Strong carbon nanofibers from electrospun polyacrylonitrile

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Strong carbon nanofibers from electrospun polyacrylonitrile Salman N. Arshad, Mohammad Naraghi, Ioannis Chasiotis

*

Department of Aerospace Engineering, University of Illinois at Urbana-Champaign, 104 South Wright Street, Urbana, IL 61801, USA

A R T I C L E I N F O

A B S T R A C T

Article history:

Strong carbon nanofibers with diameters between 150 nm and 500 nm and lengths of the

Received 28 October 2010

order of centimeters were realized from electrospun polyacrylonitrile (PAN). Their tensile

Accepted 23 December 2010

strength reached a maximum at 1400 C carbonization temperature, while the elastic mod-

Available online 30 December 2010

ulus increased monotonically until 1700 C. For most carbonization temperatures, both properties increased with reduced nanofiber diameter. The tensile strength and the elastic modulus, measured from individual nanofibers carbonized at 1400 C, averaged 3.5 ± 0.6 GPa and 172 ± 40 GPa, respectively, while some nanofibers reached 2% ultimate strain and strengths over 4.5 GPa. The average tensile strength and elastic modulus of carbon nanofibers produced at 1400 C were six and three times higher than in previous reports, respectively. These high mechanical property values were achieved for optimum electrospinning parameters yielding strong PAN nanofibers, and optimum stabilization and carbonization temperatures, which resulted in smooth carbon nanofiber surfaces and homogeneous nanofiber cross-sections, as opposed to a previously reported core–shell structure. Turbostratic carbon crystallites with average thickness increasing from 3 to 8 layers between 800 C and 1700 C improved the elastic modulus and the tensile strength but their large size, discontinuous form, and random orientation reduced the tensile strength at carbonization temperatures higher than 1400 C.  2010 Elsevier Ltd. All rights reserved.

1.

Introduction

Carbon nanostructures are emerging multifunctional materials for advanced polymer matrix composites because of their high strength, elastic modulus, thermal and electrical conductivity and relatively low density [1–3]. Their applications include structural laminate and woven composites with improved matrix toughness for the aerospace and automotive sectors, and filters, scaffolds and fuel cells for bio-medical and energy applications [4,5]. Existing carbon nanomaterials that serve as polymer matrix reinforcements include carbon nanotubes (CNTs), vapor grown carbon nanofibers (VGCNFs) and other advanced structural forms of 1D carbon [6–8]. While CNTs and VGCNFs can provide matrix toughening [2,8–14], their discontinuous and entangled form does not support strengthening. On the other hand, carbon nanofibers in

continuous and aligned form have been derived from electrospun polymer nanofibers, such as polyacrylonitrile (PAN) and pitch [15–19]. PAN is the predominant precursor for microscale carbon fibers suitable for structural applications due to its high yield and the flexibility to tailor the fiber strength and elastic modulus by controlling the carbonization and graphitization temperatures [20]. While microscale carbon fibers from PAN precursors developed over the last four decades have typical diameters in the range of 5–10 lm and tensile strengths that reach 7 GPa, significantly less research has been devoted to their nanoscale counterparts which are derived from electrospun PAN nanofibers with diameters in the range of 100–500 nm [16,21]. A major advantage of nanofibers is their 1000–10,000 times smaller cross-sectional area which provides tremendous refinement and improved interaction with polymer matrices, thus

* Corresponding author: Fax: +1 217 244 1474. E-mail addresses: [email protected], [email protected] (I. Chasiotis). 0008-6223/$ - see front matter  2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.carbon.2010.12.056

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increasing the matrix shear strength. Since carbon nanofibers derived from electrospun PAN are long (centimeter long nanofibers are fabricated routinely), continuous and relatively well aligned, they are ideal for strengthening, stiffening and toughening of polymers [22–24]. PAN nanofibers are converted into carbon nanofibers by the processes of stabilization, carbonization and graphitization [25]. The carbon precursor PAN nanofibers are fabricated by electrospinning [26–28], in which a high voltage of 10–30 kV is applied between a fine nozzle, containing PAN solution, and a nanofiber collector. Upon application of a high voltage, a droplet of PAN solution, held together by surface tension at the nozzle tip, forms a Taylor cone and nanofibers are driven towards the collector because built-up electric charges overcome the surface tension that holds the droplet together and carry with them the polymer molecules [29]. While traveling towards the collector, the polymer jet undergoes several instabilities whereby its diameter decreases and major portion of the solvent evaporates [30]. The nanofibers gathered on the collector are continuous and can be aligned depending on the collector type [17]. Fabrication of carbon nanofibers from PAN proceeds with stabilization in air at 200–300 C while the PAN nanofibers are held in tension. During stabilization, PAN undergoes cyclization and partly dehydrogenation which makes it denser and stable to retain its fibrous structure during subsequent high temperature carbonization [25,31]. The stabilized PAN nanofibers are carbonized at >800 C in inert atmosphere, when the carbon content increases dramatically maintaining an amorphous structure with limited crystallinity. The structure and mechanical properties of commercial microscale carbon fibers as a function of heat treatment are well established [31–33]: at carbonization temperatures 1000–1500 C the tensile strength increases because of increased carbon content, while the turbostratic carbon crystallite size remains small. The maximum strength is achieved at 1500–1700 C while the elastic modulus increases monotonically with temperature due to the increased crystallite size and volume fraction, especially at temperatures between 2000 C and 3000 C. Microscale carbon fibers have tensile strengths and elastic moduli between 3.8–7 GPa and 230–440 GPa, respectively [34]. Unfortunately, equally high tensile strength and modulus have not been demonstrated for PAN-derived carbon nanofibers. Zussman et al. were among the first to report on PAN derived carbon nanofibers [17]: they presented tensile strengths between 0.32 GPa and 0.9 GPa and Young’s modulus of 63 ± 7 GPa, which are about six times lower than those of microscale carbon fibers. The authors identified the nanofiber skin-core cross-sectional structure as the reason for the low mechanical properties [17]. Similarly, Zhou et al. [15] reported on nanofiber bundles with 300–600 MPa tensile strength and 40–60 GPa Young’s modulus, which showed increasing trends with carbonization temperatures between temperatures 1000 C and 2200 C but they were still well below the properties of microscale carbon fibers. Higher strengths and moduli have only been reported for large and small microscale carbon fibers from gel-spun PAN and PAN-carbon nanotube composites by Chae et al. [35,36]. Their experiments on carbon fiber bundles resulted in tensile strength and modulus of 3.2 GPa and 337 GPa, respectively, while experiments performed on CNT reinforced carbon fibers resulted in tensile strength and

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modulus of 4.5 GPa and 463 GPa, respectively, which are comparable to high quality commercial carbon fibers. Thus, continuous carbon nanofibers with diameters of the order of 100 nm, or more, still have not reached their potential due to structural homogenization and defect reduction issues that have limited their mechanical strength [22]. The only carbon nanofibers with diameters of the order of 150–300 nm whose different grades have high tensile strengths between 2.74 GPa and 3.34 GPa, and Young’s modulus between 180 GPa and 245 GPa are the VGCNFs [37], which, however, are discontinuous, only about 100 lm long and wavy. As a consequence, they do not provide the expected stiffening in the low strain regime when they are incorporated into epoxy matrices [37,38]. This work reports on strong carbon nanofibers following a process optimization with experiments on individual PAN nanofibers in order to identify conditions for improved molecular orientation in the precursor nanofibers, and experiments on individual carbon nanofibers to identify the optimum carbonization conditions for high tensile strength and modulus. The experimental results are corroborated with TEM images that provided information about the size, orientation and distribution of the turbostratic carbon crystallites as a function of carbonization temperature.

2.

Experimental

In order to fabricate PAN nanofibers, Polyacrylonitrile (Sigma Aldrich) with molecular weight Mw = 150,000 g/mol was dissolved in N,N-dimethylformamide (Sigma–Aldrich) at room temperature for 24 h to form a 9 wt.% solution. A home-built electrospinning apparatus with a high voltage power supply was used to spin the PAN solution. The electrospinning voltage and the distance to the collector were varied between 15–25 kV and 15–25 cm, respectively, and individual PAN nanofibers were tested under each condition. Based on the mechanical property results from individual PAN nanofibers discussed in the Results section, only nanofibres fabricated at 25 kV and 25 cm distance from the collector were stabilized and carbonized. Continuous PAN nanofibers were collected on grounded parallel steel wires with 1 cm spacing, thus forming a unidirectional net of fibers. The PAN nanofibers were picked-up from the collector on metallic clips designed to thermally expand with increased temperature and, therefore, maintain tension on the nanofibers during stabilization and carbonization. Stabilization of PAN nanofibers was conducted in a furnace by heating in air from room temperature to 300 C at a rate of 5 C/min and 1 h hold time at the peak temperature. The optimum temperature and time of stabilization were determined by differential scanning calorimetry (DSC). Four sets of nanofibers, stabilized at optimum conditions, were carbonized in a high temperature tube furnace for 1 h in a N2 atmosphere and at peak temperatures 800 C, 1100 C, 1400 C and 1700 C. A heating rate of 5 C/min was used in carbonization to reach directly the desired temperature, as opposed to two-step processes used before [17]. The PAN and carbon nanofibers were inspected for uniformity and surface defects under a scanning electron microscope (SEM), while transmission electron microscopy (TEM) was employed to investigate the nanofiber structure at different

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Fig. 1 – (a) PAN nanofiber mounted on a MEMS loading platform. (b) Mechanical behavior of PAN nanofibers fabricated by different electrospinning conditions. The legend entries are (in order) voltage (kV), spinning distance (cm) and nanofiber diameter (nm).

carbonization temperatures and to measure the average turbostratic carbon crystallite thickness. A microelectromechanical (MEMS)-based nanoscale testing platform with a high resolution optics-based method for mechanical property experiments at the nanoscale, developed to test individual polymer and VGCNFs [37,39–41], was used to measure the stress versus strain curves of individual PAN and carbon nanofibers. Fig. 1a shows a PAN nanofiber mounted on a MEMS platform for nanofiber testing. A focused ion beam (FIB) was used to deposit Pt at both ends of the carbon nanofiber before testing to ensure rigid mounting. The MEMS platform was actuated by an external piezoelectric device and images of the loadcell opening and the distance between the grips (i.e. the nanofiber gauge length) were recorded concurrently by a CCD camera at 400· optical magnification as described by Naraghi et al. [39]. As part of this method, digital image correlation (DIC) analysis was performed to calculate the loadcell opening and the nanofiber extension with displacement resolution of about 25 nm. The loadcell stiffness was measured by a traceable method of suspending glass spheres of known weights while recording the corresponding loadcell openings [42].

3.

Results

To date, the only work on molecular alignment in PAN nanofibers employed a rotating collector [43]. In the present research, optimization experiments were carried out to determine the optimum electrospinning conditions for improved molecular orientation and uniform cross-section of PAN nanofibers as collected by electrospinning. Fig. 1b shows the effect of different electrospinning parameters on the elastic–plastic mechanical response of PAN nanofibers. The figure legend includes the PAN nanofiber diameters which were reduced by about 50% after carbonization. Nanofibers spun at an average electric field of 1 kV/cm had higher elastic modulus, yield strength and similar ductility as those fabricated at higher electric field intensities. Furthermore, nanofibers spun at the longest distances had the highest modulus and tensile

strength, which points out to improved molecular orientation which is critical for high mechanical properties of the derived carbon nanofibers. Increased molecular orientation was confirmed by FTIR measurements [44], resulting in orientation factors that were twice as high (f = 0.52) for the nanofibers with the highest mechanical strength in Fig. 1b. Similar orientation factors have been reported from X-ray measurements for microscale PAN fibers used as precursors for carbon fibers [45]. Short electrospinning distances from the collector (15 cm) had limited or no molecule-stretching effect and perhaps resulted in increased solvent content in the nanofibers [26,46], while long electrospinning distances permitted multiple bending instabilities and evaporation of the majority of the solvent whose presence promotes (undesirable) molecular relaxations at short electrospinning distances. The optimum temperature and time for stabilization were determined by DSC. Sample curves are shown in Fig. 2a, where three large samples of PAN nanofibers were heated at 5 C/min to 250 C, 275 C, and 300 C and were held at the peak temperature for 1 h. Stabilization of PAN is an exothermic reaction and a DSC scan shows the amount of heat released as a function of time and, therefore, the completion of the reaction. The exothermic reaction was not completed at 250 C and 275 C and the nanofiber samples continued to release heat even after 1 h. However, the reaction was completed after 1 h at 300 C and the released heat was dramatically more than at 250 C and 275 C. A second scan was done at 300 C but no further heat was released, which confirmed that stabilization was completed in the first heating cycle. Stabilization temperatures higher than 300 C can result in combustion of the fibers, and therefore, were not pursued. The stabilized nanofibers were then exposed to temperatures 800–1700 C to derive the carbon nanofibers. Fourier Transform Infrared (FTIR) spectra of the as-spun PAN nanofibers, and those stabilized at 300 C and carbonized at 800 C are shown in Fig. 2b. The characteristic vibrations for the chemical groups in PAN are: at 2241–2243 cm 1 due to the C„N nitrile group [47,48], the vibrations of the aliphatic CH

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groups (CH, CH2, and CH3) at 2870–2931 cm 1, 1450– 1460 cm 1, 1350–1380 cm 1 and 1220–1270 cm 1, the strong band at 1732 cm 1 is the C@O stretching and the band at 1684 cm 1 is due to the amide group. After stabilization the most prominent structural changes are the reduction of the 2241–2243 cm 1 peak intensity, which is attributed to the C„N nitrile group, the reduction of the intensity of the aliphatic CH groups and the reduction of the peak intensity of the amide group. The appearance of the peak at 1590 cm 1 is due to a mix of C@N, C@C, and N–H groups. Most importantly, C„N is converted into C@N which results from cyclization and cross-linking and prepares the chemical structure for subsequent high temperature carbonization as reported in previous literature for carbon fibers. The appearance of the C@C group results from dehydrogenation. The FTIR spectra of the carbonized fibers do not contain any structural information because the black carbon nanofibers have very high absorbance. Fig. 3a and b shows SEM and TEM images of PAN-derived carbon nanofibers. The carbon nanofibers in Fig. 3a have smooth surfaces and uniform diameter along their length. The diameters of different nanofibers can be very different, as shown in Fig. 3b. The nanofibers were rod-like straight,

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which is an advantage compared to VGCNFs, which, because of their waviness, do not provide appreciable stiffening to a polymer matrix at strains less than 1% [37]. The TEM images of carbon nanofibers in Fig. 4a–d show the formation of randomly oriented crystallites at all carbonization temperatures, which are more pronounced at 1700 C. Individual carbon nanofibers were mounted on the MEMS nanofiber testing platform, shown in Fig. 5a, and tested as described in Ref. [37]. A representative stress–strain curve of a carbon nanofiber is shown in Fig. 5b. As expected, the nanofibers behaved in a linearly elastic manner starting at zero strain and until their failure at strains that in some cases approached 2% at strengths that exceeded 4.5 GPa. The tensile strength versus diameter for nanofibers carbonized at different temperatures is shown in Fig. 6a. Fibers carbonized at 800 C and 1100 C showed a dependence of strength on diameter, with larger diameters resulting in smaller tensile strength values. In Fig. 6b, the average nanofiber strength is plotted as a function of the carbonization temperature to identify the optimal processing conditions for maximum strength which was achieved at 1400 C. Similarly, the Young’s modulus of the carbon nanofibers did depend on the nanofiber diameter for all carbonization temperatures,

Fig. 2 – (a) DSC profiles of PAN nanofibers stabilized at 250 C, 275 C and 300 C for 1 h. (b) FTIR spectra of as-spun PAN nanofibers, nanofibers stabilized at 300 C and nanofibers carbonized at 800 C.

Fig. 3 – (a) SEM image of carbon nanofibers. (b) TEM image showing the range of carbon nanofiber diameters that are possible and their cross-sectional uniformity without evidence of skin-core structure.

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Fig. 4 – TEM images of carbon nanofibers carbonized at (a) 800 C, (b) 1100 C, (c) 1400 C and (d) 1700 C showing the increased size and density of turbostratic carbon crystallites. The crystallites are more discernible near the nanofiber edges that are sufficiently thin for TEM imaging.

as shown in Fig. 6c, and it increased monotonically with temperature as shown in Fig. 6d.

4.

Discussion

Table 1 summarizes all the mechanical and structural properties of the carbon nanofibers as a function of carbonization temperature. A reduction in the nanofiber tensile strength with increasing diameter was observed at the lower carbonization temperatures of 800 C and 1100 C: the strength of nanofibers carbonized at 800 C increased by almost 100% when the diameter was reduced from 500 nm to 200 nm. TEM images of all carbon nanofibers, as shown for example in Figs. 3b and 4a–d, revealed no porosity or other discernible defects, except for a small surface roughness. A study of the mechanical properties of the PAN nanofibers by Naraghi et al. [41,44,49] showed that their elastic modulus and engineering strength also increased by more than six times when the diameter was reduced from 800 nm to 200 nm. FTIR studies also showed that small diameter PAN nanofibers had increased molecular orientation [49], which could be the reason for the scale dependent strength and modulus of nanofibers carbonized at 800 C and 1100 C. At these temperatures the

non-carbon elements are removed during carbonization more easily in thinner than thicker nanofibers. As shown in Fig. 4a and b and discussed later in this section, the crystallite size at 800 C and 1100 C was too small to affect the scaling of the mechanical properties. Thus, a diameter scaling of the mechanical strength must be attributed to the properties of the original PAN. For nanofibers carbonized at up to 1400 C, increasing carbonization temperature increased the fiber tensile strength reaching 3.5 ± 0.6 GPa which is six times higher than the average strength reported before for carbon nanofibers of the same dimensions but carbonized at lower temperatures (1100 C) [17], or tested in a bundle form [15]. The initial rise in strength with carbonization temperature at 800 C and 1100 C is usually explained by the increasing carbon content and nanofiber densification. TEM images of carbon nanofibers produced at all temperatures, e.g. Fig. 3b, always showed homogeneous cross-sections without any evidence of a skin-core structure. Prior studies identified the heterogeneous skin-core structure as the reason for the significant suppression of the mechanical strength and modulus of carbon nanofibers [17], and thus, the homogeneity of the present nanofibers is one of the reasons for the high property values

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Fig. 5 – (a) Carbon nanofiber mounted on a MEMS device showing a detail of the specimen grips. (b) Engineering stress–strain curve from a single nanofiber carbonized at 1400 C.

reported here. It is also important to note that the nanofiber strength was not found to depend on the nanofiber diameter at 1400 C carbonization temperature. The tensile strength dropped precipitously for nanofibers produced at 1700 C. This reduction in mechanical strength is due to the evolving crystalline structure shown in Fig. 4a– d: increased carbonization temperature resulted in the formation and growth of randomly oriented turbostratic carbon crystallites which caused early fiber rupture as a consequence of the stress mismatch with the surrounding amorphous carbon. The highest stiffness constant of graphite can exceed 1 TPa [50], which is significantly larger than the average stiffness of the surrounding amorphous carbon. As the two phases are approximately under the same strain, the mismatch stress rises dramatically for larger crystallites causing crack nucleation and instant brittle fracture.

A large number of TEM images of the carbon nanofibers were obtained to measure the average crystallite thickness, Lc, and length, La, for different carbonization temperatures. Lc and La both increased with increasing carbonization temperature: as listed in Table 1, the average crystallite thickness increased from an average of 3.3 ± 0.9 carbon layers at 800 C, which is in good agreement with previous reports for micron size diameter fibers [32,33,35], commercial (T-300) fibers [35], and nanoscale fibers [17], but it is higher than that reported before by Zhou et al. for similar size nanofibers processed between 800 C and 1400 C [32], to an average of 7.9 ± 1.9 layers at 1700 C. The average crystallite thickness of microscale PAN derived carbon fibers carbonized at 1800 C has been reported to be 8–10 carbon layers [33] which is similar to the present values, suggesting that the nanoscale size of the fibers does not affect the growth of the carbon crystallites. It should be noted that the crystallite size for the carbonization

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Fig. 6 – (a) Tensile strength data versus nanofiber diameter for four carbonization temperatures. (b) Average nanofiber strength versus carbonization temperature. (c) Elastic modulus data versus nanofiber diameter, and (d) average elastic modulus versus carbonization temperature.

Table 1 – Mechanical properties and crystallite thickness as a function of carbonization temperature. The standard deviation is provided for each average property value. Carbonization temperature (C)

Carbon content (%)

Young’s modulus (GPa)

Tensile strength (GPa)

Characteristic strength rc (GPa)

Weibull modulus

Crystallite thickness (# of layers)

800 1100 1400 1700

81.2 92.7 N/A N/A

80 ± 19 105 ± 27 172 ± 40 191 ± 58

1.86 ± 0.55 2.30 ± 0.70 3.52 ± 0.64 2.05 ± 0.70

2.20 2.90 3.60 2.30

3.1 6.4 5.9 3.0

3.3 ± 0.9 3.9 ± 0.9 6.6 ± 1.4 7.9 ± 1.9

temperature of 1100 C is very comparable to that reported for PAN derived carbon nanofibers with significantly lower tensile strength and modulus implying that the dramatic improvement in the mechanical properties reported in this work is owed to the nanofiber homogeneity across its thickness. The Young’s modulus, on the other hand, depended on the nanofiber diameter for all carbonization temperatures, as shown in Fig. 6c, reaching a maximum average value of 191 ± 58 GPa at 1700 C. While for nanofibers produced at 800 C and 1100 C the elastic modulus scaling with diameter could be directly attributed to a similar scaling of PAN nanofibers [41,44,49], at the higher temperatures of 1400 C and 1700 C the structure of the nanofibers is dominated by the presence of the turbostratic carbon crystallites as clearly shown in Fig. 4c and d. The larger density and size of crystallites at higher temperatures resulted in a ‘‘composite’’ nanofiber with higher stiffness. The elastic modulus scaling with nanofiber diameter at 1400 C and 1700 C could be due to in-

creased density and size of crystallites at and near the nanofiber surface compared to its interior, especially for thicker nanofibers. Prior works reported on preferred alignment of turbostratic carbon crystallites at the nanofiber surface [17,51]. However, it was not possible to measure the crystallite size and distribution in the interior of nanofibers thicker than 100 nm with a TEM. It was evidenced, however, that very thin nanofibers with diameters 50–100 nm, had significant crystallite density and sizes in their interior, as shown in Fig. 7a and b. It should be noted that even in the thinnest nanofibers, the crystallites were not aligned with the fiber axis. In comparison with other reports on PAN-derived and other forms of carbon nanofibers, the tensile strength and the elastic modulus of the present carbon nanofibers were six and three times larger than previously reported, respectively, as a result of selecting optimal conditions for PAN electrospinning. More importantly, the commercial carbon T-300 (Toray Industries Inc.) have mechanical strength of 3.53 GPa [1,31,35,52], which is very close to that reported in this work

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Fig. 7 – TEM image and detail of a carbon nanofiber carbonized at 1400 C with large, randomly oriented, crystallites.

for PAN nanofibers carbonized at the same temperature as the T-300 fibers, namely 1400 C. Similar strength (3.2 ± 0.7 GPa) but higher modulus (337 ± 38 GPa) have been reported for highly drawn microscale PAN-derived carbon fibers [36], obtained by the islands-in-a-sea method, which indicates that mechanical drawing before stabilization does increase the elastic modulus of the final carbon fibers. Finally, it is worth mentioning the force-bearing capacity of the nanofibers reported here exceeds that of other forms of nanoscale carbon such as CNTs. PAN nanofibers carbonized at 1400 C with 200 nm diameter carried at least 50 lN of force before failure, which is 20 times higher than the 2.68 lN sustained by 26 nm diameter (gage length of 2.1 lm) as-grown multi-walled carbon nanotubes (MWCNTs), and comparable to that of 49 nm diameter (gage length of 1.9 lm) irradiated MWCNTs that have been reported to sustain 60.5 lN [53]. The carbon nanofibers were brittle and potential extrapolations of their failure properties could be made by fitting the Weibull probability density function to the strength data, which yields the two Weibull parameters: the characteristic strength, rc, and the Weibull modulus m. Their values are tabulated in Table 1. As the characteristic strength increased from 2.2 GPa to 3.6 GPa for nanofibers produced between 800 C and 1400 C, the Weibull modulus also increased to about 6, which is an average value for brittle materials. The Weibull modulus is a measure of the distribution and variability of the flaw sizes in a material. Large values (>10–15) indicate small dependence of the mechanical strength on the specimen size and, therefore, for large values of m a well-defined flaw size and distribution exist. Small values of m (<5–6) indicate a diverse population of flaws in size and/or in orientation. The mechanical strength scales with the specimen size as r1/r2 = (‘2/‘1)1/m, where r1 and r2 are the failure strengths of specimens with ‘‘sizes’’ ‘1 and ‘2, respectively [54]. ‘1 and ‘2 may denote the specimen length, surface area or volume depending whether the flaws that cause failure are evenly distributed along the specimen length, its surface or its volume. It is evident from this equation that for m  6 (fibers produced at 1400 C) the nanofiber strength scales rather weakly with its length. This favorable trend changes for carbonization at 1700 C for which m  3. As described earlier, this was due to the large and randomly distributed turbostratic carbon crystallites which acted as stress concen-

trations and sites for failure initiation. This random distribution and size of crystallites were captured by the low Weibull modulus and the reduced characteristic strength of carbon nanofibers produced at 1700 C.

5.

Conclusions

This investigation pursued an optimization process to establish fabrication-structure–property relationships for strong carbon nanofibers derived from electrospun PAN. The tensile strength and the elastic modulus of the carbon nanofibers were six and three times larger than previously reported as a result of identifying appropriate conditions for PAN electrospinning, stabilization and carbonization. The homogeneous nanofiber cross-section eliminated the failure prone skin-core structure that has been reported before as a major structural weakness of this class of nanofibers. The nanofiber tensile strength increased monotonically reaching a maximum at 1400 C, while the elastic modulus increased steadily until 1700 C. The formation of turbostratic carbon crystallites with 3–8 layers in thickness and the initial molecular orientation in the PAN nanofibers increased their elastic modulus with processing temperature. However, the increased crystallite size was also the source for the drastic reduction in strength at 1700 C carbonization temperature. Moreover, the random orientation of the crystallites points out to the necessity for better molecular orientation in the PAN nanofibers. Finally, compared to the equally strong but discontinuous VGCNFs, the present nanofibers can provide immediate load transfer when embedded in a polymer matrix because of their rod-like geometry as opposed to the wavy structure of VGCNFs. The improved mechanical properties owed to the smooth nanofiber surface and their homogeneous cross-sections supported average tensile strengths as high as those of the T-300 commercial carbon fibers and the load bearing capacity of the strongest cross-linked MWCNTs, which emphasizes the potential of PAN-based carbon nanofibers.

Acknowledgements The authors acknowledge the support by the Solid Mechanics Program on Composites for Marine Structures under ONR

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grant #N00014-07-1-0888. The authors also thank Mr. Wacek Swiech in Center for Microanalysis at University of Illinois at Urbana-Champaign for his invaluable help with TEM imaging.

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