, it can still revert backto its original shape (Fig. 3b and Movie S2), illustrating anoutstanding flexibility and impressive elastic property of PGFs evenafter high-temperature annealing. The interconnections of gra-phene layers do not show any change after annealing, as observedin the SEM images shown in Fig. S4 (Supporting Information),establishing the cornerstone for the excellent flexibility of the PGFs.The content of relatively closed cell in the 200 mm-thick APGF ismeasured to be ~84.5%, similar with the value of PGF beforeannealing, indicating that the thermal annealing process has littleeffect on the pore structure of porous films. Furthermore, theexcellent fire resistance of the PGF is confirmed by burning it overan alcohol lamp (up to ~500 �C), and the result demonstrates thatsuch PGFs can tolerate the flame for more than an hour (Fig. S5 andMovie S3). A thicker 1000 mmporous film is folded and crumpled inliquid nitrogen, and the ultralow temperature (�196 �C) treatedand crumpled PGFs could also recover completely at a noticeablyvisible rate without leaving creases (Fig. 3c and Movie S4). Theseresults indicate that the extremely high flexibility of our porousgraphene films is not affected by thermal annealing and ultralowtemperature treatment; such a wide temperature range and largeoptional expansion ratio scope (up to a hundred-fold) greatlyenhance its application. For comparison, the flexibility of the un-constrained foaming film and rGO film thermally reduced at1000 �C under the same conditions as those used for PGF annealing,inwhich the same 10 mm-thick GO film as a precursor is used, is alsoevaluated. Cracks and breaks are observed for the unconstrainedfilm after folding due to the weak layer connections and disorderedpore structure with many large voids (Fig. 3d). It is noted that if thepores in the structure are very large, they will readily evolve to

g and even crumpling of the confined porous graphene film (200 mm-thick). The porousled porous film (200 mm-thick) and (c) liquid nitrogen temperature treated porous filmO film thermally reduced at 1000 �C for 2 h as those used for PGF annealing. (A colour

D. Lai et al. / Carbon 158 (2020) 728e737732

large cracks during folding. The thermally reduced graphene filmeven directly breaks apart and cracks after bending (Fig. 3e), andthis brittleness is attributed to its substantially condensed structure(the thickness decreases from 10 mm to 5 mm after thermalreduction).

Supplementary video related to this article can be found athttps://doi.org/10.1016/j.carbon.2019.11.047.

The uniform microvoids and the separated parallel graphenelayers are responsible for the extreme flexibility of PGFs. Whenfolded, a film experiences tensile stress at the outer surface andsimultaneous compressive stress at the inner surface, and the stressreaches highest at the folding site. It has been detected that thecritical radius of curvature to break the graphene film (at whichfacture occurs) was proportional to thickness [27]. In this work, thethickness of the PGFs is tens to hundreds of micrometers, but thePGFs is comprised of very thin and highly separated graphenelayers, which are less than 30 nm in thickness as shown in Fig. 2g.Owing to the large and uniform microvoids in the PGFs, theseseparated graphene layers can flex and bend independent of eachother. The critical radius of curvature required to break the ultra-thin and highly separated graphene layers is expected to be verysmall (several nanometers), this holds the key to the exceptionalflexibility of PGFs. In other words, the curvature radius (severalmicrometers) of the separated graphene layers under folding ismuch higher than the critical curvature radius, thus these

Fig. 4. Structural characterization of porous graphene films. (a) XRD patterns, (b) Ramanfilm with thickness of 200 mm and its annealed porous film. (A colour version of this figur

separated laminates can withstand fracture on the tensile side. Onthe compression side, the compress stress can be accommodatedby the large microvoids that are formed between the separatedlayers. Therefore, the PGFs exhibit the excellent elastomeric foldingbehavior rather than plastic deformation. Furthermore, it shouldemphasized that the formed uniform pores by confined foamingand the well interconnected graphene network are also crucialfactors determining the flexibility of PGFs. Such unprecedentedflexibility and elastic properties render the porous graphene film anideal candidate shielding material for flexible and foldable smartelectronics.

There is an average weight loss of 42%e44% after hydrazine-induced reduction for PGFs, while a similar weight loss of 43% isobserved for the compact rGO film reduced by HI (film-HI), indi-cating that the deoxygenation and reduction degree by hydrazineand HI have no significant difference. Meanwhile, the thickness offilm-HI decreases from ~10 mm to ~8 mm, which should beattributed to the interlaminar consolidation of graphene layersduring the reduction. SEM cross-sectional observation (Fig. S6,Supporting Information) together with XRD analysis (Fig. 4a) revealgood stacking of graphene layers with an interlayer spacing of0.36 nm. A summary of the interlayer spacing and the full width athalf-maximum (FWHM) extracted from the XRD patterns of all thecharacterized films is given in Table S1 (Supporting Information).The XRD pattern of the original GO film features a typical sharp and

spectra, and (c) C1s XPS spectra for the original GO film, film-HI, and porous graphenee can be viewed online.)

D. Lai et al. / Carbon 158 (2020) 728e737 733

strong peak at 2q of 10.87�, corresponding to a d-spacing of 0.81 nmand FWHM of 0.61�. The intensity of the main peaks for PGFs ismuch lower than that for the GO film, and the peaks shift to higherangles that correspond to d-spacing values in the range of0.35e0.43 nm. The decrease in the d-spacing is due to the elimi-nation of interlamellar water and most oxygen-containing groups.The XRD patterns for PGFs with different foaming thicknessesappear to be similar, but their intensities gradually decrease withincreasing expansion ratio (see details in Fig. S7 and Table S1),revealing a gradually higher exfoliation level. The broadening of theFWHM suggests a decrease in the mean size of the ordered stack ofrGO platelets [28]. The presence of the main peaks at the higherangles of 24.73�e25.45� indicates that the expanded films containnumerous domains of densely stacked/overlapping layers of rGOplatelets on the surface of the PGFs, which is consistent with theSEM results discussed above and shown in Fig. 2i. The presence ofpeaks around 20� is originated from the residual oxygen-containinggroups and induced multimodal character of the interlayer spacingof PGFs [28]. After thermal annealing, the peak at 20� disappearsand the main peak shifts to a higher angle of 26.86�, correspondingto a decreased d-spacing of 0.33 nm and FWHM of 1.74�, whichdemonstrates the further removal of oxygen-containing groups andrepairing of the defects of rGO nanosheets.

The structural integrity of the samples is evaluated by Ramanspectroscopy, as shown in Fig. 4b, and two peaks at 1350 cm�1 and1590 cm�1 correspond to the well-documented D and G bands,respectively. The ID/IG ratio determined by peak intensity is an in-dicator of the degree of disorder and average size of the sp2 carbondomains of the graphene platelets. It is evident that the intensityratio of ID/IG increases from 0.91 for the GO film to 1.22 for film-HIand 1.21 for PGFs, indicating the decrease in the average size of thesp2 carbon domains and successful reduction of the GO sheets,which is ascribed to the removal of oxygen-containing groups andthe generation of topological defects and vacancies in the graphenelattice [14]. The quite similar ID/IG values for film-HI and PGFsstrongly suggest that they possess nearly the same structuralintegrity. After thermal annealing, ‘‘self-healing’’ of graphitic lat-tices takes place, and the ID/IG value decreases to 1.08, which isattributed to the restoration of the original sp2 lattice by newlyformed smaller sp2 clusters [29].

X-ray photoelectron spectroscopy (XPS) is used to analyze thesurface chemical composition and bonding states of the GO andGFs. The C 1s XPS spectra for the original GO film, film-HI and thePGFs are shown in Fig. 4c, with the full surveys given in Fig. S8. TheC 1s spectrum of the GO film can be deconvoluted into three fittedpeaks at binding energies of 284.8 eV for C]C groups, 286.8 eV forthe CeO (hydroxyl and epoxide) groups and 287.9 eV for the C]O(carbonyl) groups, while for the film-HI and PGFs, two additionalpeaks at binding energies of 285.8 eV for CeC bonds and 289.2 eVfor OeC]O groups are fitted. Obviously, the intensity of the oxygengroup peaks in the film-HI, PGF and APGF decreased noticeablyrelative to that in GO film because of the successful reduction. TheC/O ratio for the GO film is 2.74, and the corresponding values forfilm-HI and PGF are 8.26 and 6.31, respectively, further showing thedeoxidization of GO film. The XPS spectrum of the APGF has anincreased C/O ratio of 13.88, strongly confirming the restoration ofthe aromatic carbon atom network.

Materials with large electrical conductivity are typically neededto obtain high EMI SE values. Fig. 5 presents the electrical con-ductivity (EC) of PGFs. Owing to the presence of a microcellularstructure that would impair the conductive network, PGFs formedby confined foaming of a 10 mm-thick GO film show a lower EC thanthat of compact film-HI (~14840 S/m). The EC values decrease from1931 S/m to 49 S/m as the expansion ratio increases from 6 to 200(Fig. 5a). Note that the conductivity initially decreases rapidly at a

low expansion ratio and then tends decrease slowly when theexpansion ratio is higher than 30. To elucidate the effect of filmexpansion on the conductivity, confined foaming of GO films withdifferent thicknesses of 5, 10, 15 and 20 mm is conducted, and theresults are depicted in Fig. 5b. As the original GO film thicknessincreases from 5 mm to 20 mm, the EC remarkably increases from668 S/m to 2102 S/m at the expanded thickness of 100 mm, a morethan threefold improvement. Additionally, a remarkable improve-ment of ~300%e400% is observed for all PGFs when considered atthe same final thickness of 200, 400 and 1000 mm. By foaming tothe same thickness, the exfoliation degree of graphene laminatesdecreases with increasing the initial GO film thickness; inversely,the laminate connection area increases, resulting in enhancedconductivity. In addition, all the EC values follow a similardecreasing trend with increasing foaming thickness (expansionratio) for the same GO film precursor, which is in agreement withthe results in Fig. 5a. To further improve the EC of materials, a1000 �C annealing treatment is performed. As plotted in Fig. 5b, adramatic 3~5-fold increase in conductivity is achieved after ther-mal annealing due to the further restoration of defects in rGOsheets. Such an annealing process does not affect the flexibility offilms (Fig. 3b) but greatly increases the conductivity and reducesthe weight of the PGFs, making them more promising as light-weight EMI shielding materials. More importantly, the highestspecific electrical conductivity (EC divided by density) of the APGFsreaches 840 S cm2 g�1, which is far superior to the correspondingcompact film (72 S cm2 g�1) and most graphene-polymer com-posites (0.42 S cm2 g�1) [30], graphene foams (52 S cm2 g�1) [18],pristine graphene aerogels (450 S cm2 g�1) [31] and CVD-graphenefoams (33 S cm2 g�1) [2] reported in the literature.

To further explore the foaming performance, confined expand-ing of GO films with different original thicknesses at the sameexpansion ratio is also conducted (Fig. 5c). Surprisingly, theexpanded porous films with the same expansion ratio from variousGO film thicknesses possess similar EC values, indicating that thefoam structure and graphene laminate interconnections are quitesimilar at the same expansion degree. The reduction-inducedconfined foaming process does not depend on the GO thickness,and the construction of the pore configuration is stable, uniformand identical. Furthermore, the foldability of PGFs and APGFs isevaluated through manual folding and unfolding for 1000 cycles(Fig. 5d and e). We do not observe crack formation or breaks in thefilm after repeating the folding process. Even after such abuse, theability of the PGFs and APGFs to rapidly spring back to their originalshape is retained. The electrical conductivity shows a negligiblechange even after folding and unfolding 1000 times for both of thePGFs and APGFs, which confirms that the structural integrity of theporous film is retained despite repeated folding. The mechanicalproperties of porous graphene films are confirmed by tensile andcompression tests. As illustrated in Fig. 5f, the tensile strength of200 mm-thick PGF and APGF from confined foaming its precursor10 mm-thick GO film is 7.12 MPa and 6.92 MPa, respectively. Suchquite similar tensile strength indicates that the thermal annealingprocess has little effect on the mechanical property of PGFs. Itshould be noted that, those tensile strength values are much higherthan previously reported porous graphene films or graphene foams(<5 MPa) [14,18,32], strongly demonstrating the strong intercon-nection and effective load transfer between assembled graphenesheets. Moreover, the compressive stress-strain responses of PGFsand APGFs are conducted to further evaluate their foldability. Theresults in Fig. S9 (Supporting Information) illustrate that the com-pressing stress-strain curves of 200 mm-thick PGF and APGF during50 folding cycles are highly overlapping, indicating the excellentcompressibility and recoverability of PGFs and APGFs. Thecompressive stress (resilience) of PGF and APGF during folding and

Fig. 5. Electrical conductivity of porous graphene films. (a) Variation in electrical conductivity of expanded porous film from 10 mm-thick GO film in terms of expansion ratio. (b)Effect of initial GO film thickness on the electrical conductivity of expanded porous films and their further annealed films. (c) Comparison of the conductivity of porous graphenefilms with different thicknesses as a function of expansion ratio. Electrical conductivity changes of the 200 mm-thick (d) PGF and (e) APGF versus the number of folding cycles; theinset optical images show the folded (left) and released (right) PGF and APGF. (f) Tensile stress-strain testing of 200 mm-thick PGF and APGF. (A colour version of this figure can beviewed online.)

D. Lai et al. / Carbon 158 (2020) 728e737734

compressing reaches 30.0 kPa and 27.5 kPa, respectively, directlyproving the outstanding elastic property of our porous graphenefilms.

To explore the EMI shielding property of PGFs and to emphasizethe impact of porous structure on EMI SE, the direct comparisonbetween compact graphene film (film-HI) and the correspondingmicrocellular PGFs with various thicknesses originating from thesame 10 mm-thick GO film is investigated, as shown in Fig. 6a.Despite the lower electrical conductivity of the as-prepared PGFs,they exhibit an observably improved average SE value of higherthan 27.2 dB over the whole X band, which is 90% larger than14.3 dB measured for film-HI. Although the increased expansionratio brings a dramatic decrease in film conductivity, it still resultsin a considerable improvement in both overall shielding effective-ness (SEtotal) and absorption (SEA), even with a slight decrease in

reflection (SER). The highest EMI SE reaches 34.4 dB for the1000 mm-thick PGF at the expansion ratio of 100, improving by140% comparedwith film-HI. Such remarkable improvement of EMISE is primarily attributed to the absorption, and the considerablylarger absorption and lower reflection suggests an absorption-dominant EMI shielding for PGFs. The implication of these resultsis that transforming the compact layered graphene structure into aporous structure can lead to a remarkable improvement in EMIshielding, which is ascribed to the formation of enhanced internalmultiple reflections at the larger cell-matrix interface area owing tothe microcellular structure in PGFs. This foam construction coupledwith an interconnected conductive network could harness thepotential of graphene as a superior shielding material to a greaterextent, that is, obtaining a higher EMI SE with lower grapheneamount, greatly improving the graphene effectiveness and

Fig. 6. EMI shielding performance of porous graphene films. (a) EMI SE of compact film-HI and expanded porous graphene films (PGFs) with different expansion ratios. (b) EMISE of 100 mm- and 200 mm-thick PGFs expanded from GO films at thicknesses of 5, 10 and 20 mm. (c) Shielding performance of annealed PGFs (APGFs) with thicknesses of 100 and200 mm by confined foaming of 10 mm and 20 mm-thick GO films. (d) Comparison of the total EMI SE (SETotal) and the absorption (SEA) and reflection (SER) mechanism of PGFs andAPGFs at thicknesses of 100 mm and 200 mm by expanding a 20 mm-thick GO film. (A colour version of this figure can be viewed online.)

D. Lai et al. / Carbon 158 (2020) 728e737 735

reducing the cost. Furthermore, with the expansion and exfoliationof graphene sheets, the larger interface area and integrated porousconfiguration facilitate a higher absorption attenuation than thebulk structure of film-HI. The repeated reflection and scattering ofmicrowaves within porous films significantly prolong the trans-mission pathways, thus naturally enhancing the microwave ab-sorption and attenuation combined with the highly conductivenetwork. It should be noted that this dramatic enhancement isattributed to both the microcellular structure and the 3D conduc-tive network, and high EMI SE cannot be achieved without eitherone of them.

Normally, EMI SE is directly proportional to electrical conduc-tivity. By confined foaming to the same thickness of 100 mm and200 mm, the PGFs derived from the 20 mm-thick GO film exhibit amuch higher EMI SE of 33.1 dB and 38.6 dB, respectively, comparedto the SE values of PGFs from the 5 mm-thick GO film (22.4 dB and23.6 dB) and 10 mm-thick GO film (27.2 dB and 28.9 dB), increasingby 47.6% and 63.5% for the 100 mm- and 200 mm-thick PGFs as theinitial GO thickness increases from 5 mm to 20 mm (Fig. 6b). Thisenhancement is mainly attributed to the higher conductivity andmore graphene layers for microwave shielding in the PGFs foamedfrom a thicker GO film. At the same expansion ratio of 10, despitethe quite similar conductivity, the EMI SE of 200 mm-thick PGF from20 mmGO film (1050 S/m, 38.6 dB) shows an obvious improvementin contrast to that of 100 mm-thick PGF from 10 mm GO film (998 S/m, 27.2 dB), indicating that the thicker PGFs possess higher EMI SE.To strip out other variables, except the conductivity, that could

influence the final EMI shielding performance, the method ofthermal annealing is used to increase the conductivity whileretaining the pore structure of the PGFs. Remarkably, the SE valuesof 100 mm- and 200 mm-thick PGFs (PGF-100 and PGF-200,respectively) from the 10 mm GO film improved to 44.3 dB and47.8 dB after annealing (Fig. 6c). For PGFs expanded from 20 mm-thick GO film, the annealed PGF-100 (APGF-100) and annealed PGF-200 (APGF-200) exhibit an extremely high SE value of 56.1 and63.0 dB, respectively, enough to block and absorb 99.9998% and99.99995% of incident radiation with only 0.0002% and 0.00005%transmission. Such superior EMI shielding performance of porousgraphene films with small thicknesses and elastic properties hasnever been reported before. The comparison of SEA and SER for20 mm-thick GO film foamed PGFs and APGFs in Fig. 6d confirmsthat the improvement of shielding performance by annealing is alsomainly attributed to the SE absorption, while the SE reflectioncontributes little to the increase. It is thus clear that the porousstructure coupled with enhanced intrinsic conductivity is respon-sible for the outstanding EMI shielding performance of APGFs.

The superiority of lightweight and elastic PGFs over other ma-terials including metal- and carbon-based composites is high-lighted by comparing their specific shielding effectiveness (SSE) (SEdivided by density). As shown in Fig. 7a and Table S2 in the Sup-porting Information, our porous graphene films clearly outperformall of the known synthetic materials and rank at the top of thecomparison chart when compared with other shielding materials,including carbon-based solid materials (lower right), metal-based

Fig. 7. Comparison of EMI SE with the previous literature and shielding mechanisms. (a) Comparison of the specific SE as a function of thickness. A detailed description of eachdata point is presented in Table S2 in the Supporting Information. (b) Proposed EMI shielding mechanism. (A colour version of this figure can be viewed online.)

D. Lai et al. / Carbon 158 (2020) 728e737736

solids including MXene films (lower left), carbon-based foams(upper right) and metal-based foams (left). By considering thematerial thickness, a more realistic parameter of SSE/t (SSE dividedby thickness) is introduced to determine the effectiveness of thismaterial. Thus far, although there have been encouraging advances,few materials that integrate lightweight, ultrathin, highly flexibleand even elastic properties and ease of processing with highshielding performance have been reported. The comprehensiveliterature review of previously studied materials of SSE/t presentedin Table S2 clearly indicates that porous graphene films are amongthe best EMI shielding materials known to date. The highest SSE/tvalue is 49750 dB cm2 g�1, which is 500% higher than that of itscorresponding solid graphene film (8635 dB cm2 g�1) and ranksfirst of all the carbon-based EMI shielding materials (Fig. S10,Supporting Information), outperforming most of the shieldingmaterials studied thus far (Table S2). Normally, to satisfy thecommon commercial EMI shielding requirements (above 20 dB)[30], graphene-polymer composites with large thicknesses, usuallymore than 2 mm, are needed, which prevents the practical use ofthese materials in applications in aerospace and communicationdevices. The PGFs are far superior in performance compared tographene-based materials in EMI shielding applications. For com-parison, the thermally reduced graphene oxide film that possessedhigher electrical conductivity is also plotted and is far below ourporous films (Table S2) [7,17,33,34]. Very recently, superb shieldingperformance over 50 dB was achieved for MXene (2.5 mm) [1] andMXene foams (18 mm) [35] compared with all other syntheticmaterials. However, their potential applications may be hinderedby poor chemical instability, inferior high-temperature resistance,heavy weight and unsatisfactory flexibility. By contrast, the SSE/tvalue of our porous graphene films is comparable or even higherthan that of MXene-based materials. Moreover, these PGFs alsopresent excellent corrosion and temperature resistance, low den-sity, and outstanding folding and elastic properties, which shouldfind wider potential applications in the fields of aerospace, defense,and smart foldable electronics.

The massive EMI SE of PGFs can be understood from the pro-posed mechanisms shown in Fig. 7b. As the incident electromag-netic (EM) waves strike the surface of highly conductive PGFs, somewaves are immediately reflected due to the large impendencemismatch between air and graphene layers. In particular, thelaminated structure of densely stacked graphene platelets on thePGF surface is more favorable for reflecting radiation. Theremaining waves that pass through the surface graphene laminatesinduce a microcurrent by charge carriers that contribute to ohmiclosses, resulting in a drop in energy of the EM waves. After passingthrough the surface layer of PGFs, the surviving waves encounter

the next barrier layers with a cell-like configuration, and the phe-nomenon of EM wave attenuation repeats. Simultaneously, theinternal layer acts as a reflecting surface and gives rise to multipleinternal reflections, and the repeated reflection and scatteringgreatly enhance the transfer of EM energy, which is dissipated asheat in the form of microcurrent, leading to the enhancement in SEabsorption. Remarkably, the relatively closed cellular pores in themicrocellular structure constructed by exfoliated, dense and highlyconductive graphene flakes may entrap the incoming EM waves.The entrapped EM waves continue to bounce off the graphene cellwalls until they are completely absorbed and quenched within thisstructure (see inset of Fig. 7b). This is in marked contrast to film-HIwith a condensed structure and no interlayer reflecting surfaceavailable to provide the internal multiple reflection phenomenonand entrapping of waves. Thus, the multilayered laminated gra-phene structure and internal pore configuration provide PGFs withthe ability to behave as a multilevel shield and “trap” to attenuateand dissipate waves in the form of heat within the materials.Furthermore, local dipoles of the residual oxygen-containinggroups may be created when subjected to an alternating electro-magnetic field. As these dipoles interact with the incoming EMwaves, such induced dipole polarization leads to polarization los-ses, which in turn improve the overall shielding. The propertycombining high electrical conductivity (charge carriers and dipoles)with a porous structure is responsible for the extremely high EMIshielding efficiency of PGFs.

3. Conclusions

Herein, a graphene filmwith engineered porosity is prepared bya simple leavening-accompanied reduction treatment of its pre-cursor GO films. This treatment is found to introduce an abundanceof homogeneous microvoids between the graphene layers. In ourmethod, by confining the GO film between two glass plates, the filmexperiences a highly controlled expansion in thickness. The resul-tant porous graphene films can be folded and even crumpled butspring back to their original shape without creases, akin to anelastomeric material, after the applied stress is removed. Even afterthermal annealing at 1000 �C and ultralow liquid nitrogen tem-perature treatment, the folding performance of the porous gra-phene film is not compromised. The microvoids in the PGFs areresponsible for such excellent elastomeric scaffold-like foldingbehavior. We have shown here that such porous and elastic gra-phene films exhibit excellent electrical conductivity and EMIshielding capabilities. A considerably enhanced EMI shieldingperformance is achieved compared to the corresponding compactfilm due to the favorable foam structure for microwave absorption.

D. Lai et al. / Carbon 158 (2020) 728e737 737

The highest EMI SE value of 63.0 dB is enough to obstruct andabsorb 99.99995% of the incident radiation, which is among thehighest values of any known synthetic materials with similar smallthicknesses. The largest specific SE (SSE/t) of 49750 dB cm2 g�1 isthe highest of all the graphene-based shields and outperformsmostof the shielding materials known to date. These outstandingproperties of thinness, light weight, high foldability and elasticity,as well as excellent conductive and superior EMI shielding perfor-mance, make such porous graphene films a promising shieldingmaterial for application in foldable and wearable electronics,aircraft and spacecraft.

4. Experimental section

The experimental details and characterizations are provided inthe Supporting Information.

Notes

The authors declare no competing financial.

Declaration of competing interest

The authors declare that they have no known competingfinancial interests or personal relationships that could haveappeared to influence the work reported in this paper.

Acknowledgments

This work is supported by the China-Japan Research CooperativeProgram (2016YFE0118000), the Industry Leading Key Projects ofFujian Province (2019H0056), the National Natural Science Foun-dation of China (21806163) and the “Strategic Priority ResearchProgram (A)" of the Chinese Academy of Sciences (XDA23030301).

Appendix A. Supplementary data

Supplementary data to this article can be found online athttps://doi.org/10.1016/j.carbon.2019.11.047.

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