Multiscale Engineering of Carbon Nanotube Fibers

CHAPTER

MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

6

Bel
en Alema´n, Juan J. Vilatela IMDEA Materials Institute, Getafe, Spain

CHAPTER OUTLINE 1 Introduction ................................................................................................................................... 113 2 Synthesis of CNT Fibers: Toward Molecular Control .......................................................................... 114 2.1 Control of CNT Number of Walls and Morphology by Means of the Promoter ........................ 114 2.2 Catalyst Structure and Growth Model ............................................................................... 119 2.3 Chiral Angle Distribution and n,m Indices Assignation ...................................................... 123 3 Structure-Properties: A Complex Hierarchical Structure .................................................................... 127 3.1 Control of Orientation Along Fiber Axis ............................................................................ 127 3.2 Structural Studies .......................................................................................................... 127 3.3 Structure and Tensile Properties ..................................................................................... 136 4 Low-Dimensional Properties ............................................................................................................ 138 5 Conclusions and Outlook ................................................................................................................ 142 References ........................................................................................................................................ 143

1 INTRODUCTION Continuous macroscopic fibers of carbon nanotubes (CNTs) were first produced via coagulation spinning in 2000 [1], followed rapidly by other methods: liquid-crystal spinning [2], drawing from CNT forests [3], or directly spinning from the gas phase [4]. The following 18 years have seen enormous efforts to optimize fabrication processes, establish clear structure-property relations, and develop applications and upscale spinning facilities. This chapter deals primarily with the first two aspects. In particular, it focuses on the direct spinning process. This method consists in drawing a web (aerogel) of CNTs directly from a tubular reactor where they continuously grow by chemical vapor deposition (CVD) at around 1200°C (Fig. 1A). As the web is drawn out of the reactor, it densifies as a fiber/ film, which can then be wound onto a bobbin. The process has an attractive simplicity: it transforms a cheap carbon source (e.g., methane) into a high-performance fiber at rates above 50 m min
1. Chemical engineering efforts over the last years have led to reproducible spinning of uniform kilometer-long fibers (Fig. 1B) in laboratory [5] and industrial settings [6]. This has unlocked many applications in Nanotube Superfiber Materials. https://doi.org/10.1016/B978-0-12-812667-7.00006-9 Copyright # 2019 Elsevier Inc. All rights reserved.

113

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CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

(A)

(B)

Precursors

1200°C

Catalyst nanoparticles

Entangled carbon nanotubes (aerogel)

Carbon nanotube fibre

Liquid

FIG. 1 Fiber spinning directly from the gas phase during CNT growth by CVD. (A) Schematic of a CVD fiber spinning reactor and (B) photograph of a bobbin with 1 km of CNT fiber.

composites and energy [7, 8] and, very importantly, enabled a gradual transition from trial-and-error synthesis recipes to ground material science strategies for synthesis-structure-property control.

2 SYNTHESIS OF CNT FIBERS: TOWARD MOLECULAR CONTROL CNT fiber produced by different methods has mechanical properties already in the high-performance range and mass-basis electric conductivity close to copper but still modest compared with the properties of individual CNTs. Further improvements in properties require fibers made up of CNTs as similar as possible in terms of diameter, number of layers, metallicity, and chiral angle. These have to be realized in the macroscopic CNT fiber, with diameter in the range of 5–100 μm and continuous length (Fig. 2), implying that there are around 106 individual CNTs per cross section and thus about 109 in a meter of sample.

2.1 CONTROL OF CNT NUMBER OF WALLS AND MORPHOLOGY BY MEANS OF THE PROMOTER The challenge is to produce continuous fibers with uniform molecular structure and to have the possibility of tailoring composition covering the whole CNT “molecular” spectrum from single-walled to multiwalled CNTs (SWNTs and MWNTs, respectively). In the case of direct spinning process, it was early observed that the addition of sulfur (S) to the CVD reaction leads to large increases in CNT growth rate and length relative to conventional low-temperature CVD (around 800°C), which are key for the association of the in-forming CNTs as an aerogel in the gas phase and the subsequent extraction as continuous fiber [4]. It was also recognized that S plays a critical role in the reaction by limiting C diffusion to the catalyst surface and accelerating CNT growth [9].

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FIG. 2 (A) Optical and (B) scanning electron micrographs of CNT fibers.

FIG. 3 Electron microscopy images of nanotubes and nanotube bundles within the fiber showing the increase in the number of walls when increasing thiophene in the precursor mixture: (A) 0.1 wt%, (B) 0.4 wt%, and (C) 1.5 wt%.

By varying the concentration of thiophene (S or S/C) in the precursor mixture, it is possible to tailor the type of CNTs that make up the fiber [5]. Fig. 3 shows Raman spectra for different concentrations of thiophene in the precursor mixture. The graph shows the progression from SWNTs to MWNTs with increasing S in the reaction. For midconcentration (0.2–0.4 wt%), the nanotubes are predominantly large diameter and have few layers (2–4) and therefore autocollapse and form closed-packed stacks of graphitic tubular ribbons at turbostratic separation [10, 11] (Fig. 3B). Raman spectroscopy also enables the identification of other features of SWNTs: radial breathing modes ωRBMs  100–300 cm
1, split of G peak into G+ and G
ωG  1590 cm
1, and a very low ID/IG ratio. Similarly, MWNT evidences are present by a lower resonance, broader G peak, and a higher ID/IG ratio [12]. In this sense, note that the ID/IG ratio increases from 0.005 to 0.23 when going from SWNT to MWNTs not implying a lower degree of graphitization (Fig. 4A). A comparison of spectra in absolute units (Fig. 4A inset) confirms that the D band remains fairly constant whereas the G band dramatically loses intensity with the loss of resonance when going to CNTs with more layers. Current work in our group is exploring other characterization techniques to determine sample crystallinity. The effect of increasing the number of walls in CNTs also produces an upshift of the 2D (G0 ) band [13] (Fig. 4B), similar to that observed in graphene samples, although in the case of CNTs the addition of layers has an effect not only on the interaction between layers but also on curvature.

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FIG. 4 Raman spectroscopy data showing the evolution from SWNT to MWNT with increasing S promoter during CNT fiber synthesis. (A) Normalized full spectra showing SWNT features (RBM and G split) at low thiophene concentration and the loss of resonance when increasing the number of walls (inset). (B) Plot of Raman 2D (G0 ) peak position for samples produced with different amounts of sulfur precursor. Error bars represent standard deviation. The upshift with increasing sulfur is due to the increment in the number of layers of the CNTs.

FIG. 5 (A) Effect of increasing sulfur precursor on fiber linear density and on the relative mass per unit length of individual CNTs (CNT linear density) determined from TEM. (B) Reaction yield (carbon input/carbon output  100) against thiophene concentration.

The increment in the number of walls through a higher S/C concentration in the reaction has a direct effect on the fiber linear density. Interestingly, such increase follows a similar trend as the increase in individual CNT linear density estimated from TEM measurements (Fig. 5A). This implies that the use of more sulfur in CVD reaction does not result in the formation of more CNTs but only in the increase of their diameter and number of walls. The reaction yield can be calculated from the knowledge of the fiber linear density, winding rate (fixed at 7 m min
1), and precursor throughput, which in terms of injected carbon/output carbon goes from 0.7% to 9% from the data in Fig. 5B (yield (%) ¼ linear density (g km
1)  17.11 (km g
1)). The model for CNT growth in the direct spinning process is that ferrocene decomposes and forms Fe catalyst nanoparticles that saturate rapidly with C in their trajectory through the reactor. Most of them (99.95% of the particles) encounter S when their C concentration is already very high. This

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FIG. 6 Iron-rich corner of the isothermal 1400°C section of the Fe-S-C ternary diagram [5] showing the location of the active and inactive catalyst nanoparticles based on experimentally approximated composition.

corresponds to a region of the equilibrium Fe-S-C phase diagram (Fig. 6) with two immiscible liquids (one S-rich L1 and one C-rich L2) in equilibrium with solid carbon C(s) (i.e., graphite). Thus, catalyst “poisoning” occurs by encapsulation of a graphitic carbon layers before a graphitic cap is lifted off to start CNT growth. Remnant C in the supersaturated particles precipitates upon cooling [14]. In contrast, the very small proportion of active particles (0.05%) corresponds to those in which S diffusion occurs at lower C concentrations where the liquid(s) is (are) not equilibrium with C(s), and therefore, the incoming carbon would upon saturation of the catalyst particle get extruded to form a CNT. As the amount of available S in the reactor is increased, more particles become active, and the probability of them coalescing increases and so does their size [15]. Their larger diameter implies that they can accommodate a few additional layers for the same composition as the smaller ones. During CNT growth, the rate of incoming C scales as particle radius squared [9], whereas the rate of extrusion scales as number of layers (n) times radius; thus with increasing size of the active catalyst particle, the greater availability of C can produce the formation of additional layers at the same elemental composition. In the past few years, selenium (Se) or tellurium (Te) has been used as alternative promoters in the direct spinning process leading to a similar reaction as when using S [16]. All these promoters have the same valence shell electronic configuration (group 16) and play an equivalent role in the CVD reaction: adsorbing into the catalyst surface, reducing its surface tension, assisting in stabilizing the edge of nascent tubes, and promoting the growth of graphitic tubular layers. These roles are consistent with the findings in similar systems studied in metallurgy and catalytic methanation reactions. Irrespective of the choice of promotor, increasing its concentration relative to that of C leads to an increase in the CNT number of layers. However, Se-grown CNTs present an unusually large diameter (D) for the number of layers (n) that favors the formation of collapsed tubes (Fig. 7B). The ratio n/D is

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FIG. 7 (A) Layer-to-diameter ratio (n/D) of nanotubes grown by S and Se promoter. The lower n/D in Se-grown CNTs favors the formation of (B) collapsed nanotubes (collapsed DWNT with a diameter around 8.75 nm).

FIG. 8 Self-folding of large-diameter collapsed SWNTs. (A) HRTEM micrograph of a SWNT folded over itself. The inset shows a higher magnification view of the folding area. (B) Comparison of energy per atom of individual SWNT as a function of diameter in round, collapsed, and self-folded morphologies. The theoretically stable regions are shadowed, and examples of tubes included (30,30) round CNT of 4.03, (60,60) collapsed CNT of 8.13, and self-folded (100,100) of 13.4 nm diameter.

fairly constant across Se contents though around half that of S-grown CNTs (Fig. 7A). For example, at 3  10
4 atomic ratio promoter, the SWNTs grown from Se have a diameter of 6.2 nm, whereas those grown using S have a diameter of 2.45 nm. The low n/D ratio of Se-grown CNTs favors their self-collapsed into ribbons. In the case of CNTs using Se promoter, the large diameter of SWNTs and double-walled CNTs (DWNTs) over 10 nm implies that they can collapse and self-fold and still remain stable due to the additional layer contact and relatively small bending energy penalty (Fig. 8A). A molecular dynamics (MD) model of individual armchair SWNTs predicts that they are stable as cylindrical tubes up to 5 nm, as collapsed tubes between 5 and 7.3 nm, and as self-folded ribbons for larger diameter (Fig. 8B). This simulation confirms the stability of the conformation observed experimentally, although actual cutoff diameters depend strongly on bundle composition, chemical environment, and temperature.

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2.2 CATALYST STRUCTURE AND GROWTH MODEL Although it is difficult to directly confirm that the catalyst particles are in molten state during the CVD reaction at 1250°C, postsynthesis analysis indicates a structure whose formation is consistent with a molten stage. The particles have a core-shell structure, with a core being predominantly distorted FCC Fe (austenite) or martensite, both nonequilibrium phases with remnant C in the crystal lattice trapped after rapid cooling and with ferromagnetic properties [17]. Cementite is present in very large inactive particles, and no BBC Fe or iron oxides are found (Fig. 9). Sulfur is located in the shell of the particle [9], as a subnanometric surface layer of intermetallic Fe-S [5, 18], and does not get incorporated into the nanotube during its growth, as confirmed by X-ray photoelectron spectroscopy (XPS) (Fig. 10).

Intensity (a.u.)

(A)

(D)

Section (b) Fe FCC (200) CNT (004)

Fe FCC (220) CNT (110)

Section (c)

40

60 2q (degrees)

50

70

Martensite (C0.07 Fe1.93) (101) & (110)

Intensity (a.u.)

Fe FCC (111) CNT (100)

(E)

Fe BCC (111) C Fe3 (031) & (112)

40

42

Martensite (C0.07 Fe1.93) (002)

(C)

44 46 2q (degrees) (200)

48

Martensite (C0.07 Fe1.93) (200)

50

(F)

Intensity (a.u.)

(B)

80

60

62

64

66

68

70

2q (degrees)

FIG. 9 Catalyst particle crystal structure. (A) High-resolution TEM micrograph of an active catalyst particle with its corresponding FFT and inverse FFT images and the assignation to FCC Fe by means of comparison with the simulated (B) SAED pattern and (C) FCC Fe crystal structure. XRD patterns showing the composition of CNT fibers: In the range (D) from 40 to 80 degrees, the presence of FCC Fe and graphitic peaks related to CNTs are clearly observed. Magnified data in the ranges of (E) 40–50 degrees and (F) 60–70 degrees clearly show the presence of FCC Fe and martensite with 0.77 wt% C as the major constituents of the catalyst while confirming the absence of BCC Fe.

120

(B)

(Fe2p3/2)Fe (Fe2p1/2)Fe

730

725 720 715 710 Binding energy (eV)

735

730

725

700

720

715

710

705

700

Binding energy (eV)

S2p / Fe-S compounds

(E)

705

(Fe2p1/2)Fe

172

(C)

0.2 wt% of thiophene

(Fe2p3/2)Fe

(Fe2p1/2)Fe-S

Intensity (a.u.)

(A)

Intensity (a.u.)

Intensity (a.u.)

(D)

735

(Fe2p3/2)Fe-S

168

164

Binding energy (eV)

C 1s C 1s fitted Deconvolution components 284.5 eV (sp2) 285.7 eV (sp3) 290.5 eV (p–p*) Shirley background

160 300

(F)

295

290

285

280

Binding energy (eV)

FIG. 10 (A) STEM micrograph and (B) EDS map showing the presence of S at the surface of the Fe-rich catalysts. The intensity line profile across the particle (C) shows higher S concentration at the particle edged, implying higher concentration of S at the surface of the particle. XPS spectra show the presence of sulfur that is associated with the catalyst but not incorporated into the CNTs as a dopant: (D) Fe is observed in a metallic state, associated with S; (E) the signal from S further confirms its presence in the sample and (F) the deconvolution of C-related peaks showing the predominance of sp2 C]C and the absence of S doping.

CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

Thiophene (wt%) 0.2% 1%

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FIG. 11 Thermodynamic analysis of the Fe-S-C catalyst phase composition during CNT growth. (A) Fe-S equilibrium phase diagram and calculation of fast cooling processes superimposed. Lower line: the Scheil model with suspended eutectic transformation (lower limit). Upper line: the Scheil model and a possible second liquid phase richer in S (upper limit). (B) Phase fraction evolution showing the presence of a liquid shell down to at least 1000° C. (C) Schematic showing the growth of the CNT from an FCC core-Fe-S shell catalyst particle. S lowers the interface energy at the CNT edge and favors its nucleation, but the interface of Fe-S-C and the basal plane of graphite are nonwetting, leading to the ejection of the graphitic layer (i.e., extrusion of the CNT).

In determining the catalyst evolution throughout the CVD reaction, it is instructive to consider the corresponding equilibrium phase diagrams. Inspection of the ternary diagram at 1250°C (Fig. 6) suggests the formation of a S-rich liquid phase (L2), which is likely to be related to the evolution of the Fe-S shell. The Fe-S equilibrium phase diagram (Fig. 11A) predicts more than a 15% fraction of the liquid even in temperature as low as 1000°C, before the eutectic transformation, in agreement with the view that the catalyst is in a molten sate.

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FIG. 12 Chiral angles of SWNTs extracted from TEM measurements. (A) SWNT bundle and its corresponding SAED pattern presenting the characteristic armchair diffraction pattern. (B) Chiral angle distribution obtained from SAED patterns showing that 80% of the CNT bundles analyzed present chiral angles in the range of 20–30 degrees. For reference, the dashed line shows the distribution line shape of theoretical predictions for molten catalysts based on the model proposed by Yakobson and coworkers [19].

The core-shell path can be analyzed from thermodynamic calculations of the phase fraction of solid phases during fast quenching (cooling rate of 150°C s
1) considering 9 at.% of S in Fe [18] (Fig. 11B). Taking into account the effect of undercooling under fast cooling, Fig. 12B shows lower and upper limits for the phase fraction evolution of the catalyst by suspending/retarding the formation of FCC/Fe1
xS (pyrrhotite) eutectics and yet preserving primary FCC. The lower limit represents conventional Scheil model that assumes no back diffusion in a solid but fast diffusion in a liquid allowing an instant uniform composition distribution [20]. Because the solubility of S is negligible, the predicted solid fractions using the conventional Scheil model and under equilibrium conditions are the same, so this lower limit is thus simply the extension of the liquidus line down to a divorced eutectic point (i.e., the eutectic without lamella structure even in bulk). The upper limit of the phase transformation not only considers the conditions of the Scheil model but also allows for the possibility of a phase transformation into a second liquid richer in S (L0 in Fig. 11A) upon fast cooling, increasing the phase fraction of the primary FCC phase but still with a liquid surrounding it. Therefore, the outer layer of the catalyst particle in the high-temperature CVD process is in the liquid state at temperatures relevant for CNT growth, leading to a formation of a core (Fe)-shell (Fe1
xS) structure through a phase transformation pathway in between the two calculated limits. The core-shell catalyst evolution suggests that CNT growth is enabled by the formation of a S-rich liquid at the surface that limits the C diffusion into the “bulk” particle compared with a S-free system. There, it acts as a promoter in the catalytic decomposition of the incoming C-containing molecules from the gas phase, very much in line with its role in methanation of transition metal catalysts, where in small quantities S prevents subsurface phase diffusion but leads to the formation of layered graphitic deposits [21]. The formation of transient C-S bonds would stabilize the edge of a nascent CNT and lower the line energy between the graphite lattice edge {hk0} and the Fe-S-C molten catalyst. However, once a graphitic ring or hemisphere is formed around the particle, an interfacial area between the graphite lattice basal plane (002) and the catalyst is formed (Fig. 11C). Its corresponding interfacial energy in fact carries a large energy penalty. Molten Fe-S-C is known not to wet graphite, even at small concentrations <1 at.% S [22], so the rejection of graphitic carbon probably plays an important role in the exceptionally fast extrusion of the CNTs in the direct spinning process in the order of mm s
1.

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2.3 CHIRAL ANGLE DISTRIBUTION AND N,M INDICES ASSIGNATION The synthesis of fibers of SWNTs has enabled the study of the chiral angle of the constituent CNTs by a combination of Raman spectroscopy and selected area electron diffraction (SAED) in TEM [18]. Because of the strong tendency of SWNTs to aggregate in bundles, due to the enormous challenge to exfoliate CNTs in fibers and the large number of nanotubes in a single fiber cross section, a distribution of chiral angles statistically significant could only be obtained by SAED on bundles (20 nm average diameter), at the expense of sacrificing angular resolution. Fig. 12A shows an example of an inspected bundle and the corresponding diffraction pattern (inset). It shows the strong {110} reflections perpendicular to the (002) and a chiral angle of 30 degrees, corresponding to armchair nanotube. Analysis of over 300 CNTs in more than 20 bundles gives a distribution of chiralities that is clearly biased toward high angles and peaking at the armchair end, with a complete absence of zigzag nanotubes (Fig. 12B). Interestingly, such distribution was found to be independent of C precursor (e.g., toluene instead of butanol) or promoter (e.g., Se instead of S) [16]. Its line shape is in good agreement with the theoretical prediction for molten catalyst CNT growth [19] based on a screw dislocation growth model, in which ˚
1) comhigh-chiral angle tube predominates through the lower edge energy of armchair (γ A  0 eV A
1 ˚ ). pared with zigzag (γ Z  0.45 eV A The assignation of chiral indexes (n,m) in SWNT fibers through the radial breathing mode frequency ωRBM and the analysis of the G
line shape from resonant Raman spectroscopy is not as straightforward as in standard SWNTs [23, 24]. The large number of closed-packed nanotubes (104) under the laser beam when analyzing a fiber results in the strong overlap of several RBM peaks (sometimes almost a continuum) and a G
band line shape dominated by metallic nanotubes masking the contribution of semiconducting ones [18, 25]. To partially overcome these limitations, we have recently carried out an extensive Raman study on SWNT fibers with a low degree of bundle aggregation and thus with better resolved RBMs and G peak features. The first step toward assignation of chiral indexes (n,m) is to compare experimental RBM frequencies with the theoretical optical transitions for different CNTs, captured in the Kataura plot (Fig. 13A). Fig. 13 shows the SWNT diameter ranges expected based on HRTEM observations of nanotubes in CNT fibers, superimposed onto the Kataura plot. The laser energies used in the study are also marked. RBM frequencies are related to SWNT diameters through the dependence of transition energy on 1/ diameter (d), leading to a convenient relation w ¼ A = d + B, where A and B are empirical constants that depend strongly on SWNTs and their interaction with the environment (substrate, surfactants, bundling, …) [23]. After the analysis of a large number of spectra using various laser lines, we obtained A ¼ 214  5 cm
1nm and B ¼ 17  5 cm
1 for CNT fibers on silicon substrate. These values are in agreement with data obtained for SWNTs in aqueous suspensions wrapped by a surfactant [26,27] or alcohol-assisted CVD isolated SWNTs and bundles [28]. Obtaining experimental values for A and B is the first necessary step to index SWNTs based on Raman data. With such relation in hand, it is possible to carry out a detailed comparison of a specific region of the Kataura plot and the experimental Raman spectra, as shown in Fig. 14, for a SWNT bundle in a fiber, probed with 785 and 633 nm laser excitations. The approximate resonance window of 100 meV based on the resonance profile of the optical transitions of 60 meV [27] is also included. After careful analysis, it is possible to assign the Raman signals to different nanotube families (2n + m ¼ constant) based on the semiconducting S22 and metallic M22 transitions, for example, f27 (2n + m ¼ 27) assigned to the RBM at around 225 cm
1. Accurate determination of n,m indexes is then made considering the van der Waals effects on transition energies, such as those observed in the presence of surfactant [27], which curve the lower energy branches toward lower transition energies.

124

Metallic Semiconductor Armchair

40 35 2.33 eV (532 nm)

2.0

1.96 eV (633 nm) 1.58 eV (785 nm)

1.5

S44 S33

1.0

M11

0.0

25 20 15 10

0.5

1.0

1.5

2.0

SWNT diameter (nm)

2.5

10.1 7.9

6.7

5

S11

SWNT fiber

32.0

30

S22

0.5

(A)

Relative frequency (%)

Transition energy (eV)

2.5

SWNT fiber 183 SWNTs

37.6

5.1 0.6

0 1.0

3.0

(B)

1.5

3.0 2.0 2.5 SWNT diameter (nm)

3.5

4.0

FIG. 13 (A) Kataura plot with the most common Raman excitation laser energies and the diameter range of SWNT within the fiber. The (B) SWNT diameter distribution is obtained from HRTEM images of approximately 180 nanotubes.

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3.0

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125

FIG. 14 Kataura plot obtained by the relationship ω ¼ 214/d + 17 and its comparison with RBM peaks of a single SWNT bundle for two different Raman excitation laser lines (785 and 633 nm). Each group of peaks is assigned to different SWNT families (e.g., f25 for 2n + m ¼ 25).

In Raman spectroscopy of SWNTs, G band reflects metallic or semiconducting behavior of the nanotubes depending on the line shape of the G
component, which can present Breit-Wigner-Fano (BWF) or Lorentzian line shape, respectively. In standard isolated SWNT bundles where metallic and semiconducting tubes are present, metallic nanotubes dominate G-band line shape due to the strong enhancement of the BWF Raman line caused by the plasmon band formation in SWNT bundles [29]. The same effect is present when probing CNT fiber due to the high density of nanotubes packed in bundles resulting in a metallic behavior of the G band at the macroscopic scale. However, probing more

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FIG. 15 G peaks of SWNT bundles showing (A) the domination of the metallic G
line shape (BWF) and (B) the resolution of metallic (BWF) and semiconducting (Lorentzian) line shapes of the SWNTs in the bundle.

individualized bundles in low-density SWNT fiber samples has enabled to resolve metallic and semiconducting contributions to the G band depending on the metallic/semiconducting nanotube ratio (Fig. 15). Work is in progress to apply this method to determine the fraction of nanotubes of a different metallicity in a bulk CNT fiber.

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3 STRUCTURE-PROPERTIES: A COMPLEX HIERARCHICAL STRUCTURE 3.1 CONTROL OF ORIENTATION ALONG FIBER AXIS The orientation of individual CNTs in continuous macroscopic fibers was achieved by dispersing them in polymer solution [1] or lyotropic nematic liquid-crystal phases [2] in a process similar to polymer fiber wet spinning due to the fact that CNTs are similar to polymer chains [30]. Applying the same polymer wet spinning principles to CVD direct spinning process, the dilution in the gas phase of the very long CNTs (1 mm) reduces the entanglements in the aerogel, thus enabling the drawing of the in-form fiber and alignment of the building blocks [31]. The rate at which the aerogel is drawn is a fundamental parameter controlling the final properties of the fiber, and the dilution of the aerogel through either a decrease in precursor feed rate or an increase in hydrogen carrier gas flow rate enables much faster fiber drawing rates (Fig. 16). The maximum winding rate, Wmax, is a good indicator of how much an aerogel can be drawn during the spinning process. Independently of the type of CNTs and hence of the thiophene content (SWNTs, double-walled nanotubes (DWNTs), or multiwalled nanotubes (MWNTs)), the highest winding rates are enabled by low feed rates and/or high H2 flow rate (Fig. 17A). The role of the feed rate and the H2 flow rate is equivalent in diluting the system and leading to higher values of Wmax (Fig. 17B and C, sections of Fig. 17A at a constant feed rate or H2 flow rate). The draw ratio is obtained by defining zero draw as the minimum winding rate achievable without the accumulation of fiber in the reaction, corresponding to 3 m min
1 under synthesis conditions used here. The type of CNTs in terms of the number of layers is dictated by the precursor mixture S/C and is not altered by changes in the precursor feed rate and hydrogen gas flow that not only enables the catalytic reaction but also has a secondary role of fixing the aerogel constituents and thus the possibility of CNT assembly into the final fiber. The direct dilution effect in the aerogel is the increase of its diameter as the precursor feed rate increases (Fig. 18), so at a constant hydrogen flow, the aerogel diameter increases a 20% when the precursor feed rate is reduced from 5 to 2 mL h
1, going from a concentrated aerogel that cannot be substantially drawn (Wmax ¼ 8 m min
1) to a more diluted one that can be spun much faster (Wmax >20 m min
1) and produce oriented fibers.

3.2 STRUCTURAL STUDIES Like other macroscopic ensemble of nanobuilding blocks, CNT fiber has a complex hierarchical structure arising from the confluence of multiple length scales. Electron microscopy observation is a convenient way to inspect the structure of the fibers, and attempts have been made to describe their hierarchical structure of CNT fibers by normal electron microscopy analysis [32] and inspection of focused ion beam (FIB) sections [33,34]. As the examples of electron micrographs at different magnifications in Fig. 19 show, an additional unusual feature of CNT fibers is the coexistence of a large porosity and extended closedpacked bundle domains. Indeed, few materials are simultaneously considered promising as a highperformance fiber and as a porous electrode. Not surprisingly, the structure of CNT fibers and similar ensembles is both difficult to probe experimentally and to characterize precisely.

3.2.1 Pore structure The most natural method to study porous systems is gas adsorption. However, the need for a relatively large sample size (1–10 km of individual filament) has meant that there are very few studies reporting on the specific surface area (SSA) and pore structure of CNT fibers. Although initial reports put SSA of

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FIG. 16 (A) Schematic representation of the direct CNT fiber spinning process showing the effect of aerogel dilution on the orientability of the final fibers. SEM micrographs of (B) nonoriented and (C) oriented CNT fibers.

LC-spun SWNT fibers as high as 500 m2 g
1, more precise measurements converge to values around 100 [35] to 250 m2 g
1 [36]. The latter values are in line with theoretical predictions [37] considering the strong tendency of CNT to form bundles. Extensive BET N2 adsorption measurements in our research group indicate that artificially high values of SSA are very common when sample size is too small (below around 50 mg) or the degassing process is not adequate. With respect to measurement conditions, it is also important to ensure enough adsorption equilibration time and sufficient measurement points in an adequate range of relative pressures [38]. Under those controlled measurement conditions, it is possible to carry a pore size analysis by methods such as Barrett, Joyner, and Halenda (BJH). Fig. 20A shows an example of a N2 isotherm, accompanied by its multipoint analysis (Fig. 20B) used to ensure measurement accuracy. The resulting BJH pore size distribution, shown in Fig. 20C for CNT fibers with low degree of orientation, confirms that most pores are in the mesoscopic range, which is consistent with the relatively low values of SSA compared with activated carbons, but on the other hand implies much faster diffusion rates of small molecules. Comparison of BJH

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129

FIG. 17 Dependence of the maximum winding rate on precursor feed rate and hydrogen flow shows that faster rates are obtained in dilute conditions when the concentration of particles (CNTs) in the gas phase is lower, and therefore, the possibility to draw the aerogel and orient the CNTs is higher. (A) Surface plot of parameter space for different aerogel compositions. Constant hydrogen flow (B) and constant feed rate (C) sections of parameter space showing the equivalent role of these two parameters in diluting the aerogel.

results obtained for fiber spun at different draw ratios, which produces different degrees of CNT alignment (vide infra), shows that drawing has no substantial effect on SSA (Fig. 20D) or the distribution of pore sizes in the mesoporous range. Pore structure can be further studied by small-angle X-ray scattering (SAXS). In a recent study, we found that large samples of low-orientation CNT fibers have a structure consistent with a surface fractal with dimension of 2.5 and 2.8 for samples made of predominantly few-layer CNT and SWNTs, respectively. The dimension can be extracted from the (Porod) slope of a plot of log scattering intensity versus log q. Gas adsorption measurements give a similar dimensions when calculated using the method developed for porous carbons [39]. It is also of interest to determine more precisely the pore size and shape, particularly of anisotropic CNT fibers. SAXS corresponds to intertube separations in the range of 5–100 nm, which are ideal to inspect the opening of pores at bundle bifurcations. Yet, the wide range of interbundle separations and inherently irregular structure make a quantitative analysis of experimental data a challenge,

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Aerogel diameter (a.u.)

1.2 1.1 1.0 0.9 0.1 0.0 2.0

(A)

(B)

2.5

3.0

3.5

4.0

4.5

5.0

Precursor feed rate (mL h–1)

FIG. 18 (A) Optical images showing the aerogel expanding as feed rate is decreased. (B) Comparison of aerogel diameter from image analysis showing a 20% increment with dilution.

FIG. 19 Electron micrographs showing the hierarchical structure of CNT fibers. (A) Individual CNT fiber filament. (B) At low magnification, the sample resembles a porous network. (C) Higher-resolution SEM micrographs show the mesoporous structure arising from imperfect bundle packing, and (D) high-resolution TEM show that pores are delimited by CNTs that branch out of bundles.

3 STRUCTURE-PROPERTIES: A COMPLEX HIERARCHICAL STRUCTURE

131

1200 5.0

800

4.5

1/[Q(Po/P–1)]

600 400

2.5 2.0

0 0.0

0.2

0.4

0.6

0.8

1.0

(B)

Relative pressure (P/P0)

2.5 2.0

FWHM2 = 43.6 1.5

FWHM1 = 43.3

1.0 0.5

0.10

0.15

0.20

0.25

0.30

Relative pressure (P/Po) 2.5

300

Surface area (m2 g−1)

Draw ratio: 126 Draw ratio: 4.7 Lorentz fit (1) Lorentz fit (2)

3.0

dV (cc/Å/g)

3.5 3.0

200

(A)

4.0

250 2.0 200 1.5 150

0.0 0

(C)

20

40

60

80

Pore diameter (nm)

100

120

100

(D)

0

20

40

60

80

100

120

Cumulative volume of pores (cm3 g−1)

Volume adsorbed (cm3 g–1)

5.5 1000

1.0 140

Draw ration (m min−1)

FIG. 20 Textural properties of CNT fibers. (A) N2 adsorption-desorption isotherm, (B) linear regression on the multipoint BET plot, (C) pore size distribution calculated by the BJH method for nonaligned (draw ratio, 4.7) and aligned (draw ratio, 126) samples, and (D) dependence of specific surface area and total pore volume on the draw ratio [38].

compounded by the anisotropy of the fibers. Fig. 21A presents a simplified schematic of the main parameters extracted from SAXS data, namely, the average length of the scattering elements through the section and length of the sample, Lp and L3, respectively. Calculation of these parameters from SAXS data for oriented samples leads to a projected structure of elongated slit-shaped pores, as shown in Fig. 21B. With increasing alignment in the fibers, pores narrow and elongate, as observed in Table 1 for a set of fibers with the same composition. Yet, obtaining precise absolute values for Lp and L3 requires further considerations not taken into account in previous studies [35,39]. The first is the need to analyze individual CNT fiber filaments, since the use of multiflaments, while reducing collection times, introduces a large misorientation [39] that has an effect on the final analysis. More importantly, the analytic treatment of data needs to treat the equatorial scattering as a 3D structure and to take into account density fluctuations inherently present in the graphitic structure of the fibers [40]. Establishing a structural model for CNT fibers based on SAXS has proved equally challenging and interesting; few materials have the combination of irregular hierarchical pore structure and anisotropy of these nanocarbon aggregates. The combined use of wide- or small-angle X-ray scattering (WAXS/SAXS) and gas adsorption has emerged as an attractive method to characterize more precisely the complex structure of CNT fibers across different length scales. Further work should be directed at studying samples with controlled composition and orientation that thus enable deconvolution of these interlinked effects on SSA. This

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CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

Lp

Bf

Fiber axis

CNT bundle L3

(A) SAXS

L3 = 221 nm

WAXS

Bf = 17.9°

Lp = 13.1 nm

(B) FIG. 21 Fiber structure according to SAXS results. (A) Parameters determined from 2D SAXS and their assignation. (B) Schematic showing the resulting fiber structure for a highly drawn sample (DR  126).

Table 1 Structure Properties of CNT Fibers Produced at Different Draw Ratios SAXS Draw Ratio

FWHM (Degrees)

Lp (nm)

L3 (nm)

4.7 31.5 63.0 95.0 126.0

52 31 25 25 26

20.3 17.6 17.3 13.8 13.1

159 164 210 227 221

will clarify the role of features such as number of CNT layers and bundle size while also helping refine current structural models built based on these techniques.

3.2.2 Orientation The dominant structural feature to determine most bulk physical properties of CNT fibers and similar ensembles is the orientation of CNTs. This can be measured mainly by three methods: Fourier transform of electron micrographs, polarized Raman spectroscopy, and X-ray diffraction by 2D WAXS/ SAXS. They provide a distribution of signal intensity relative to an azimuthal angle, typically chosen to reflect the “real” distribution of elements relative to the fiber main axis. The full width at half maximum (FWHM) of the associated azimuthal distributions, an indicator of misalignment, is often similar when using these techniques. But in spite of this apparent agreement, a quantitative structure-property prediction based on orientation measurements requires additional considerations in terms of the significance and analytic interpretation of the data. In the first case, the question is

3 STRUCTURE-PROPERTIES: A COMPLEX HIERARCHICAL STRUCTURE

133

how representative local measurements are of bulk features. Electron micrographs of bundles inherently probe small areas, and thus, one has to confirm that the sample under study is uniform by collecting extensive maps and analyzing magnification effects [41]. Raman measurements, on the other hand, have a resolution limit of around 1 μm, depending on resonant conditions and measurement parameters, although with the benefit of being able to use larger laser spot sizes to probe large sample areas. A limitation of both SEM and Raman signals is that they originate from the sample surface and may thus not be representative of the whole sample volume. In this respect, X-ray studies suggest that there is no substantial core-sheath structure in CNT fibers produced by either direct spinning [42] or coagulation [43], but a more direct confirmation is desired. Two-dimensional X-ray scattering methods are attractive, as they probe the whole sample volume through the beam path, and in the case of WAXS, the data have a direct correspondence with the crystal structure analyzed. However, the intrinsically low scattering of C implies that measurements have often to be conducted at synchrotron facilities. Even then, WAXS measurements are only possible with a microfocus beam (around 10 μm2), since multifilament samples introduce a large misalignment that hinders further study of individual fiber filament orientation (Fig. 22C). In this respect, SAXS measurements have emerged as a key method to probe orientation. Although SAXS signal arises from larger elements than individual CNTs (e.g., bundles and pores), studies have confirmed the correspondence between WAXS and SAXS orientation [42] (Fig. 22A, D, and E). Having identified the limitations of the methods used to determine alignment in CNT fibers, the questions move onto the parameters extracted from each technique. The simplest method, often used in the process of optimizing longitudinal fiber properties, is to compare relative values of degree of orientation, for example, as the FWHM of the azimuthal distribution of intensities. A more meaningful alignment metric is Herman’s orientation factor, which takes values of 1 for perfectly aligned fibers and 0 for random orientation [44]. It corresponds to the first coefficient of the first Legendre polynomial of the orientation distribution function in spherical coordinates, which is quadratic, hence why this term is commonly known as P2. P2 is obtained from Eq. (1): hP2 ð cos θÞi ¼

  1
3 cos 2 θ
1 2

(1)

Here, hcos2θi takes the form of a weighted average of the orientation distribution function (ODF). For SAXS data, it can be determined by Eq. (2): h cos n θi ¼

ðπ

cos n θIðθÞ sin θ dθ

(2)

0

with the ODF of the normal to the basal planes defined by Eq. (3): ΨðθÞ ¼ ð π

I ðθ Þ

(3)

I ðθÞsin θ dθ

0

Interestingly, WAXS/SAXS X-ray measurements on direct-spun CNT fibers consistently show a Lorentzian azimuthal distribution. Irrespective of azimuthal breath, such distribution leads to low values of P2, around 0.6. The implication is that even in samples with highly CNTs, the presence

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CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

FIG. 22 (A) WAXS and SAXS azimuthal distributions and (B) 2D X-ray scattering pattern of aligned CNT fiber covering both SAXS and WAXS ranges. (C) WAXS azimuthal FWHM evolution with draw ratio for fibers with different number of filaments. (D) and (E) WAXS and SAXS azimuthal FWMH mappings.

of misaligned elements indicated by the “tail” in the Lorentzian distribution leads to P2 values far lower than those found in polymer high-performance fibers, typically >0.9. Polarized Raman spectroscopy measurements under sample rotation can provide the prefactors of the first two Legendre polynomials (P2 and P4 [45]). The example in Fig. 23 shows the drop in G-band intensity as a function of angle between the fiber axis and the polarization direction of the incident beam (parallel to the analyzer, i.e., IVV). Using the maximum entropy formalism, P2 and P4 can be used to obtain the ODF. Ferna´ndez et al. have recently compared orientation measurements by polarized Raman and SAXS for a range of CNT fibers and discussed the resulting ODF and related parameters in the context of tensile property prediction [46]. Improvements in alignment in polymer fibers have the additional effect of increasing the size of crystalline domains, with the corresponding improvement in tensile properties. Heat tensioning of rigid-rod polymer fibers, for example, has the dual effect of improving orientation and increasing lateral coherence size, thus doubling both tensile strength and modulus [47]. In the case of CNT fibers, determining a coherent domain size from WAXS has been a challenge because of the contribution from internal layers of CNTs and the wide distribution of intertube spacings in the (002) region and the overlap between other graphite reflections with residual catalyst peaks. WAXS measurements on multifilament samples confirm that higher alignment increases the fraction of CNTs at turbostratic separation

3 STRUCTURE-PROPERTIES: A COMPLEX HIERARCHICAL STRUCTURE

Raman (lw) SAXS

1.0 Normalized intensity (a.u.)

135

0.8

0.6

0.4

0.2

0.0 0

20

40

60

80

100

Angle (degrees)

FIG. 23 Experimental data obtained to determine CNT alignment. Azimuthal profile from SAXS and polarized Raman spectroscopy G-band intensity as a function of the angle between the fiber and the polarization direction, in Ivv mode (see text for details).

FIG. 24 Effects of increased in fraction of graphitic domains. (A) Radial profile from WAXS in the range q ¼ 15–22 nm
1 for fibers produced at different draw ratios, with scattering intensity normalized by the invariant. (B) More oriented fiber shows an increase in the area of the (002) reflection.

(Fig. 24) [39], that is, a sufficiently close separation to take part in stress transfer by shear. A “degree of crystallinity” can be calculated from the integral of the (002) reflection normalized by the invariant (Q). Comparing this quantity (I(002)/Q) for samples produced at different draw ratios confirms that higher alignment leads to better packing of elements in the fibers and a larger fraction of graphitic layers in close proximity.

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CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

The results in Fig. 24 provide a method to compare similar fibers, for example, in the process of improving longitudinal properties by increased alignment. But there is also interest in determining the actual size of coherent lengths between adjacent CNTs in fibers. Defining precisely the nature of these domains and establishing a method to characterize them remains a critical challenge for further development of fiber properties.

3.3 STRUCTURE AND TENSILE PROPERTIES The development of theoretical models able to successfully describe the physical properties of CNT fibers as a function of their structure has proved an elusive challenge. Most efforts on establishing structure-property relations have focused on tensile properties. A fundamental difficulty arises because of the inherently complex hierarchical structure of CNT fibers. Such complexity stems from the confluence of many parameters determining bulk properties, including those linked with the physical and chemical properties of constituents (number of layers in CNTs, chiral angle, diameter, the presence of impurities, etc.), their spatial arrangement (orientation and bundle formation), and interaction between building blocks. CNT fibers have similar features to both staple yarns and high-performance polymer fibers. Experimental results have confirmed expectations about higher alignment [31,48,49] and longer CNT length (Fig. 25) leading to higher tensile properties. Fig. 25A shows literature data for the strength of CNT fibers produced by different methods as a function of CNT length, illustrating this effect. For comparison, we show data for cotton yarns, whose tensile properties are strongly determined by the effective length of the microfibers (Fig. 25B). Similarly, large-diameter few-layer CNTs are recognized as the optimal type to maximize intertube contact per unit weight [55]. Although length and orientation were early recognized as key parameters affecting fiber tensile properties, going a step further and determining the dependence of bulk tensile properties on finer structural features, such as the number of layers of the constituent CNT, have proved challenging even from a theoretical point of view. The problem is compounded by the fact that often the key properties are those of the bundles, and bundle size is difficult to determine and more so to control. Nevertheless, there are emerging mechanical models going a step further in establishing quantitative structure-property relations. This has been possible by analyzing samples with controlled structure that thus enable testing the effect of individual structural parameters independently. A recent study by Tsentalovich et al. has looked at the dependence of tensile strength on aspect ratio, CNT type, and purity, in fiber samples produced by wet spinning from liquid crystalline lyotropic nematic phases [56]. The use of this method ensures a very high, nominally identical orientation in all samples. Nominal tensile strength is found to scale almost linearly (0.9) with aspect ratio, which is in agreement with simplified shear-lag models assuming axial stress to build up linearly with overlapping length, for CNT lengths below the critical length for fracture and thus failing in shear. This work has also shown similar properties in fibers produced from SWNTs or DWNTs. A parallel study has looked at the role of orientation for samples with fixed composition and demonstrated that the fibers can be treated as a network of crystallites, similar to the structure of highperformance polymer fibers, with their tensile properties determined by the crystallite: ODF, shear modulus, and shear strength [46]. Synchrotron SAXS measurements on individual CNT fibers were used to determine the initial ODF and its evolution during in situ tensile testing. This enables the prediction of fiber tensile modulus, strength, and fracture envelope with good agreement with

3 STRUCTURE-PROPERTIES: A COMPLEX HIERARCHICAL STRUCTURE

137

FIG. 25 (A) Fiber tensile strength dependence with CNT length obtained from different works: dark square [2], dark downpointing triangle [50], dark circle [51], dark left-pointing triangle [52], dark right-pointing triangle [53], and dark triangle [48]. (B) Fiber tensile strength dependence with effective length of microfibers [54].

experimental data. Fig. 26 shows that compliance scales linearly with the degree of orientation (hcos2Φ0i), at least for highly aligned fibers. This is in agreement with the model, which predicts E
1 by Eq. (4):
 cos 2 θ0 1 1 ¼ + g E ec

(4)

where E is the fiber modulus and ec and g the in-plane and shear moduli of graphite, respectively. Extrapolation to perfect orientation gives ec ¼ 540 GPa, which is not far from the modulus of single crystal graphite (around 1000 GPa). The fibrillar crystallite model helps decouple orientational from compositional effects. Thus, the comparison of fibers made up of collapsed CNTs and of few-layer mutiwalled indicates that the higher

138

CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

FIG. 26 Compliance (E
1) of CNT fiber plotted against the orientation parameter hcos2Φ0i for both in-house fibers (dark square and dark circle [46]) and data obtained from the literature (dark hexagonal [57]; dark triangle [58]; dark down-pointing triangle, dark dice, and dark star [59]; and dark pentagonal [32]).

strength of the former is associated to a higher crystallite shear strength and shear modulus. This is attributed to the higher tube-tube contact area for collapsed CNTs. It is also of interest to determine the reorientation of the crystallites upon axial deformation, as this gives a measure of the crystal shear modulus. This has been achieved by the determination of the ODF from in situ SAXS during the stretching of a single 10 μm diameter filament. The resulting shear modulus in the elastic deformation regime and the subsequent secant shear modulus in the plastic regime can provide useful predictions of fracture envelope and insight into the associated deformation mechanisms. They also point to the unusually low ratio between elastic and “plastic” shear modulus in these fibers, which is thought to be responsible for their exceptionally high energy to break (50–100 J g
1), resulting in a very high-ballistic figure of merit above that of Kevlar [46]. We anticipate that the fibrillar crystal model can be a guiding formalism to better describe the mechanical properties of CNT fibers, particularly with the view to tune interfaces to favor specific failure mechanisms and thus obtain desired properties. Furthermore, it is likely that transport properties of highly aligned fibers can be similarly described by reducing the system to an ensemble of oriented crystalline domains, whose properties are mostly defined by the ODF and the coherent length for charge/heat transfer.

4 LOW-DIMENSIONAL PROPERTIES In most envisaged applications of CNT fibers developed so far, they are essentially exploited as tough porous conductors, sometimes combined with ion intercalation or a catalytic process. There has been comparatively much less work on measuring and exploiting the low-dimensional properties of the constituent CNTs in bulk samples. Here, low dimensional refers to optoelectronic properties arising from

4 LOW-DIMENSIONAL PROPERTIES

139

the quantized one-dimensional electronic structure of few-layer CNTs, in contrast to that of graphite or multilayer graphene. The discussion earlier shows that there is indeed a high degree of control of CNT composition in fibers produced by the direct spinning method, including the possibility to produce predominantly SWNTs. A further prerequisite to observe low-dimensional properties in these systems is a high degree of crystallinity in the CNT hexagonal lattice. In the context of low-dimensional properties, crystallinity does not refer to graphitization, which can be interpreted as Bernal stacking, but rather a high degree of conjugation in terms of sp2 bonding, measured, for example, by XPS and often including the emergence of the π-π* plasmonic band (see examples in [60, 61]). A simple direct method to access the low-dimensional electronic structure of large CNT fiber samples is through their quantum (chemical) capacitance measured by standard electrochemical techniques [36]. Fig. 27 shows examples of cyclic voltamograms (CV) of CNT fiber electrodes in different electrolytes. In all three cases, the curves show a “butterfly” shape, notably different from the rectangular shape of an ideal metallic electrode or conventional activated carbon. The saddle-point shape of the curve indicates a dependence of differential capacitance (I/V) on electrochemical potential. While CV helps to rule out pseudocapacitive redox processes, which also produce an irregular CV line shape, it is more instructive to measure capacitance at different potentials using electrochemical

1.5

Pyr14TFSI

1M TBAPF6

3

1.0 2 I (A/g)

I (A/g)

0.5 0.0

1 0

–0.5

–1

–1.0

–2 –3

–1.5 –1.5

(A)

–1.0

–0.5

0.0

0.5

1.0

1.5

–1.0

–0.5

0.0

(B)

E (V) 1.0

0.5

1.0

1.5

E (V) 1M KOH

I (A/g)

0.5 0.0 –0.5 –1.0 –1.5 –1.2 –1.0 –0.8 –0.6 –0.4 –0.2

(C)

0.0

0.2

E (V)

FIG. 27 Cyclic voltamograms of CNT fibers in different electrolytes as (A) Pyr14TFSI, (B) TBAPF6, and (C) KOH, showing the typical line shape arising from the contribution from the quantum capacitance of low-dimensional CNTs.

140

CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

impedance spectroscopy (EIS) measurements. Fig. 28A presents a plot of total capacitance normalized by electrode surface, against electrochemical potential. It also includes data for electrode conductance, determined separately from EIS by two-probe resistance measurements. Both total capacitance and relative conductance increase symmetrically from the point of zero charge over the wide electrochemical range enabled by the highly stable ionic liquid used (around 3 V). The minimum of differential capacitance of 3.2 μF cm
2 at pzc is in excellent agreement with values for highly oriented pyrolytic graphite (HOPG) and high-quality CVD graphene samples [63]. Assuming a linear dispersion relation in the range (0  1 V), the left and right slopes give slopes of 5 μF cm
2 V
1, comparable with few-layer high-quality graphene with aqueous electrolyte (2.5–7 μF cm
2 V
1 [63]). The dependence of CV on electrochemical potential can be understood considering that it represents the series combination of two capacitances therefore dominated by the smaller one. One is purely electrostatic and dependent on surface area, with a constant value close to 3.2 μF cm
2 for a graphitic carbon. The other is the quantum (chemical) capacitance, accessible only materials with a low density of states (DOS) near the Fermi level, such as low-dimensional semiconducting or quasimetallic nanocarbons. (For a rigorous treatments of quantum capacitance in semiconductors and nanocarbons, see [64, 65].) 6 5 3

7 6

2

5 4

s /s PZC

CA (mF cm–2)

8

JDOS (a.u.)

9

1

3

4 3 2 1

2 1

0 –1.0

(A)

JDOS of 16 SWNTs from families 2n+m={25,33,36}

4

10

–0.5

0.0

0.5

0 –2.0 –1.5 –1.0 –0.5

1.0

E vs pzc (V)

(B)

0.0

0.5

1.0

1.5

2.0

Energy (eV)

Normalized quantum capacitance (Cq/a) (a.u.)

1.0 0.8 0.6 0.4 0.2 0.0 –1.0

(C)

–0.5

0.0

0.5

1.0

Chemical potential (eV)

FIG. 28 Quantum capacitance in CNT fiber electrodes and its relation to the joint density of states (JDOS) of SWNTs. (A) Experimental area-normalized capacitance and conductance against electrochemical potential. (B) JODS of 16 SWNTs (semiconducting and metallic), obtained by the addition of individual DOS [62]. (C) Theoretical quantum capacitance of a SWNT bundle assuming a linear relation of the JDOS arising from the smoothing of features in individual DOS after their superposition.

4 LOW-DIMENSIONAL PROPERTIES

141

The close correspondence between the capacitance “V-shape” plot (Fig. 28A) and electronic structure of the CNT fiber electrode is more evident by plotting a theoretical joint density of states (JDOS) of the constituent CNTs. We take JDOS as the sum of theoretical individual DOS of SWNTs, a valid assumption for bundles with CNTs of different chiral angles or in the absence of crystallographic registry [66], and consider diameters from 1 to 3 nm considering equal probability of chiral angles. Individual DOS were retrieved from S. Mayurama work [62]. The resulting graph (Fig. 28B) is indeed very similar to the experimental result, especially considering that it was produced without considering interactions between CNTs and the electrolyte and ignoring thermal broadening effects. The correspondence is captured by a couple of simple equations. Quantum capacitance (per unit length or per unit area) is given by Eq. (5): Cq ¼


∂q ∂q ¼
e ∂φ ∂μ

(5)

where q is the charge density, φ the electrostatic potential, and μ the chemical potential. The charge density is given by the integral of the occupied DOS, which in the case of CNTs (and graphene) must consider electron-hole symmetry, leading to Eq. (6): Cq ¼ e2

ð∞
∞

gðEÞFth ðE, μÞdE

(6)

where g(E) is the DOS and Fth a thermal broadening function defined as the Eq. (7): Fth ¼

  ∂F 1 E
μ sech 2 ¼ ∂μ 4kB T 2kB T

(7)

There are analytic expressions for the individual DOS of SWNTs, but solving Eq. (6) for the large number of SWNT types expected in a CNT fiber is not practical. A less rigorous but convenient approach is to note that the JDOS of superimposed SWNTs is approximately linear, g(E)  aj Ej, coincidentally similar to that of graphene. This assumption reduces Eqs. (6)–(8): h  i Cq ¼ ae2 2kB Tln 1 + eμ=kB T
μ

(8)

As shown in Fig. 28C, Eq. (8) has a V-shaped profile, qualitatively similar to that observed experimentally. In a recent example demonstrating the sensitivity of Cq to CNT chemistry, Iglesias et al. carried out CV and EIS electrochemical measurements on CNT fiber samples subjected to a gas-phase oxidative treatment using ozone [61]. Raman spectroscopy and XPS provided direct evidence of the introduction of functional groups for different treatment times, evidenced as defects increasing the D band (Fig. 29A) and the O1s/C1s ratio (Fig. 29B) due to the formation of oxygen functional groups. Fig. 29C presents the corresponding CVs. Clearly, their line shape gradually becomes more rectangular for the samples subjected to longer oxidative functionalization times. Similarly, a CV plot from EIS shows a nearly flat profile with respect to electrochemical potential. These observations confirm that quantum capacitance in nanocarbon systems is only accessible when they not only are nanoscopic but also have a high crystallinity in terms of sp2 conjugation, leading to a low DOS near the Fermi level. The presence of functional groups is expected to introduce new localized energy states, as observed in simulations on individual graphene layers (Fig. 29D) [67] and suggested by the small peak in CV near 0 V.

CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

Normalized intensity (a.u.)

1.5

120 min 30 min 15 min 5 min Pristine fibers

3

p-CNT fibers 5 min 15 min 30 min 120 min

2.0

2 1

Increase of the valley intensity

I (A/g)

142

1.0

0 –1

0.5

–2 0.0 1200

(A)

–2.0 –1.5 –1.0 –0.5 0.0

1400 1600 Raman shift (cm–1)

(C) 25

120 min ozone treated

0.0 AreaO1s/C1s=0.29

2.0

Pristine fibers

15

10

30 min ozone treated

0.0

5

AreaO1s/C1s=0.26

0 –1.0

0.4

15 min ozone treated

–0.5

(D)

0.0

0.5

1.0

E (V)

0.0

Pristin graphene

–C6H5

–C6H4NH2

–C6H4NO2

–NH2

1.5 0.8

AreaO1s/C1s=0.24

0.4

5 min ozone treated

0.0 0.8

AreaO1s/C1s=0.09

DOS(1022/eV .g)

Norm. intensity (a.u.)

0.4

0.8

1.5

20

0.4

0.8

1.0

Functionalized fibers (120 min)

AreaO1s/C1s=0.40 Capacitance (F/g)

0.8

0.5

E (V)

1

0.5

0.4 Pristine

0.0 700

(B)

600

500

400

300

Binding energy (eV)

0

200

–1

(E)

–0.5

0

0.5

1

E (eV)

FIG. 29 Evidence of O3 oxidative treatments removing the contribution from Cq to CNT fiber electrodes. (A) Raman spectra and (B) XPS spectra showing the introduction of functional groups (defects) as a function of treatment time. (C) CV showing gradual loss of the quantum capacitance contribution leading to a more rectangular line shape. (D) EIS capacitance against voltage showing a flat profile after extensive functionalization. (E) Theoretical DOS of a graphene layers with different functional groups showing the emergence of energy states near Ef [67].

5 CONCLUSIONS AND OUTLOOK Advances in producing CNT fibers using the direct spinning process have enabled control over CNT number of layers and exposed a chiral angle distribution biased toward large angles (i.e., armchair). Such distribution is consistent with the screw dislocation growth model assuming a molten catalyst. Thermodynamic simulation of catalyst quenching and analysis of ternary Fe-S-C and binary (Fe-C and Fe-S) equilibrium phase diagrams suggest that the catalyst is molten during the CVD reaction and evolves into an Fe-rich core and a S-rich shell. Such structure is confirmed by postsynthesis

REFERENCES

143

analysis of catalysts. CNT fiber spinning is demonstrated using other group 16 elements (Se and Te), with a similar increase in CNT layers with increasing ratio of promotor to C in the reaction. The promotor has three roles: restrict C diffusion to the catalyst particle surface, stabilize the nascent CNT by minimizing the line energy of the CNT edge, and contributing to CNT extrusion through a high-surface energy with the CNT basal plane (002). CNT fibers have a complex hierarchical structure and an intrinsic confluence of length scales, ranging from the molecular identity of the building blocks in terms of CNT chiral angle and number of layers to the continuum macroscopic structure (kms). The distribution of chiral angles can be obtained from extensive measurement of electron diffraction patters in TEM. Raman spectroscopy measurements provide additional data in terms of metallicity, with various spectral features observable in the G-band region under high-spectral-resolution measurements. These features are accessible in CNT fibers purposely assembled with a low degree of bundle aggregation. Implicit in these results is the observation that CNT and bundle aggregation smooth out spectroscopic features and suggest an overlap of electronic structure and phonon modes. This issue deserves further study, particularly because of the importance of CNT-CNT interactions on bulk fiber properties. Analysis of the pore structure of CNT fibers by a combination of gas adsorption and SAXS/WAXS indicates that the material can be best described as a porous network of long crystalline domains, giving rise of a high-specific surface area. An interesting question in this respect is CNT number of layers, which on the one hand affects individual SSA but on the other influences aggregation and packing into bundles and is thus not easily related to bulk fiber SSA. Bulk fiber properties are largely dominated by the degree of CNT orientation. There has been a recent move from simple comparison of angular breath of azimuthal distributions to determination of the orientation distribution function. This is key to establish accurate structure-property correlation with predictive capability. The most advanced example is the use of a fibrillar crystal model originally developed for high performance, to describe tensile properties of CNT fibers based on their ODF and its evolution during axial deformation. It is anticipated that the interest in the electronic properties of CNT fibers and its relation to electrochemical methods will increase in the coming years. This stems not only from the interest in applications directly linked to low-dimensional CNT properties, such as sensing, but also the possibility to relate macroscopic fiber properties to more fundamental charge transfer and accumulation processes in the constituent CNTs.

REFERENCES [1] B. Vigolo, A. P
enicaud, C. Coulon, C. Sauder, R. Pailler, C. Journet, P. Bernier, P. Poulin, Macroscopic fibers and ribbons of oriented carbon nanotubes, Science 290 (2000) 1331–1334, https://doi,org/10.1126/ science.290.5495.1331. [2] L.M. Ericson, H. Fan, H. Peng, V.A. Davis, W. Zhou, J. Sulpizio, Y. Wang, R. Booker, J. Vavro, C. Guthy, A. N.G. Parra-Vasquez, M.J. Kim, S. Ramesh, R.K. Saini, C. Kittrell, G. Lavin, H. Schmidt, W.W. Adams, W. E. Billups, M. Pasquali, W.-F. Hwang, R.H. Hauge, J.E. Fischer, R.E. Smalley, Macroscopic, neat, single-walled carbon nanotube fibers, Science 305 (2004) 1447–1450, https://doi,org/10.1126/science. 1101398. [3] M. Zhang, K.R. Atkinson, R.H. Baughman, Multifunctional carbon nanotube yarns by downsizing an ancient technology, Science 306 (2004) 1358–1361, https://doi.org/10.1126/science.1104276.

144

CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

[4] Y.-L. Li, I.A. Kinloch, A.H. Windle, Direct spinning of carbon nanotube fibers from chemical vapor deposition synthesis, Science 304 (2004) 276–278, https://doi.org/10.1126/science.1094982. [5] V. Reguero, B. Alema´n, B. Mas, J.J. Vilatela, Controlling carbon nanotube type in macroscopic fibers synthesized by the direct spinning process, Chem. Mater. 26 (2014) 3550–3557. [6] M.W. Schauer, M.A. White, Tailoring industrial scale CNT production to specialty markets, MRS Online Proc. Libr. Arch. 1752 (2015) 103–109, https://doi.org/10.1557/opl.2015.90. [7] J.J. Vilatela, R. Marcilla, Tough electrodes: carbon nanotube fibers as the ultimate current collectors/active material for energy management devices, Chem. Mater. 27 (2015) 6901–6917. [8] J.J. Vilatela, Nanocarbon-based composites, in: Nanocarbon-Inorganic Hybrids: Next Generation Composites for Sustainable Energy Applications, De Gruyter, Berlin/Boston, MA, 2014. https://doi.org/10.1515/ 9783110269864 (Chapter 8). [9] M.S. Motta, A. Moisala, I.A. Kinloch, A.H. Windle, The role of sulphur in the synthesis of carbon nanotubes by chemical vapour deposition at high temperatures, J. Nanosci. Nanotechnol. 8 (2008) 2442–2449, https:// doi.org/10.1166/jnn.2008.500. [10] N.G. Chopra, L.X. Benedict, V.H. Crespi, M.L. Cohen, S.G. Louie, A. Zettl, Fully collapsed carbon nanotubes, Nature 377 (1995) 135–138, https://doi.org/10.1038/377135a0. [11] M. Motta, A. Moisala, I.A. Kinloch, A.H. Windle, High performance fibres from ‘dog bone’ carbon nanotubes, Adv. Mater. 19 (2007) 3721–3726, https://doi.org/10.1002/adma.200700516. [12] M.S. Dresselhaus, G. Dresselhaus, R. Saito, A. Jorio, Raman spectroscopy of carbon nanotubes, Phys. Rep. 409 (2005) 47–99, https://doi.org/10.1016/j.physrep.2004.10.006. [13] A.C. Ferrari, J.C. Meyer, V. Scardaci, C. Casiraghi, M. Lazzeri, F. Mauri, S. Piscanec, D. Jiang, K. S. Novoselov, S. Roth, A.K. Geim, Raman spectrum of graphene and graphene layers, Phys. Rev. Lett. 97 (2006) 187401, https://doi.org/10.1103/PhysRevLett.97.187401. [14] F. Ding, A. Ros
en, E.E.B. Campbell, L.K.L. Falk, K. Bolton, Graphitic encapsulation of catalyst particles in carbon nanotube production, J. Phys. Chem. B 110 (2006) 7666–7670, https://doi.org/10.1021/jp055485y. [15] A. Moisala, A.G. Nasibulin, D.P. Brown, H. Jiang, L. Khriachtchev, E.I. Kauppinen, Single-walled carbon nanotube synthesis using ferrocene and iron pentacarbonyl in a laminar flow reactor, Chem. Eng. Sci. 61 (2006) 4393–4402, https://doi.org/10.1016/j.ces.2006.02.020. [16] B. Mas, B. Alema´n, I. Dopico, I. Martin-Bragado, T. Naranjo, E.M. P
erez, J.J. Vilatela, Group 16 elements control the synthesis of continuous fibers of carbon nanotubes, Carbon 101 (2016) 458–464, https://doi.org/ 10.1016/j.carbon.2016.02.005. [17] B. Alema´n, R. Ranchal, V. Reguero, B. Mas, J. Vilatela, Carbon nanotube fibers with martensite and austenite Fe residual catalyst: room temperature ferromagnetism and implications for CVD growth, J. Mater. Chem. C (2017)https://doi.org/10.1039/C7TC01199G. [18] B. Alema´n, M.M. Bernal, B. Mas, E.M. P
erez, V. Reguero, G. Xu, Y. Cui, J.J. Vilatela, Inherent predominance of high chiral angle metallic carbon nanotubes in continuous fibers grown from a molten catalyst, Nanoscale 8 (2016) 4236–4244. [19] V.I. Artyukhov, E.S. Penev, B.I. Yakobson, Why nanotubes grow chiral, Nat. Commun. 5 (2014) 4892, https://doi.org/10.1038/ncomms5892. [20] S.-L. Chen, Y. Yang, S.-W. Chen, X.-G. Lu, Y.A. Chang, Solidification simulation using Scheil model in multicomponent systems, J. Phase Equilib. Diffus. 30 (2009) 429, https://doi.org/10.1007/s11669-0099568-0. [21] B. Delmon, Catalyst Deactivation, International Symposium Proceedings, 1980, (1980). [22] C. Wu, V. Sahajwalla, Influence of melt carbon and sulfur on the wetting of solid graphite by Fe-C-S melts, Metall. Mater. Trans. B Process Metall. Mater. Process. Sci. 29 (1998) 471–477, https://doi.org/10.1007/ s11663-998-0126-7. [23] P.T. Araujo, P.B.C. Pesce, M.S. Dresselhaus, K. Sato, R. Saito, A. Jorio, Resonance Raman spectroscopy of the radial breathing modes in carbon nanotubes, Physica E Low Dimens. Syst. Nanostruct. 42 (2010) 1251–1261, https://doi.org/10.1016/j.physe.2010.01.015.

REFERENCES

145

€ u, B.B. Goldberg, M. [24] A. Jorio, A.G. Souza Filho, G. Dresselhaus, M.S. Dresselhaus, A.K. Swan, M.S. Unl€ A. Pimenta, J.H. Hafner, C.M. Lieber, R. Saito, G-band resonant Raman study of 62 isolated single-wall carbon nanotubes, Phys. Rev. B 65 (2002) 155412, https://doi.org/10.1103/PhysRevB.65.155412. [25] R.M. Sundaram, K.K.K. Koziol, A.H. Windle, Continuous direct spinning of fibers of single-walled carbon nanotubes with metallic chirality, Adv. Mater. 23 (2011) 5064–5068, https://doi.org/10.1002/ adma.201102754. [26] C. Fantini, A. Jorio, M. Souza, M.S. Strano, M.S. Dresselhaus, M.A. Pimenta, Optical transition energies for carbon nanotubes from resonant Raman spectroscopy: environment and temperature effects, Phys. Rev. Lett. 93 (2004) 147406, https://doi.org/10.1103/PhysRevLett.93.147406. [27] J. Maultzsch, H. Telg, S. Reich, C. Thomsen, Radial breathing mode of single-walled carbon nanotubes: optical transition energies and chiral-index assignment, Phys. Rev. B 72 (2005) 205438, https://doi.org/ 10.1103/PhysRevB.72.205438. [28] P.T. Araujo, S.K. Doorn, S. Kilina, S. Tretiak, E. Einarsson, S. Maruyama, H. Chacham, M. A. Pimenta, A. Jorio, Third and fourth optical transitions in semiconducting carbon nanotubes, Phys. Rev. Lett. 98 (2007) 067401, https://doi.org/10.1103/PhysRevLett.98.067401. [29] C. Jiang, K. Kempa, J. Zhao, U. Schlecht, U. Kolb, T. Basch
e, M. Burghard, A. Mews, Strong enhancement of the Breit-Wigner-Fano Raman line in carbon nanotube bundles caused by plasmon band formation, Phys. Rev. B 66 (2002) 161404, https://doi.org/10.1103/PhysRevB.66.161404. [30] M.J. Green, N. Behabtu, M. Pasquali, W.W. Adams, Nanotubes as polymers, Polymer 50 (2009) 4979–4997, https://doi.org/10.1016/j.polymer.2009.07.044. [31] B. Alema´n, V. Reguero, B. Mas, J.J. Vilatela, Strong carbon nanotube fibers by drawing inspiration from polymer fiber spinning, ACS Nano 9 (2015) 7392–7398. [32] A.M. Beese, X. Wei, S. Sarkar, R. Ramachandramoorthy, M.R. Roenbeck, A. Moravsky, M. Ford, F. Yavari, D. T. Keane, R.O. Loutfy, S.T. Nguyen, H.D. Espinosa, Key factors limiting carbon nanotube yarn strength: exploring processing-structure-property relationships, ACS Nano 8 (2014) 11454–11466, https://doi.org/ 10.1021/nn5045504. [33] A.S. Wu, X. Nie, M.C. Hudspeth, W.W. Chen, T.-W. Chou, D.S. Lashmore, M.W. Schauer, E. Towle, J. Rioux, Carbon nanotube fibers as torsion sensors, Appl. Phys. Lett. 100 (2012) 201908, https://doi.org/10.1063/1.4719058. [34] R.J. Mora, J.J. Vilatela, A.H. Windle, Properties of composites of carbon nanotube fibres, Compos. Sci. Technol. 69 (2009) 1558–1563, https://doi.org/10.1016/j.compscitech.2008.11.038. [35] J. Qiu, J. Terrones, J.J. Vilatela, M.E. Vickers, J.A. Elliott, A.H. Windle, Liquid infiltration into carbon nanotube fibers: effect on structure and electrical properties, ACS Nano 7 (2013) 8412–8422, https://doi.org/ 10.1021/nn401337m. [36] E. Senokos, V. Reguero, J. Palma, J.J. Vilatela, R. Marcilla, Macroscopic fibres of CNTs as electrodes for multifunctional electric double layer capacitors: from quantum capacitance to device performance, Nanoscale 8 (2016) 3620–3628, https://doi.org/10.1039/C5NR07697H. [37] A. Peigney, C. Laurent, E. Flahaut, R.R. Bacsa, A. Rousset, Specific surface area of carbon nanotubes and bundles of carbon nanotubes, Carbon 39 (2001) 507–514, https://doi.org/10.1016/S0008-6223(00)00155-X. [38] E. Senokos, Carbon Nanotube Fiber Supercapacitors, UPM-IMDEA Materials Institute, Madrid, Spain, 2018. [39] H. Yue, V. Reguero, E. Senokos, A. Monreal-Bernal, B. Mas, J.P. Ferna´ndez-Bla´zquez, R. Marcilla, J. J. Vilatela, Fractal carbon nanotube fibers with mesoporous crystalline structure, Carbon 122 (2017) 47–53, https://doi.org/10.1016/j.carbon.2017.06.032. [40] C. Santos, E. Senokos, J.C. Ferna´ndez-Toribio, A. Ridruejo, R. Marcilla, J.J. Vilatela, Pore structure and electrochemical properties of CNT-based electrodes studied by in-situ small/wide angle X-ray scattering, J. Mater. Chem. A., submitted for publication. [41] E. Brandley, E.S. Greenhalgh, M.S.P. Shaffer, Q. Li, Mapping carbon nanotube orientation by fast Fourier transform of scanning electron micrographs, Carbon 137 (2018) 78–87, https://doi.org/10.1016/j.carbon. 2018.04.063.

146

CHAPTER 6 MULTISCALE ENGINEERING OF CARBON NANOTUBE FIBERS

[42] R.J. Davies, C. Riekel, K.K. Koziol, J.J. Vilatela, A.H. Windle, Structural studies on carbon nanotube fibres by synchrotron radiation microdiffraction and microfluorescence, J. Appl. Crystallogr. 42 (2009) 1122–1128, https://doi.org/10.1107/S0021889809036280. [43] V. Pichot, M. Burghammer, S. Badaire, C. Zakri, C. Riekel, P. Poulin, P. Launois, X-ray microdiffraction study of single-walled carbon nanotube alignment across a fibre, Europhys. Lett. 79 (2007) 46002, https://doi.org/10.1209/0295-5075/79/46002. [44] J.J. Hermans, P.H. Hermans, D. Vermaas, A. Weidinger, Quantitative evaluation of orientation in cellulose fibres from the X-ray fibre diagram, Recl. Trav. Chim. Pays Bas 65 (1946) 427–447, https://doi.org/10.1002/ recl.19460650605. [45] T. Liu, S. Kumar, Quantitative characterization of SWNT orientation by polarized Raman spectroscopy, Chem. Phys. Lett. 378 (2003) 257–262, https://doi.org/10.1016/S0009-2614(03)01287-9. [46] J.C. Ferna´ndez-Toribio, B. Alema´n, A. Ridruejo, J.J. Vilatela, Tensile properties of carbon nanotube fibres described by the fibrillar crystallite model, Carbon 133 (2018) 44–52. [47] S.R. Allen, R.J. Farris, E.L. Thomas, High modulus/high strength poly-(p-phenylene benzobisthiazole) fibres: part 2 structure-property investigations, J. Mater. Sci. 20 (1985) 4583–4592, https://doi.org/10.1007/ BF00559348. [48] K. Koziol, J. Vilatela, A. Moisala, M. Motta, P. Cunniff, M. Sennett, A. Windle, High-performance carbon nanotube fiber, Science 318 (2007) 1892–1895, https://doi.org/10.1126/science.1147635. [49] W. Lu, M. Zu, J.-H. Byun, B.-S. Kim, T.-W. Chou, State of the art of carbon nanotube fibers: opportunities and challenges, Adv. Mater. 24 (2012) 1805–1833, https://doi.org/10.1002/adma.201104672. [50] S. Zhang, L. Zhu, M.L. Minus, H.G. Chae, S. Jagannathan, C.-P. Wong, J. Kowalik, L.B. Roberson, S. Kumar, Solid-state spun fibers and yarns from 1-mm long carbon nanotube forests synthesized by water-assisted chemical vapor deposition, J. Mater. Sci. 43 (2008) 4356–4362, https://doi.org/10.1007/ s10853-008-2558-5. [51] X. Zhang, Q. Li, T.G. Holesinger, P.N. Arendt, J. Huang, P.D. Kirven, T.G. Clapp, R.F. DePaula, X. Liao, Y. Zhao, L. Zheng, D.E. Peterson, Y. Zhu, Ultrastrong, stiff, and lightweight carbon-nanotube fibers, Adv. Mater. 19 (2007) 4198–4201, https://doi.org/10.1002/adma.200700776. [52] K.R. Atkinson, S.C. Hawkins, C. Huynh, C. Skourtis, J. Dai, M. Zhang, S. Fang, A.A. Zakhidov, S.B. Lee, A. E. Aliev, C.D. Williams, R.H. Baughman, Multifunctional carbon nanotube yarns and transparent sheets: fabrication, properties, and applications, Phys. B Condens. Matter 394 (2007) 339–343, https://doi.org/ 10.1016/j.physb.2006.12.061. [53] S. Zhang, K.K.K. Koziol, I.A. Kinloch, A.H. Windle, Macroscopic fibers of well-aligned carbon nanotubes by wet spinning, Small 4 (2008) 1217–1222, https://doi.org/10.1002/smll.200700998. [54] J.J. Vilatela, A.H. Windle, A multifunctional yarn made of carbon nanotubes, J. Eng. Fibers Fabr. 7 (2012) 23–28. [55] J.J. Vilatela, J.A. Elliott, A.H. Windle, A model for the strength of yarn-like carbon nanotube fibers, ACS Nano 5 (2011) 1921–1927, https://doi.org/10.1021/nn102925a. [56] D.E. Tsentalovich, R.J. Headrick, F. Mirri, J. Hao, N. Behabtu, C.C. Young, M. Pasquali, Influence of carbon nanotube characteristics on macroscopic fiber properties, ACS Appl. Mater. Interfaces 9 (2017) 36189–36198, https://doi.org/10.1021/acsami.7b10968. [57] N. Behabtu, C.C. Young, D.E. Tsentalovich, O. Kleinerman, X. Wang, A.W.K. Ma, E.A. Bengio, R.F. ter Waarbeek, J.J. de Jong, R.E. Hoogerwerf, S.B. Fairchild, J.B. Ferguson, B. Maruyama, J. Kono, Y. Talmon, Y. Cohen, M.J. Otto, M. Pasquali, Strong, light, multifunctional fibers of carbon nanotubes with ultrahigh conductivity, Science 339 (2013) 182–186, https://doi.org/10.1126/science.1228061. [58] J.J. Vilatela, L. Deng, I.A. Kinloch, R.J. Young, A.H. Windle, Structure of and stress transfer in fibres spun from carbon nanotubes produced by chemical vapour deposition, Carbon 49 (2011) 4149–4158, https://doi. org/10.1016/j.carbon.2011.05.045.

REFERENCES

147

[59] T.S. Gspann, S.M. Juckes, J.F. Niven, M.B. Johnson, J.A. Elliott, M.A. White, A.H. Windle, High thermal conductivities of carbon nanotube films and micro-fibres and their dependence on morphology, Carbon 114 (2017) 160–168, https://doi.org/10.1016/j.carbon.2016.12.006. [60] B. Alema´n, M. Vila, J.J. Vilatela, Surface chemistry analysis of carbon nanotube fibers by X-ray photoelectron spectroscopy, Phys. Status Solidi A, 1800187, https://doi.org/10.1002/pssa.201800187. [61] D. Iglesias, E. Senokos, B. Aleman, L. Cabana, C. Navio, R. Marcilla, M. Prato, J.J. Vilatela, S. Marchesan, Gas-phase functionalization of macroscopic carbon nanotube fiber assemblies: reaction control, electrochemical properties, and use for flexible supercapacitors, ACS Appl. Mater. Interfaces (2018) https://doi.org/10.1021/acsami.7b15973. [62] S. Maruyama, Kataura-Plot for Resonant Raman, (n.d.). http://www.photon.t.u-tokyo.ac.jp/maruyama/ kataura/kataura.html#DOS. [63] H. Ji, X. Zhao, Z. Qiao, J. Jung, Y. Zhu, Y. Lu, L.L. Zhang, A.H. MacDonald, R.S. Ruoff, Capacitance of carbon-based electrical double-layer capacitors, Nat. Commun. 5 (2014)3317https://doi.org/10.1038/ ncomms4317. [64] J. Bisquert, Nanostructured Energy Devices: Equilibrium Concepts and Kinetics, CRC Press, 2014.https:// www.crcpress.com/Nanostructured-Energy-Devices-Equilibrium-Concepts-and-Kinetics/Bisquert/p/book/ 9781439836026. Accessed 30 November 2017. [65] H.-S.P. Wong, D. Akinwande, Carbon Nanotube and Graphene Device Physics, Cambridge University Press, Cambridge, 2010. https://doi.org/10.1017/CBO9780511778124. [66] A.A. Maarouf, C.L. Kane, E.J. Mele, Electronic structure of carbon nanotube ropes, Phys. Rev. B 61 (2000) 11156–11165, https://doi.org/10.1103/PhysRevB.61.11156. [67] S.M. Mousavi-Khoshdel, E. Targholi, Exploring the effect of functionalization of graphene on the quantum capacitance by first principle study, Carbon 89 (2015) 148–160, https://doi.org/10.1016/j.carbon.. 2015.03.013.