3D printing of CNT- and graphene-based conductive polymernanocomposites by fused deposition modelingCitation for published version (APA):Gnanasekaran, K., Heijmans, T., van Bennekom, S., Woldhuis, H., Wijnia, S., de With, G., & Friedrich, H.(2017). 3D printing of CNT- and graphene-based conductive polymer nanocomposites by fused depositionmodeling. Applied Materials Today, 9, 21-28. https://doi.org/10.1016/j.apmt.2017.04.003

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Applied Materials Today 9 (2017) 21–28

Contents lists available at ScienceDirect

Applied Materials Today

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D printing of CNT- and graphene-based conductive polymer nanocomposites byused deposition modeling

. Gnanasekaran a, T. Heijmans a,b, S. van Bennekom b, H. Woldhuis b, S. Wijnia b,

. de With a,∗, H. Friedrich a

Laboratory of Materials and Interface Chemistry, Eindhoven University of Technology, The NetherlandsUltimaker B.V., Geldermalsen, The Netherlands

r t i c l e i n f o

rticle history:eceived 1 March 2017eceived in revised form 6 April 2017ccepted 7 April 2017

eywords:

a b s t r a c t

Fused deposition modeling (FDM) is limited by the availability of application specific functional materi-als. Here we illustrate printing of non-conventional polymer nanocomposites (CNT- and graphene-basedpolybutylene terephthalate (PBT)) on a commercially available desktop 3D printer leading toward print-ing of electrically conductive structures. The printability, electrical conductivity and mechanical stabilityof the polymer nanocomposites before and after 3D printing was evaluated. The results show that 3D

D printingused deposition modelingolymer nanocompositesNTrapheneozzle wear

printed PBT/CNT objects have better conductive and mechanical properties and a better performancethan 3D printed PBT/graphene structures. In addition to that, printing more than one material (multi-materials) and challenges in using abrasive conductive fillers (i.e., CNT and graphene) are also discussed.Overall this study demonstrates that a commercially available desktop 3D printer can be used to fabricatelow-cost functional objects.

© 2017 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY license

. Introduction

Additive manufacturing, also known as 3D printing, is revolu-ionizing the world, mainly because computer designed objectsabricated by 3D printers can be more complex than convention-lly machined parts [1]. This allows for larger design flexibilitynd enables improved performance, often at reduced weight. Fromhe ever increasing number of additive manufacturing approaches,used deposition modeling (FDM) has pervaded most into science,ducation and home use. In FDM, the feedstock is a thermoplas-ic polymer filament, which is heated above its glass transitionemperature (Tg) and extruded through a movable nozzle in the–Y plane to form a 3D structure by layer-wise addition [2]. Whilelmost every base material, that is, metals, ceramics and poly-ers, can be 3D printed into complex shapes, nowadays a strong

emand exists for 3D printable materials having additional func-ionality [3]. Significant progress has been made in 3D printing

aterials with functionality, such as electrochemical electrodes [4],

hotochromatic materials [5], chromatographic devices [6], mem-rane modules [7], conductive composites and blends [8–14]. Withhese developments unique possibilities for creating multifunc-

∗ Corresponding author.E-mail address: g.dewith@tue.nl (G. de With).

ttp://dx.doi.org/10.1016/j.apmt.2017.04.003352-9407/© 2017 The Authors. Published by Elsevier Ltd. This is an open access article u

(http://creativecommons.org/licenses/by/4.0/).

tional devices, implants, and even organs, by 3D printing come intoscope [15,16].

Desktop 3D printing of functional devices opens the next level inthe market. The cost of manufacturing small series or unique partscan be reduced significantly, facilitating easy manufacturing evento the small-scale private sectors. However, the main limitationis the availability of application specific functional raw materialsfor specific 3D printers and the optimized printing conditions toobtain good esthetics (without surface roughness) as well as suffi-ciently good functional properties. Desktop 3D printing of polymernanocomposites (PNCs) with a relatively low amount of layeredsilicate fillers or conductive nanofillers, such as carbon nanotubes(CNT), graphene and metal particles, allows one to build objectswith multifunctional properties having good electrical conductiv-ity, thermal conductivity, mechanical strength, and stiffness at arelatively low cost [10,12,17]. In conductive PNCs, the electricalconductivity is governed by electrical percolation, which requires aminimum filler content (volume fraction) to convert an insulatingpolymer matrix into a conductive composite [18]. This minimumvolume fraction of nanofillers, referred as percolation threshold,strongly depends on factors like, shape and size distribution ofthe nanofillers [19,20], attractive interactions [21], and processing

methods (dispersion and agglomeration of nanofillers) [22]. Sev-eral studies have been reported on 3D printing of conductive PNCs,based on conventionally used 3D printing polymers like poly lactic

nder the CC BY license (http://creativecommons.org/licenses/by/4.0/).

22 K. Gnanasekaran et al. / Applied Materials Today 9 (2017) 21–28

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ig. 1. (a) Extruded PBT/CNT composite filament. (b) 3D printed monolayer of PBxtruded PBT/G composite filament. (e) 3D printed monolayer of PBT/G compositend white scale bars are 500 �m.

cid (PLA) and acrylonitrile butadiene styrene (ABS) [2,9,12,13,23].owever, development of new conductive PNC materials for a desk-

op 3D printer is highly desirable to achieve better printability,echanical properties and electrical conductivity.

Several factors need to be considered while printing conduc-ive composite materials. One of the major bottlenecks is nozzleam, caused by agglomeration of the nanofillers and resulting inoor printability and surface roughness [23]. Hence, in order toontrol these effects, the size and size distribution of the conduc-ive fillers as well as the printing parameters/conditions, such asrinting temperature, printing speed, residence time and printinged temperature, need to be optimized for printability. In addi-ion to that, certain physical properties of the polymer matrix andhe subsequent conductive PNC are required to obtain suitablerintability and functional behavior. The conductive PNC filamenthould have enough flexibility to be spooled after melt extru-ion and fed through the Bowden tube of the 3D printer withoutreaking for 3D printing. The 3D printed object should also haveufficient mechanical strength, stiffness, and chemical resistance tonsure the durability of the end-product for the specific application.he polymer material needs to withstand appropriate temperatureuring composite fabrication via melt extrusion and subsequentrinting. A material with a higher storage modulus requires 3Drinting at slower speed, or using a higher nozzle temperature.owever, if the printing temperature is too high or the residence

ime of the polymer in the hot end at slower printing speed isoo long, polymer degradation may occur, typically resulting inoids in the printed structure [24]. This also results in a lower elec-rical conductivity of the composite materials. Most importantly,he chosen printing temperature of the material (conventionallylightly above the melting temperature of the polymer) shouldatch with the 3D printer’s temperature regime. For instance, theltimaker 2 FDM printer used in this work has a processing win-ow of 150–260 ◦C. Based on these criteria (see SI, section I and

I), we chose polybutylene terephthalate (PBT) as the base poly-er matrix for the development of conductive PNC for 3D printing

pplications and investigated the change in electrical conductivity,rystallinity, morphology and viscoelastic properties of PBT/CNT

composite. (c) SEM image of the PBT/CNT monolayer illustrating the ridges. (d)M image of the PBT/G monolayer illustrating the ridges. Black scale bars are 1 cm

and PBT/graphene structures fabricated using the aforementionedcommercially available 3D printer (Ultimaker 2). Surface modifica-tion of the CNTs or the graphene could lead to a better dispersibilitybut will also affect the conductivity adversely and, hence, we optedfor using non-modified materials. In addition, we extend the scopeof printing multi-materials and discuss the potential degradation of3D printer parts (such as printer head and nozzles) upon 3D printingof abrasive materials like CNT and graphene.

2. Materials and methods

2.1. Materials

Multi-walled carbon nanotubes (CNT, � ∼ 2.1 g/cm3) weresupplied by Nanocyl, Belgium (Product No. C7000). Thermallyexpanded graphite (graphene, � ∼ 2.1 g/cm3) was prepared bymicrowaving hydrolyzed graphite bisulphite for 1 min at 2450 MHz(700 W) as explained by Arapov et al. [25]. Polybutylene terephtha-late (� ∼ 1.32 g/cm3, Tg ∼ 55 ◦C) was supplied by Sabic InnovativePlastics, The Netherlands.

2.2. Sample preparation

CNT and graphene dispersions were prepared by adding 50 mgof CNT in 100 ml of isopropanol and sonicating for 2 h in an ice bath(Bransonic 1510) to prevent any rise in temperature. Appropriateamounts of PBT powder were added to the mixture and furthersonicated for an hour. Subsequently, isopropanol was allowed toevaporate at 90 ◦C using a water bath while stirring the solutionvigorously to avoid macroscale inhomogeneity and phase separa-tion. However, some degree of inhomogeneity could not be avoided,especially in the last stages of drying due to the higher viscos-ity of the mixture. Nevertheless, this method is proven to resultin a good dispersion [26]. The resulting PBT powder coated with

conductive fillers was dried at room temperature for 24 h. The com-posite powder was fed into a DSM Xplore 15 ml mini extruder. Thetemperature of the mini extruder was set to 240 ◦C and allowed tomix for 5 min at 50 rpm. Thereafter the melt mixed composite was

ed Materials Today 9 (2017) 21–28 23

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K. Gnanasekaran et al. / Appli

xtruded via a die with a diameter of 3 mm and allowed to quencho room temperature.

An Ultimaker 2 3D printer was used for 3D printing the compos-te filaments. The composite filaments were printed into squared

onolayers with dimension of 3 cm × 3 cm × 0.3 mm. Brass nozzlesith 0.4 mm diameter were used. The printing head was set at the

ate of 20 mm/s. The nozzle temperature varied between 240 ◦Cnd 260 ◦C depending on the amount of conductive filler contentsee SI, section III). Double-sided tape was used as a printing sub-trate with the printing bed temperature set at 70 ◦C to prevent therinted material from warping (see SI, section IV). While printing,ulti-materials using dual extrusion head, the nozzle temperature

f PLA and PBT/CNT was set at 200 ◦C and 260 ◦C, respectively, andhe printing bed temperature was varied between 65 ◦C and 80 ◦C.

.3. Characterization

DC conductivity measurements were carried out using a four-oint setup with parallel electrodes separated by 5 mm. The currentas applied through the outer electrodes by a Keithley 237 sourceeasure unit, and the potential difference was measured between

he inner electrodes by a Keithley 6517A electrometer. Detailednformation on the electrical measurements is given in the SI, sec-ion V. Scanning electron microscopy (SEM) images were acquiredsing a Quanta 3D SEM (Thermo Fischer Scientific) operated at

keV on a secondary electron detector. Profilometry measurementere performed on a Veeco Dektak 150 surface profiler instrument

t a speed of 80 �m/s. Thermal stability and the transition temper-tures were determined using a differential scanning calorimetert 10 ◦C/min (DSC Q2000 – TA Instruments) and thermogravimetricnalysis (TGA Q500 – TA Instruments) at 2 ◦C/min. The viscoelasticroperties were measured using a dynamic mechanical analyzerDMA Q800 – TA Instruments) at a constant frequency of 1 Hz.ll wide-angle X-ray scattering experiments were performed on aAXSLAB’s Ganesha instrument equipped with a GeniX-Cu ultralowivergence source producing X-ray photons with a wavelength of.54 A and a flux of 1 × 108 photons/s. Scattering patterns were col-

ected using a Pilatus 300 K silicon pixel detector with 487 × 619ixels and with a pixel size of 172 �m2 placed at a sample-to-etector distance of 91 mm. The beam center and q-range werealibrated using the diffraction peaks of silver behenate. Conver-ion of 2D into 1D data was carried out with Saxsgui V2.13.01 usingn automated beam stop mask. Samples were exposed to X-rays for800 s. Electron transparent samples were prepared by microtom-

ng the 3D printed samples into ∼500 nm thick sections using aiatome 35◦ 3.0 mm diamond knife mounted on a Reichert-Jungltracut-E ultramicrotome and collected on a continuous carbonlm coated 200 mesh copper TEM grid. All scanning transmis-ion electron microscopy (STEM) images were acquired using aitan electron microscope (Thermo Fischer Scientific) operated at00 keV with a convergence angle of 4 mrad on a Fishione HAADF-TEM detector at camera length of 450 mm.

. Results and discussions

.1. 3D printing of conductive PBT composites

Several PBT composite filaments with varying volume fractionsf CNT and graphene were prepared by melt extrusion process andubsequently 3D printed to monolayers to determine the electri-al percolation threshold. Images of prepared composite filaments

nd 3D printed monolayers are shown in Fig. 1. All the graphene-ased filaments and 3D printed monolayers appeared to be roughnd brittle as compared to CNT-based composites. No differencen surface quality was obtained by slightly changing the nozzle

tion for 3D printed polymer nanocomposites. Note that the electrical conductivityof pure PBT is ∼1 × 10−15 S/cm and of the pure CNT/graphene is ∼1 × 105 S/cm.

temperature or the printing speed. Low convergence angle scan-ning transmission electron microscopy (LC-STEM) imaging of thePBT/G composite (Section 3.3) reveals that the graphene plateletswere dispersed finely in the PBT matrix without any significantagglomeration. However, during processing at elevated temper-ature, the evaporation of moisture in the thermally expandedgraphite platelets (graphene) results in the formation of voids alongthe surfaces of the printed material, thereby causing surface rough-ness and brittleness.

The surface of the 3D printed samples was examined by SEM.Multiple 3D printed layers are clearly seen and uniformly placedwith layer width closely matching to the nozzle diameter. The inter-face of the neighboring layers was merged to form ridges betweenthe layers. This is caused by the nozzle touching the printed layersand pushing the printed material aside. Such discrete interfaceswere distributed at some regions in some of the graphene-basedmonolayers as the molten material spreads across covering severallayers as shown in Fig. 1f. This can be avoided by further reducingthe temperature (slightly increasing the viscosity); however, it maycause the nozzle jam and results in loss of printing resolution.

3.2. Electrical conductivity analysis

Fig. 2 shows the electrical conductivity plotted as a func-tion of volume fraction of conductive fillers for the various 3Dprinted monolayers. The minimum amount of conductive filler (ϕc)required to form conductive composite can be estimated by elec-trical percolation theory [27,28]. A detailed analysis of electricalpercolation is given in the SI, section VI. Based on the analysis of fil-ament formation and electrical percolation, we suggest that at least∼0.49 wt.% of CNT (∼0.31 vol.%) and ∼5.2 wt.% (ϕc ∼ 3.3 vol.%) ofgraphene is required for the FDM fabrication of PBT/CNT and PBT/Gconductive filaments, respectively. Previous studies using conven-tionally used polymer matrices reported a similar level of aboutϕc ∼ 0.67 wt.% for PLA/CNT composites [12] and of about ϕc ∼ 2 wt.%for ABS/G composites [23]. The differences could be related to theuse of different preparation methods and/or different raw mate-

rials. The difference in PBT/G and PBT/CNT composite behavioris mainly caused by the dimensionality of the filler particles (1Dfor CNT, 2D for graphene) [29] and potential aggregate formationfor graphene which increases the electrical resistance [26], possi-

24 K. Gnanasekaran et al. / Applied Materials Today 9 (2017) 21–28

Fig. 3. STEM images of 3D printed PBT/CNT and PBT/G composites. White dots represent the gold markers used for tracking. Scale bar: 500 nm.

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ig. 4. (a) 1D WAXS diffraction patterns of filaments and 3D printed PBT/CNT (ϕ ∼

BT/CNT composite (ϕ ∼ 0.036).

ly reinforced by air entrapment between the exfoliated graphiteakes [25].

The anisotropy in distribution of the conductive fillers wasbserved by measuring conductivity at various angles betweenhe deposition direction of the FDM print head and the measur-ng probes. For instance, by placing the 4 probes in line with therinted layers – parallel measurement, 90◦ to the printed layers

perpendicular measurement, and also at intermediate anglessee Fig. S4). In PBT/CNT composites the anisotropy in conductiv-ty is a factor of ∼7 at the percolating volume fraction (ϕ = 0.0062),ecreasing to a factor of ∼1.3 well above the percolating volume

raction (ϕ = 0.036). This factor converges to 1, as a result of a con-uctive network formed across the printed layers that results in

lesser effect of additional CNT paths to the overall conductivity.ere it needs to be mentioned that ridges were formed between

he connecting interfaces of the layers during printing, which mayause the observed variation in conductivity. Therefore, the pre-ise variation in the thickness and the profile of the 3D printedonolayers were measured using a surface profiler (see, SI, section

II). For the 3D printed PBT/CNT ϕ = 0.036 monolayer, the thick-ess increases by a factor of 1.38 due to the ridges on the topurface. This factor matches closely with the variation in conduc-ivity (∼1.3). This shows that the apparent increase in conductivity

hile measuring parallel to the printed lines is due to the uncor-

ected increase in thickness of the ridges at high volume fractionomposites. At the percolating threshold of conductive fillers, thelectrical conductivity is limited by network connectivity across

and PBT/G composites (ϕ ∼ 0.041). (b) 2D WAXS diffraction patterns of 3D printed

the interface between neighboring deposition lines. The changein resistance as a function of angle of the measurement (anglebetween the probes and the sample) is shown in Fig. S7. The slopeof the curve changes depending on the angle of the measurement.For small angles, the resistance increases slowly, but above about45◦ the slope significantly increases. This shows that the connec-tivity at the interfaces of the neighboring printed lines was limitedand the interface between the 3D-printed layers predominantlyacts as series resistors. By changing the angle of measurement, thenetwork acts like parallel resistors (layers), connected by severalseries resistors (interface between layers). This also suggests that, inorder to get a uniform connectivity across the FDM deposition lines,it is advisable to have concentrations (of conductive nanofillers)far above the percolating volume fraction. Hence, we chose 2 3Dprinted composite monolayers that are percolating as well havenearly the same amount of conductive fillers – PBT/CNT ϕ ∼ 0.036and PBT/G ϕ ∼ 0.041 − for further analyses.

3.3. Morphological analysis

The distribution of conductive fillers within the 3D printedstructures was characterized by low convergence angle scanningtransmission electron microscopy (LC-STEM) [30,31] of relatively

thick (∼500 nm) sections with large area imaging toolbox [32].Detailed information on large-scale LC-STEM imaging and theacquired representative images is given in the SI, section IX. A detailfrom the large-scale images of 3D-printed samples of PBT/CNT

K. Gnanasekaran et al. / Applied Materials Today 9 (2017) 21–28 25

Fig. 5. (a) Cooling scans of DSC measurements. (b) Heating scans of DSC measurements. Arrow indicates the recrystallization. (c) Crystallinity as calculated from meltingenthalpies for various samples.

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ig. 6. (a) TGA curves of PBT/CNT (ϕ ∼ 0.036) and PBT/G (ϕ ∼ 0.041) composites. (rinted PBT/CNT (ϕ ∼ 0.036) and PBT/G (ϕ ∼ 0.041) composites.

ϕ ∼ 0.036) and PBT/G (ϕ ∼ 0.041) composites is shown in Fig. 3a and, respectively. The graphene platelets (marked in circles in Fig. 3b)roduce very low contrast as compared to the CNTs. Both graphenend CNTs were uniformly distributed and no large agglomeratesere found. The light regions represent the crystalline part of the

BT matrix. The amount of crystallinity appears to be roughly sameor both PBT/CNT and PBT/G composites despite the slightly higherolume fraction of graphene in PBT/G composite. It is known thatn semi-crystalline polymers like PBT, interactions between theraphitic rings of the conductive filler (CNT and graphene) andhe benzene rings of the terephthalate moiety induces crystallinetacking of polymer chains [33,34]. In the PBT/CNT composites, theNTs were predominately located within the crystalline regions,

ndicating the strong tendency of PBT to crystallize on CNT walls.The morphology of the 3D printed composites was character-

zed by wide-angle X-ray scattering (WAXS). Fig. 4 shows the WAXSiffraction patterns of 3D-printed samples of PBT/CNT (ϕ ∼ 0.036)nd PBT/graphene (ϕ ∼ 0.041). Because of the very low concentra-ion of conductive fillers, their characteristic peaks (for instance,02 plane of graphene at 26.5◦) are not visible. The peaks associ-ted with the crystalline lattices of PBT are sharper in PBT/CNT thann PBT/G. During the 3D-printing process, at each pass of the noz-le in the X–Y plane, the already printed layers were reheated. This

llows for uniform heat transfer and relatively less super-coolingnd induces the crystallization. However, such a phenomenon wasot observed in PBT/G composites. This suggests that CNTs possess

higher nucleation ability as compared to graphene. A possible

rage modulus E′ and loss modulus E′ ′ plotted as a function of temperature for 3D

explanation is that the polymer chains require more time to adjusttheir conformations across the surface area of the graphene sheetsas compared with the CNTs and therefore the rate of crystalliza-tion and the percentage crystallinity decrease [35]. In addition,the isotropic distribution of intensities along the diffraction rings(Fig. 4b) suggests that the flow conditions in the micro-compounderand the 3D printer do not induce preferentially oriented morpholo-gies [22].

3.4. Crystallization behavior

To access the change in crystallinity during the 3D printing pro-cess, differential scanning calorimetry (DSC) analyses of the 3Dprinted samples were carried out and the results are shown in Fig. 5.Fig. 5a shows the cooling curves where a shift of the crystallizationpeak toward higher onset temperatures for the 3D-printed com-posites as compared to pure PBT is observed. This clearly illustratesthe above mentioned phenomenon that the conductive fillers actas nucleating sites and induce crystallization of the polymer matrix[36,37]. The increase in crystallization temperature is not propor-tional to the amount of CNT added (see SI, section X), illustratingthat above a certain critical filler content, the onset of crystallizationlevels off with further increase in filler content. Fig. 5b shows the

corresponding heating curves. The crystals of the pure PBT meltedat lower temperatures than the 3D-printed composites, whichindicates a difference in the distribution of crystalline lamellaethickness in the 3D-printed composites. In addition, simultaneous

26 K. Gnanasekaran et al. / Applied Ma

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ig. 7. 3D printed objects using a dual head extruder. The left object is printedith constant bed temperature while the right object is printed with variable bed

emperature. Scale bar: 1 cm.

elting and recrystallization of the pure PBT filament was observedshown by an arrow), i.e., the thinner lamellae melted, resulting inhicker remaining lamellae with higher melting point. This phe-omenon was not observed in the 3D-printed composites; instead

broad shoulder is observed indicating the broad size distribu-ion of the lamellae. The degree of crystallinity was calculated fromhe melting enthalpies [38], and the results are shown in Fig. 5c. Byncreasing the conductive filler content, for the PBT/CNT compositehe crystalline fraction shows a maximum. At higher volume frac-ion both composites show a decrease in crystallinity. This decreaseuggests that the presence of large number of nucleating sites dis-ersed over the entire polymer matrix immobilizes the polymerhains and limits the formation of crystalline lamellae. Moreover,n increase in conductive filler content could increase the prob-bility of agglomeration, which probably decreases the degree ofrystallinity as well.

.5. Thermal stability and viscoelastic behavior

To assess durability, thermal and mechanical stability, ther-ogravimetric analysis (TGA) and dynamic mechanical analyses

DMA) of 3D-printed samples were conducted and the results arehown in Fig. 6. Fig. 6a shows the TGA curves of pure PBT andhe 3D-printed composites performed under oxidative conditions.egradation of PBT occurred in two steps. A weight loss of 5% occurst 303 ◦C, 339 ◦C, and 332 ◦C for the PBT, PBT/CNT, and PBT/G com-osite, respectively, while they completely decompose at about50 ◦C. The addition of conductive fillers, not surprisingly, improveshe thermal stability, as both the onset and the maximum degra-ation temperature shift to a higher value. The presence of theonductive fillers facilitates the heat conduction and acts as bar-ier to inhibit the emission of decomposition products during theombustion [39]. The total mass loss of the PBT/CNT and PBT/Gomposite was 94% and 99.5% at 550 ◦C, respectively. This showshat the PBT/CNT composite has higher thermo-oxidative stabilityhan the PBT/G composite.

Fig. 6b shows the storage E′ and loss modulus E′ ′ of theD-printed composites. The stiffness of the PBT/CNT 3D-printedonolayer was significantly higher than that of the PBT/G com-

osite; more precisely, the storage modulus was 28% higher for

he PBT/CNT at 25 ◦C. Since the crystalline fraction is higher forraphene composites (see Fig. 5c) as compared to CNT composites,he observed inverse relation in storage modulus was unexpected.ften the formation of an interphase between the surface of the

terials Today 9 (2017) 21–28

conductive fillers and the polymer chains is referenced for out-standing mechanical properties [40,41]. In addition, reinforcementis usually described in terms of the critical filler length lc, belowwhich the filler cannot transfer its stiffness or strength proper-ties to the polymer matrix, and results in premature failure of thecomposite. It is generally accepted that, in order to obtain a goodreinforcement, the filler length should be at least 10 times lc [26].In addition, the loss factor E′ ′/E′ was observed to be ∼0.1 for boththe PBT/CNT and PBT/G samples. Previous studies on 3D printedgraphene-based ABS composites reported E′ ′/E′ ∼ 0.5 at comparablevolume fraction. This shows that our PBT-based composites behavemore elastically.

4. Further discussions

4.1. 3D printing of multi-materials

Printing multiple materials would take the development offunctional devices to the next level. Recently, Gonzalez et al. [8]showed printing electrical circuit like structures of a conductivepolymer nanocomposite on an insulating base of the same poly-mer. Here we show the printing of 2 of different materials, which wecall “integrated printing”. This requires the use of a multi-extrusionhead and the printing parameters/conditions are very critical asthey need to be optimized for more than one material, especiallywhen the Tgs of the materials are significantly different. For propernecking, diffusion, and adhesion between polymers of different Tgs,the temperature of the printing polymer should be higher so thatit heats the already printed polymer layer (usually the printingtemperatures should be higher than the Tg of the already printedpolymer) [42].

For our multi-materials printing tests, we used our PBT com-posite and conventionally used pure PLA. Using the same printingparameters as used for printing monolayers, significant thermalstress at the interface of the materials (indicated by the arrow inFig. 7, left sample) arise due to the difference in the coefficient ofthermal expansion of different materials. Hence, a variable printingbed temperature (from 65 ◦C to 80 ◦C upon increasing layer thick-ness) was used to overcome these difficulties (Fig. 7, right sample).The resulting electrical resistance was also measured and appearedto be more than 2 times lower (∼600 �) than for the constantbed temperature sample (∼1350 �). We also observed that PLAbecomes softer and tends to sag at 100 ◦C printing bed temperature,while the PBT/CNT composite does not bond with PLA at a printingbed temperature as low as 50 ◦C. Hence, choosing an appropriateprinting bed temperature while printing multi-materials is crucial.

4.2. Nozzle wear

The exceptionally high specific Young’s modulus of CNTs andgraphene (higher than that of diamond) is expected to lead to noz-zle wear during the 3D printing process [43,44]. Fig. 8 shows the3D printer’s nozzle surface before and after printing several CNT-and graphene-based composite materials. Upon prolonged print-ing, the brass nozzle was abraded both inside the nozzle as well ason the front surface where it touches the printed object (Fig. 8c).A SEM image of the surface of the corresponding 3D printed PBT/Gcomposite material is shown in Fig. 8d. The surface is clearly roughand the ridges were irregularly placed with loss of printing resolu-tion due to the degradation of the nozzle and this effect potentially

degrades the esthetics and the functional properties. Therefore,for printing abrasive materials like CNTs and graphene, FDM noz-zles made from harder materials, such as silicon carbide, may berequired; however, because of the difference in the thermal con-

K. Gnanasekaran et al. / Applied Materials Today 9 (2017) 21–28 27

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ig. 8. Optical micrographs showing the surface of a 3D printing nozzle before andfter printing ∼1.5 m of PBT/CNT. (d) SEM image of PBT/G composite printed with a

uctivity of harder materials, the printing conditions need to beptimized appropriately [45].

. Conclusions and outlook

Desktop 3D printing has an enormous potential to fabricateomplex structures that can be applied for large-scale manufac-uring at low-cost. In this work, we have shown the preparation,lament extrusion and FDM-based 3D printing of CNT- andraphene-based conductive polymer nanocomposites into func-ional 3D model structures. More specifically, PBT was used aase polymer, as compared to the commonly employed poly-ers PLA and ABS, in order to obtain highly functional as well asechanically robust structures. The printability, electrical conduc-

ivity, crystallinity, morphology and viscoelastic properties of theD printed structures were evaluated. The analysis showed thatBT/CNT 3D printed structures have better functional propertieselastic behavior and conductive properties) as well as esthet-cs than the PBT/G 3D printed structures. The optimum printingarameters/conditions do not change much for the PBT/CNT andBT/G composites due to the very low volume fraction dispersed inhe PBT matrix.

We illustrated the use of multiple printing heads such that morehan one material can be printed at once. We expect that this willecome the standard for the fabrication of functional devices. For

nstance, it shows the potential to print conductive structures likelectrodes within a 3D-printed insulating matrix. We have alsohown that care must be taken while printing abrasive materi-ls that could potentially damage the printing nozzle, resultingn loss of printing resolution and causing undesirable properties.here are several research questions that still need to be addressed.ew materials need to be optimized for 3D printing grades, such

hat they have the proper flow rate, crystallization and shrinking

roperties, in addition to their functional properties. Significantrogress is required in nozzle design for printing abrasive mate-ials and also to reduce the chance of clogging of conductive fillers.o-extruding nozzles instead of double nozzles could take the

printing. (a) Unused nozzle. (b) Nozzle after printing ∼10 cm of PBT/G. (c) Nozzleded nozzle.

development to the next level, but there are several issues thatrequire attention upon using co-extruding nozzles. For instance,the nozzle has to be purged while switching to a different mate-rial and only materials with similar printing temperatures can beused. In addition, the nozzle material and nozzle geometry is prob-ably limited to one design, which forces one to print, for example,abrasive material and conventional 3D-printing material with thesame nozzle. Finally, and more importantly, health and safety issuesupon printing ultra-small nanoparticles need to be assessed. Over-all, we believe that this study demonstrates that 3D printing is aviable technology for manufacturing functional structures at verylow-cost, and shows that unconventional polymers/materials canbe exploited and optimization to obtain better printability (esthet-ics) and functional properties, and will push our materials selectionprocess and applications.

Acknowledgements

The authors would like to thank mr. M.M.R.M. Hendrix for theSAXS measurements. This project has received partially fundingfrom the European Union’s FP7-PEOPLE-2010-ITN program grantagreement number 264710, the Horizon 2020 research and inno-vation program under grant agreement No 696656 and from theDPI grant number EU-FP-002.

Appendix A. Supplementary data

Supplementary data associated with this article can be found,in the online version, at http://dx.doi.org/10.1016/j.apmt.2017.04.003.

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