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Regenerated cellulose-silk fibroin blends fibers
International Journal of Biological Macromolecules 43 (2008) 106–114
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International Journal of Biological Macromolecules journal homepage: www.elsevier.com/locate/ijbiomac
Regenerated cellulose-silk fibroin blends fibers Enrico Marsano a,∗ , Paola Corsini a , Maurizio Canetti b , Giuliano Freddi c a Dipartimento di Chimica e Chimica Industriale, Universit` a di Genova, Via Dodecaneso 31, 16146 Genova, Italy b Istituto per lo Studio delle Macromolecole C.N.R., Via Bassini 15, 20133 Milano, Italy c Stazione Sperimentale per la Seta, Via Giuseppe Colombo 83, 20133 Milano, Italy
a r t i c l e
i n f o
Article history: Received 29 February 2008 Received in revised form 28 March 2008 Accepted 31 March 2008 Available online 11 April 2008 Keywords: Cellulose Silk fibroin Blends fibers NMMO Wet spinning
a b s t r a c t Fibers made of cellulose and silk fibroin at different composition were wet spun from solutions by using N-methylmorpholine N-oxide hydrates (NMMO/H2 O) as solvent and ethanol as coagulant. Different spinning conditions were used. The fibers were characterized by different techniques: FTIR-Raman, scanning electron microscopy, wide-angle x-ray diffraction, DSC analysis. The results evidence a phase separation in the whole blends compositions. The tensile characterization, however, illustrates that the properties of the blends fibers are higher respect to a linear behaviour between the pure polymers, confirming a good compatibility between cellulose and silk fibroin. The fibers containing 75% of cellulose show better mechanical properties than pure cellulose fibers: modulus of about 23 GPa and strength to break of 307 MPa. © 2008 Elsevier B.V. All rights reserved.
1. Introduction The preparation of fibers with a variety of functions and properties has attracted the interest of scientists for many years. In particular spinning of polymers blends is followed to get highperformance fibers and composite materials [1–3]. Silk fibroin (SF), a fibrous protein consisting of Glycine, Alanine and Serine as main amino acid residues, is one of the most extensively studied materials due to its good biological compatibility as well as biodegradability [4,5]. Moreover, the natural silk fibers are one of the strongest and toughest materials mainly because of the dominance of well orientated -sheet structures of protein chains [6,7]. Nevertheless until now the obtained regenerated SF materials, in which the random coil conformation is predominant, are brittle. So the possibility of engineering silk-based biomaterials by blending with other polymers [8], is a subject of great interest and of intensive investigations due to the simplicity and effectiveness of mixing two different polymers to obtain new materials. Especially blending silk with cellulose (CE) is an important way to exploit for realizing new environmentally friendly artificial materials. The usefulness of cellulose is due to its abundance, low cost and peculiar properties. In the textile field cellulose, for instance, is widely used
∗ Corresponding author. Tel.: +39 0103538727; fax: +39 0103538727. E-mail address: [email protected] (E. Marsano). 0141-8130/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ijbiomac.2008.03.009
thanks to its superior moisture absorbing property and biodegradability [9]. Nowadays in modern industry blends based on synthetic polymers are widely used in order to develop new polymeric manufacts, able to yield property profiles superior to those of the individual components. This method is usually cheaper and less time-consuming for the creation of a novel polymeric materials with new properties, which may depend on different factors, such as molecular structure, miscibility or compatibility between the polymers, blend composition, morphology, and processing conditions [10–12]. A great advantage of polymer blends is that the materials’ properties can be tailored by combining component polymers and changing blend composition. The blending process must be performed in a fluid state that can be reached at temperature higher than the melting or glass transition temperature for semicrystalline or amorphous polymers. When polymers decompose before they undergo melt flow, the blending must be performed by the dissolution of the polymers in a common solvents followed by elimination of the solvent through evaporation or coagulation in a non-solvent. In the case of natural polymer as silk or cellulose the melting temperature is higher than the decomposition one; hence for these polymers it is necessary to find a solvent able to dissolve both. Studies regarding the phase behaviour of blend materials, often in form of film, based on SF and natural or synthetic polymers such as: cellulose [13–15], chitosan [16,17], sodium polyglutamate [18], sodium alginate [19], poly(vinyl alcohol) [20–22], nylon [23]
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have been reported. Most samples showed a good compatibility and an improvement of the mechanical properties [16–19,23]. Not only blend films based on SF were reported in the literature, but also fibers obtained by a wet spinning process blending SF with: cellulose [8,24,25], chitin [26], chitosan [27]. In our previous work [8] blend fibers composed by 70% CE and 30% SF were prepared by wet spinning solutions of the two polymers in DMAc-LiCl, using ethanol as coagulant. The fibers showed good mechanical properties, modulus and breaking strength in comparison to the pure CE fibers obtained using the same operative conditions. Once verified the spinnability of the dopes and the good compatibility between CE and SF, we studied the possibility to use an industrial, renewable and friendly solvent: Nmethylmorpholine-N-oxide (NMMO), a compound belonging to the organic cyclic amine oxide family, able to dissolve both SF [28] and CE [29]. Cellulosic fibers produced by using NMMO as solvent in the spinning process are commercialized with the trade name Lyocell® [30]. The process uses NMMO – or more generally a mixture of NMMO and water – to directly dissolve pulp without prior derivatization of the cellulose chains. This process is environmentally benign because the nontoxic NMMO solvent is used and almost all the solvent used is totally recycled [31]. Owing to these advantages, lots of technical developments of Lyocell® fibre have been reported by many manufacturers [32–35]. Lyocell® fibers have proved commercially successful because of better mechanical properties when compared with viscose rayon [36–39]. More recently reports on the swelling and dissolution of silk fibroin polymers in NMMO appeared [28,40,41] and regenerated SF were obtained by spinning these solutions at high polymer concentration [42,43]. In our previous papers [42,43] we described the results obtained by using a dry-jet wet spinning line, starting from SF–NMMO solutions. Regenerated SF fibers with good performance were obtained and it was observed that the tensile properties were strongly dependent on the draw ratio, which directly affects the molecular orientation of both amorphous and crystalline domains. The aim of the present work is to prepare regenerated CE/SF blend fibers, at different composition, by using NMMO/H2 O as solvent and to determine whether the polymers are compatible and if the resulting fibers have structural and physical-mechanical properties comparable to those of conventional regenerated cellulose fibers. 2. Experimental 2.1. Materials SF fibers were purified from silk cocoons degumming with water in autoclave at 120 ◦ C for 20 min, followed by extensive rinsing with warm distilled water to completely remove sericin, a gum protein surrounding SF filaments. The molecular weight of starting degummed SF fibers (Mv ) ranged from 220 to 240 kDa. To prepare SF for dissolution in NMMO, fibers were dissolved in saturated aqueous LiBr (9–10 M) at 60 ◦ C for 3 h. The solution was then filtrated to remove insoluble foreign matters, dialysed until complete removal of salt, and freeze-dried. The resulting SF in form of porous, sponge-like material was used for the dissolution tests. A commercial aqueous solution, NMMO/H2 O 50/50 (w/w), and the antioxidant n-propyl gallate (PG) were purchased from Aldrich Co. A regenerated Cellulose II sample having Mv = 63,000 (DP = 390) was supplied by Lenzing AG, Austria.
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2.2. Dope preparation Singles and mixtures of the two polymers were dispersed into the commercial aqueous solution of NMMO containing 0.7% (w/w) n-propyl gallate with respect to NMMO monohydrate. The dispersion was, then, concentrated under vacuum at 55 ◦ C to obtain NMMO hydrates with the desired water content (the maximum value was about 13.1% w/w). The water content was determined by Karl Fischer method. The overall concentration (Cp ) was fixed to 17 % (w/w) with respect to the solution. Blend solutions were prepared mixing silk and cellulose in the following ratio: 25/75, 50/50 and 75/25. In order to prepare a homogeneous solution in NMMO monohydrate, the dispersions were heated at 110 ◦ C for 20 min, under a nitrogen atmosphere, with a fast energetic stirring, controlling time by time the progress of dissolution with an optical microscope. Air bubbles were trapped in the solutions, due to the fast energetic stirring, and caused flow instability in the dry-jet wetspinning line. So before spinning the centrifugation of the dopes at high temperature was used in order to disengage air bubbles. After the centrifugation step, no phase separation was observed. 2.3. Spinning The dry-jet wet-spinning line is composed by an extrusion unit, a coagulation bath, ethanol, and two sets of spools, take up and roller. Take-up rate (V1 ) and roller rate (V2 ) can be varied so that fibers are collected with different drawn ratios. DR1 is defined as the ratio between V1 and Vo (DR1 = V1 /Vo ) where Vo is fiber’s rate at the spinneret hole. The post-spinning ratio, DR2 , is defined as the ratio between V2 and Vo (DR2 = V2 /Vo ). An extrusion rate Vo = 4 m/min, a 100 m spinneret and an air gap of 100 mm were used to extrude CE/SF-NMMO solutions. Collected fibers were further washed in ethanol for about 72 h in order to extract the residual NMMO and then they were dried and stored over CaCl2 until testing. 2.4. Fibre characterization 2.4.1. SEM analysis SEM characterisation was performed by using a Jeol JSM-6380LV microscope, at an acceleration voltage of 15 kV and 15 mm working distance. Films based on pure polymers and blends were fractured in liquid nitrogen and then mounted onto aluminium specimen stubs by means of double-sided adhesive tape and sputter-coated with a thin gold layer under rarefied Argon atmosphere, using a Polaron SC7620 Sputter Coater, with a current of 30 mA for 180 s. While the fibers were simply mounted onto aluminium specimen and coated with a thin gold layer. 2.4.2. Optical microscopy Fibers diameter were determined using an optical microscope. To measure the diameters, a small piece, 1 cm long adjacent to one side of each sample were observed. The average diameter was the result of at least six measurements on each adjacent piece. 2.4.3. Mechanical properties Mechanical properties of single fibers were measured using an Instron dynamometer mod. 5500, with a 50 mm gauge length at a crossbar rate of 30 m/min, corresponding to strain rate of 1.0 × 10−2 s−1 if elongation at break of specimen was up to 8%, and at a crossbar rate of 15 mm/min, relative strain rate of 6.7 × 10−3 s−1 if elongation at break was under 8% according to ASTM D3822-012004 standard method. Elastic modulus (E), stress at break ( b ), deformation to break (εb ) were calculated as the average of at least 15 measurements from stress–strain curves.
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2.4.4. Raman analysis FT-Raman spectra were recorded with 4 cm−1 resolution using a Bruker RFS100 instrument and NdYAG laser excitation line (1084 nm). 2.4.5. Thermal analysis Differential scanning calorimetry (DSC) was performed with a Mettler DSC 30 instrument, calibrated by an Indium standard. The calorimeter cell was flushed with 20 ml/min of N2 . Sample weight was about 3 mg and Al crucibles were used. The temperature range scanned was 25–500 ◦ C, at a heating rate of 10 ◦ C/min. Samples were pre-heated until 120 ◦ C, at 10 ◦ C/min, to remove moisture, then temperature was decreased to 25 ◦ C and the scan started. 2.4.6. X-ray data Wide-angle x-ray diffraction (WAXD) data were obtained at 20 ◦ C using a Siemens D-500 diffractometer equipped with a Siemens FK 60–10 2000 W tube (Cu K␣ radiation, = 0.154 nm). The operating voltage and current were 40 kV and 40 mA, respectively. The data were collected from 5 to 35 2 ◦ at 0.02 2 ◦ intervals, by using a silicon multi-cathode detector Vortex-EX (SII). The fibers were cut, mixed and pressed for a random panel to be submitted to WAXD measurements. The degree of crystallinity, Xc , was calculated from diffracted intensity data in the range 5–35 2 ◦ by using the area integration method [44]. 3. Results and discussion Before showing the results and discussing them, it is necessary to clarify the terminology that we use regarding miscibility and compatibility. As previously reported by Olabisi et al. [11], we use the term “miscibility” to describe polymer–polymer blends with behaviour similar to that expected of a single phase system. The term compatibility that has been used by many investigators involving various studies on polymer–polymer blend behaviour is a more general term. In a strict technological sense it can be used to describe whether a desired result occurs when two polymers are combined together and this can be correlated, for immiscible system, to good adhesion between two or more phases. For thermodynamic reasons most polymer pairs are immiscible and, as a result, phase separation could occur. Accordingly the dispersed phases could exist in different forms of spheres, ellipsoids, fibrils and lamellas, depending on the compositions, chemical, and physical structures of polymer blends, as well as processing conditions. Nevertheless, the degree of compatibility may vary widely. It is well known that the mechanical properties of immiscible blends are strongly dependent on their morphologies and, therefore, the control of the morphology of an immiscible polymer blend is extremely important for the tailoring of the final product properties [45,46]. Among different morphologies of the dispersed phase, the microfibrillar one is preferred because it can improve mechanical properties, as it is previously shown by different authors [47,48]. The morphology usually results from the complex history experienced by the different constituents during processing. During the wet-spun processing the final size, shape and distribution of the dispersed phase are determined by a wider variety of parameters such as the composition, miscibility in solution, coagulation rate and phase separation rate in the coagulation bath, shear rate/shear stress, interfacial tension among the polymers and also processing conditions, such as extrusion and take-up rate. In our previous work [8] blend fibers composed of 70% CE and 30% SF were prepared by wet spinning solutions in DMAc-LiCl and using ethanol as coagulant. The CE/SF fibers showed good mechanical properties, Young’s modulus, breaking strength and elongation
to break were quite similar, differences not more than 15% to those of CE fibers obtained in the same conditions were observed. Furthermore, small-angle x-rays analysis showed that CE/SF fibers were essentially amorphous and that there was a homogeneous dispersion of SF domains into a CE matrix. The cross-sectional size of the SF domain accounts for only few nanometres, equivalent to the dimension of 6–8 protein chains in an ordered conformation; the nano-dispersion of the two phases was the responsible of the good compatibility between the two polymers.
3.1. Influence of draw ratio on preparation of CE/SF fibers Fibers based on SF and CE blends obtained from NMMO solutions by using a laboratory dry-jet wet spinning line were prepared and characterized. All spinning processes are characterized by two drawing ratios: a coagulation ratio, DR1 , and a post-spinning ratio, DR2 . The influence of both was initially investigated by increasing the stretching of the filaments (DR1 and DR2 ) until a maximum value that can be used without the break up of the filaments. Our wet-spinning equipment allows a maximum DR1 = 6 and DR2 = 15 for an extrusion velocity (Vo ) equal to 4 m/min. The observed field of spinnability evidenced that CE could not undergo a second post-spinning stretching, thus DR2 was always equal to 1, while SF fibers could not be stretched more than three times in the first coagulation bath without observing the break-up of the filament at the extrusion unit. This behaviour could be due to different aspects: the molecular weight, the stiffness of the macromolecules, the viscoelasticity of the dope and the coagulation behaviour in solution. It is well know that cellulose and cellulose derivatives show mesophase formation in solution and that this correlates to the worm-like aspect of the macromolecules with persistence length ranging from 6 to 11 nm [49]. For instance, theoretical estimate for the persistence length of the cellulose chains in polar solvent, water, yields about 14 nm [50]. For this system a coagulation ratio higher than 1 is allowed, but once the filament is coagulated the high orientation of the stiff macromolecules does not allow further stretching. On the other hand, SF most probably has a coil conformation in NMMO solution and it maintains a coiled conformation also in the coagulated filament so that the fibers can be stretched in the post-spinning step. Another important factor in the stretching of fibers is the different coagulation behaviour of CE and SF filaments in ethanol. We observed that SF filaments after coagulation maintains a swollen behaviour with very high degree of elasticity; this can be correlated to the diffusion rate of ethanol into the SF/NMMO fibers, which was higher than the diffusion rate of NMMO into the ethanol coagulation bath. For coagulated CE filament we observed a slight decrease of the filament’s diameters in the coagulation bath and a formation of a stiff skin due to a high coagulation rate of CE in ethanol. The obtained filament is compact and cannot be further stretched. Another significant difference in the stretching behaviour can be surely correlated to the difference in the viscoelasticity of the dopes. Indeed, both systems present entanglements and they behave as synthetic polymers in a melt state. However, at the same concentration, the CE/NMMO system showed a higher viscosity with respect to the SF/NMMO one [51]. Moreover, we observed that all solutions based on CE/SF blends were not able to be stretched in the post-spinning step, even those with a very low content of cellulose, as show in Fig. 1. So fibers based on CE/SF blends were collected with different coagulation ratios (DR1 = 1–3–6), but a unitary post-spinning ratio (DR2 = 1).
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Fig. 1. The blocks show the field of spinnability of solutions of SF (position 1), CE (position 5) and their blends at different compositions (CE/SF 25–75 position 2; CE/SF 50/50 position 3; CE/SF 75/25 position 4) in function of the draw ratios DR1 and DR2 .
3.2. Raman Vibrational spectra are especially sensitive to the geometry of molecules and to the system of intramolecular and intermolecular interactions. For this reason FTIR and Raman have often been used as a useful tool to determine the nature of mixing between two polymers. The presence of a peak at a definite wavenumber would indicate the presence of a specific chemical bond, if interactions took place between the two polymers, the most obvious and significant difference would be the appearance of new peaks or shift of existing peaks. Fig. 2 shows the Raman spectra of CE, SF and CE/SF blend fibers collected applying only a draw ratio during coagulation (DR2 = 1). CE Raman spectrum presents a peak at 2894 cm−1 related to the stretching of the methylene group, while in the finger print region different peaks assigned to the structure of cellulose II can be detected: 1465 cm−1 due to bending COH, 1375 cm−1 and 1266 cm−1 due to bending of the methylene groups in the plane and out of the plane, respectively. Moreover, symmetric and asymmetric COC stretching of the glycoside groups are evident at 1114 cm−1 and 1097 cm−1 [52]. All these peaks clarify the presence of cellulosic chains in the structure of cellulose II [53]. On the other hand SF Raman spectrum, reported in Fig. 2, presents a peak at 3062 cm−1 due to the CH stretching in the C CH group, and asymmetric and symmetric stretching of methyl groups at 2985 cm−1 and 2935 cm−1 , respectively. In Fig. 2 SF Raman spectrum shows different typical peaks of silk fibroin. In particular Amide I vibrations indicative of antiparallel -sheet of silk occurred at 1663 cm−1 , while at 1229 cm−1 is evident the amide III, as well as CH2 bending scissoring and the CC stretching of the structure -sheet can be observed respectively at 1448 cm−1 and 1084 cm−1 [54–56]. The shape and intensity of all these bands can be associated with a -sheet molecular structure. In any of CE/SF blend fibers spectra neither bands shift superior to instrument resolution (4 cm−1 ), nor the appearance of new peaks occurred. As a consequence, no specific interaction between silk and cellulose could be detected. Therefore, the presence of a phase
Fig. 2. Raman spectra of the pure and blends fibers collected applying only a draw ratio during coagulation (DR1 = 3; DR2 = 1): (a) CE/SF 100/0; (b) CE/SF 75/25; (c) CE/SF 50/50; (d) CE/SF 25/75; (e) CE/SF 0/100. Above the 3200–2700 cm−1 region Raman spectra are shown, while below the finger print regions are reported.
separation in the fibers based on CE/SF blends at every prepared composition can be hypothesized. 3.3. Wide-angle X-ray diffraction characterization The typical WAXD pattern of CF/SF samples collected with drawn ratios DR1 equal to 3 are shown in Fig. 3. The profile of the CF/SF 100/0 sample is typical of regenerated cellulose fibers spun by organic solution. The reflection peaks can be indexed referring to the known cellulose II crystal structure [57,58]. The X-ray diffractograms of regenerated SF fibers shows the peak at 2 ∼ = 21◦ characteristic of the -sheet crystals of silk fibroin [59]. Intermediate profiles were collected for the CF/SF fibers prepared with different blend ratios. Analogous trend was observed for the fibers prepared with a value of DR1 equal 6. As reported in Table 1 the overall crystallinity (Xc ) decreases with the increase of the silk component in the blend. Slightly higher Xc values were calculated for pure CE and CF/SF blend fibers collected with DR1 equal to 6, even if the differences between the values are included in the error limits of the calculation method and thus they must be considered not significant. In Fig. 4, the experimental WAXD profiles of the blend fibers (DR1 = 3) are compared to the smoothed profile constructed by summing together the relative fraction contributions of the experimental pure components, CF/SF 0/100 and CF/SF 100/0. A good fitting was observed for all compositions confirming the presence
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Fig. 3. WAXD profiles of sample fibers at the different compositions collected with DR1 = 3 and DR2 = 1. (a) CE/SF 100/0; (b) CE/SF 75/25; (c) CE/SF 50/50; (d) CE/SF 25/75; (e) CE/SF 0/100. Table 1 Crystallinity, Xc , of pure and blends fibers prepared with different draw ratios, DR1 Samples Pure CE fibre
Blend fibre
Polymer weight ratio (CE/SF)
DR1
Xc (%)
100/0
3 6
32 36
75/25
3 6 3 6 3 6
24 26 23 25 21 22
1 3
21 21
50/50 25/75
Pure SF fibre
0/100
in the blend fibers of both components separately crystallized, confirming the hypothesis done previously. 3.4. Thermal analysis The properties of CE/SF blend fibers were further investigated by thermal analysis. Fig. 5 shows the DSC profiles of pure and blend fibers wet-spun from NMMO solutions and collected applying only a draw ratio during coagulation (DR2 = 1). The DSC thermogram of
Fig. 5. DSC curves of regenerated CE (a), SF (e) and CE/SF blend fibers with different blend compositions: (c) CE/SF 75/25; (d) CE/SF 50/50; (e) CE/SF 25/75. All the samples are collected applying only a draw ratio during coagulation (DR1 = 3; DR2 = 1).
CE fibers showed a weak and progressive endo shift until about 200 ◦ C. Heating at above 200 ◦ C caused a change of the trend from endo to exo. It has been reported that the thermal transitions occurring in this temperature range are associated with the onset of thermal degradation of cellulose, which consists of a sequence of events starting with rearrangement of chains in the amorphous regions, followed by main chain fission by cleavage of glucosidic bonds, volatilization of pyrolysis products (mainly levoglucosan), and dehydration, accompanied by increase of crystallinity and intermolecular cross-linking [60–62]. Beyond 300 ◦ C fast thermal degradation with extensive weight loss has been reported to occur [60]. Pure SF fibers showed an intense endothermic peak at 293 ◦ C representing the thermal decomposition of NMMO-regenerated SF fibers with oriented -sheet structure [42]. This peak clearly appeared also in the blend fibers (SF content ≥50 wt%), falling at a temperature very close to that of pure SF fibers (290.2 ◦ C and 288.4 ◦ C for the blends containing 75% and 50% SF, respectively). These features seem to indicate that the thermal behaviour of SF was only slightly influenced by blending with cellulose. With reference to WADX results we may suggest that well organized and crystalline SF domains were present in the blend fibers, especially at high SF content, and responded to the thermal stresses almost independently of the presence of cellulose. The situation was significantly different when the SF content decreased to 25 wt%. This fiber sample showed a CE-like profile with only a small SF contribution falling at 281 ◦ C in form of weak and broad endothermic peak. The contribution of CE to the DSC thermograms of blend fibers was identifiable as a high temperature shoulder of the main decomposition peak of SF. Interestingly; the shoulder resulted in an exothermic peak falling at about 350 ◦ C in the 50 wt% fiber blends, where it appeared with the highest intensity. Since this peak is present only in the blend fibers it can be attributed to some chemical and/or physical interactions between the degradation products of SF and CE, both of which are undergoing very fast and extensive thermal degradation in this high temperature range which leads to almost complete pyrolysis of the polymers. For this reason any interpretation of this thermal event can be hardly suggested. 3.5. SEM morphology
Fig. 4. Experimental and constructed WAXD profiles of blend fibers (DR1 = 3): (a) CF/SF 75/25; (b) CF/SF 50/50; (c) CF/SF 25/75.
The miscibility of CE and SF in the blends was systematically investigated by analysing the fracture morphology of films pre-
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Fig. 6. SEM cross-section morphology of SF, CE, and CE/SF blend films. (a) CE/SF 0/100; (b) CE/SF 100/0; (c and f) SF/CE 75/25; (d and g) SF/CE 50/50; (e and h) SF/CE 25/75. Mag. bars: (a–e) 10 m; (f–h) 2 m.
pared from the spinning dopes by smearing the polymers-NMMO hydrates onto a glass plate, followed by coagulation and solvent removal in an ethanol bath. As shown in Fig. 6, the fracture morphology of SF films was smooth and glassy (Fig. 6a) while that of CE films showed a granular and colloidal appearance (Fig. 6b). CE/SF blends displayed a gradual transition from one morphology to the other as a function of the blending ratio (Fig. 6c–e). Evidences of phase separation were detected in all the blends, irrespective of blend composition (Fig. 6f–h). Morphologically different domains randomly appeared in the fractured surfaces as areas with prevalently smooth or granular morphology sharply separated one from the other. These results correlate with the above spectroscopic and thermal analysis, and can be taken as indirect evidence that the two polymers tend to form separate domains when mixed in NMMO hydrates solvent. Regenerated SF, CE and CE/SF fibers were also characterized by SEM (Fig. 7). The whole fibers showed a circular shape with fairly regular size along the fiber axis. The surface morphology changed from granular to smooth for pure SF and CE fibers, respectively. Interestingly, a fine longitudinal striation was evidenced by CE/SF blend fiber. This feature appeared in all blends and was almost independent on the draw ratio applied. The morphological analysis of fiber cross-sections did not provide any clear evidence of phase separation observed in films, although the presence of cross-sectional features with a prevalently fibrous or granular structure was observed (pictures not shown).
3.6. Mechanical properties Generally mechanical properties, above all stress and elongation at break, are very important to evaluate fiber performances for proper application. Regenerated SF, CE and SF/CE blend fibers were tested for their cross-sectional dimension and tensile properties. The average mechanical properties were calculated from at least 10 curves for samples at different compositions and draw ratios, as it is shown in Table 2. Fig. 8 shows the diameters and the standard deviation of CE, SF and blend fibers as a function of the draw ratio. It can be noticed that for all the fibers a decreasing of the diameters with respect the drawing and not meaningful differences between samples with different composition were observed.
3.7. Tensile properties of pure polymers Fig. 9 illustrates the typical stress–strain plots of CE and SF samples. A different behaviour can be noticed, elongation at break of CE fibers decreased from 15 to 7%, and breaking strength increased from 180 to 360 MPa as the draw ratio during the coagulation (DR1 ) increased. For SF fibers a change from a fragile fiber to a ductile one was observed with increasing DR1 , as indicated by the values of both elongation at break and the breaking strength, which increased from 2 to 10% and from 52 to 116 MPa, respectively.
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Fig. 7. SEM photographs of SF (a), CE (b), and SF/CE blend fibers with different blend composition: (c) SF/CE 75/25; (d) SF/CE 50/50; (e) SF/CE 25/75. Mag. bars: (a–e) 5 m. All the samples are collected applying only a draw ratio during coagulation (DR1 = 3; DR2 = 1).
The mechanical behaviour of regenerated SF fibers can be understood from the microstructural analysis. As previously shown by the WADX analysis the existence of -sheet microcrystallites in all regenerated SF fibers is evident and the degree of crystallinity is approximately comparable, independently on the draw ratio in the coagulation. The change in tensile properties could be correlated with the alignment of protein chains in the amorphous regions that depends on the ability of the chains to slip at entanglements [63], which, in turn, depends on two microstructural features: the force required for the slippage of the chains at an entanglement and the density of the entanglements. In this context, the ductile behaviour of regenerated SF fibers subjected to high stretching suggests a reduction in the number of chain entanglements, facilitating chain slippage and plastic behaviour. In CE fibers, which are made by a more rigid polymer with a lower entanglement density, the increase of the draw ratio allows
an increase of the orientation of the amorphous phase that causing an increase of modulus, breaking strength and fragility of the fibers. It should be noted that the mechanical properties of our pure cellulose fibers were found to be significantly lower than those reported for the conventional regenerated cellulose fibers or films [64,65]. This can be correlated to our laboratory-scale dry-jet wet spinning equipment. 3.8. Tensile properties of blend fibers Typical stress–strain curves of wet-spun CE/SF blend fibers with different SF content are shown in Fig. 10 where four representative profiles per sample are plotted. The presence of SF in the blends should affect the orientation and mechanical strength of the fibers because a rigid polymer, cellulose, is mixed with a random coil polymer SF. According to Flory theory [66] a phase separation occurs upon blending rigid and coil polymers, beyond a critical concentration, in the absence of strong
Table 2 Polymer compositions (polymer weight ratio), average mechanical properties (Young’s modulus E, tensile strength b and elongation at break εb ) and diameters of the fibers obtained by using Vo = 4 m/min and different draw ratios DR1 in the coagulation bath Samples
Polymer weight ratio (CE/SF)
DR1
Diameter (m)
E (GPa)
b (MPa)
εb (%)
100/0
1 3 6
34.3 ± 5.2 16.1 ± 2.5 14.2 ± 1.0
11.3 ± 0.8 23.0 ± 1.5 24.2 ± 1.9
186 ± 9 260 ± 39 364 ± 50
15 ± 3 5±1 7±2
75/25
1 3 6 1 3 6 1 3 6
41.9 20.4 16.3 38.2 23.2 16.4 32.3 23.2 17.1
1 3
53.6 ± 6.9 24.5 ± 1.1
Pure CE fibers
50/50 Blend fibers 25/75
Pure SF fibers
0/100
± ± ± ± ± ± ± ± ±
1.6 6.0 2.6 0.6 0.2 0.2 1.0 0.8 0.2
11.7 22.7 21.7 14.4 16.2 18.1 11.2 11.7 14.0
± ± ± ± ± ± ± ± ±
0.8 2.9 1.0 0.7 0.7 0.9 0.6 0.4 0.4
5.4 ± 0.5 7.9 ± 0.4
159 307 261 193 236 315 148 155 169
± ± ± ± ± ± ± ± ±
10 20 18 4 16 39 5 7 8
52 ± 6 116 ± 5
8 3 3 8 7 6 4 5 7
± ± ± ± ± ± ± ± ±
2 1 1 1 1 2 1 3 1
2±1 10 ± 5
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Fig. 8. Diameters versus the draw ratio during coagulation, DR1 , of pure and blends fibers.
intermolecular interactions. The phenomenon is due to the interference of the random coil polymer with the mutual orientation of the molecules of the rod-like polymer. A qualitative agreement between theory and experimental results is reported in literature for system based on a rigid polymer and a flexible one [67,68]. The CE/SF 25/75 fibers showed behaviour similar to that of pure SF fibers. A slight increase of the elongation at break with the DR1 was observed (Table 2). Fibers composed of 50/50 CE/SF displayed a practically constant elongation at break irrespective of the draw ratio. Fibers composed of 75/25 CE/SF fibers showed a behaviour similar to pure CE, the fragility of the fibers increasing with the draw ratio. Fig. 11 shows the trend of Young’s modulus and stress at break for fibers obtained at DR1 = 3 with different blends compositions. The decrease in mechanical properties seems for all the samples to be less than the linear decreases with the amount of SF added. In particular the blend containing more than 75% of CE show higher modulus and strength than those of pure cellulose fiber, while these properties had lower values in the fiber with a CE/SF weight ratio of 50/50 and 25/75. In compatible blends, a linear relationship between blend composition and such mechanical properties as elastic modulus and strength at break is usually observed. In general, a negative deviation from the linear relationship is considered an indication of
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Fig. 10. Stress–strain plot of pure and blend fibers at different compositions obtained using a draw ratio during coagulation, DR1 = 3: (a) CE/SF 100/0; (b) CE/SF 75/25; (c) CE/SF 50/50; (d) CE/SF 25/75; (e) CE/SF 0/100.
Fig. 11. Variation of Young’s modulus () and stress at break () with CE content for CE/SF fibers at DR1 = 3.
poor compatibility between blend components, whereas a positive deviation is considered an indication of improved compatibility. The results reported in Fig. 11 are an indication of good compatibility between CE and SF fiber and put also in evidence the role of the DR in the formation of morphologies of the dispersed phase that can enhance the compatibility between the two polymers. The
Fig. 9. Stress–strain plot of the pure polymer fibers obtained using different draw ratios, DR1 , in the coagulation bath: (a) CE fibers; (b) SF fibers.
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results of the mechanical behaviours seem to indicate that with the increase of CE content in the fibers and for a moderate draw ratio (DR1 = 3) good interaction between the two phases occur. This increase of interactions could be due to the formation of morphologies that enhanced the amount of amine groups of SF that can be incorporated into the CE amorphous matrix and interacted with the OH groups in the fibers. 4. Conclusions CE/SF blends monofilaments, with diameters ranging between 14 and 41 m, were successfully prepared by extruding solutions in NMMO·H2 O in a coagulation bath containing ethanol, through a dry-jet wet spinning line. Blends fibers were obtained in different drawing conditions and the amount of SF did not influence the spinnability regions of solutions based on CE/SF respect to cellulose one. FTIR, DSC and WAXD analyses of CE/SF blends fibers evidenced that both polymers underwent crystallization upon solidification and two phase systems were obtained. The phase separation was observed in blends films with SEM; the microphotographs showed that the fracture of CE/SF blends films displayed a gradual transition with blend compositions, from a smooth and glassy morphology, typical of SF films, to a granular one characteristic of CE films. This behaviour was detected in all the blends, irrespective of composition and, therefore, the presence of a phase separation in the fibers at every prepared composition could be hypothesized. Respect to the linear behaviour of the main mechanical properties between the pure polymers, the tensile properties of the blends fibers showed higher values, although the two polymers were not miscible and showed phase separation in all samples. This is a clear indication of the existence of some kind of interactions between the two phases. In particular CE/SF blend fibers containing more than 75% of CE exhibit higher tensile properties than those of pure cellulose, obtained by using the same spinning condition (DR1 = 3). This result seems to indicate that good interactions between the two phases occur suggesting a good compatibility of these two natural polymers in the fibers. To this end, CE/SF blend monofilaments have great potential as a new regenerated biodegradable fiber with structural and physic-mechanical properties comparable to those of conventional regenerated cellulose fibers. Acknowledgments The authors wish to thank Mr. Mario Traverso for his helpful work on dry-jet wet spinning line. References [1] S. Arcidiacono, C.M. Mello, M. Butler, E.A. Welsh, J.W. Soares, A. Allen, D. Ziegler, T. Laue, S. Chase, Macromolecules 35 (2002) 1262–1266. [2] F. Buehler, V. Baron, E. Schmid, P. Meier, H.J. Schltze, USA USPatent 5516815 (1996). [3] R.D. Gilbert, X.C. Hu, R.E. Fornes, J. Appl. Polym. Sci. 58 (1995) 1365–1370. [4] A. Chiarini, P. Petrini, I. Bozzini, I. Dal Pra, U. Armato, Biomaterials 24 (2003) 789–799. [5] D.L. Kaplan, G.H. Altman, F. Diaz, C. Jakuba, T. Calabro, R.L. Horan, J. Chen, H. Lu, J. Richmond, Biomaterials 24 (2003) 401–416. [6] Z. Shao, F. Vollrath, R.J. Sirichaisit, R.J. Young, Polymer 40 (1999) 2493–2500. [7] Z. Shao, F. Vollrath, Nature 418 (2002) 741. [8] E. Marsano, M. Canetti, G. Conio, P. Corsini, G. Freddi, J. Appl. Polym. Sci. 104 (2007) 2187–2196. [9] C.H. Park, Y.K. Kang, S.S. Im, J. Appl. Polym. Sci. 94 (2004) 248–253. [10] D.R. Paul, S. Newman, Polymer Blends, Academic Press Inc., New York, 1978. [11] O. Olabisi, L.M. Robeson, M.T. Shaw, Polymer-Polymer Miscibility, Academic Press, New York, 1979.
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Contents lists available at ScienceDirect
International Journal of Biological Macromolecules journal homepage: www.elsevier.com/locate/ijbiomac
Regenerated cellulose-silk fibroin blends fibers Enrico Marsano a,∗ , Paola Corsini a , Maurizio Canetti b , Giuliano Freddi c a Dipartimento di Chimica e Chimica Industriale, Universit` a di Genova, Via Dodecaneso 31, 16146 Genova, Italy b Istituto per lo Studio delle Macromolecole C.N.R., Via Bassini 15, 20133 Milano, Italy c Stazione Sperimentale per la Seta, Via Giuseppe Colombo 83, 20133 Milano, Italy
a r t i c l e
i n f o
Article history: Received 29 February 2008 Received in revised form 28 March 2008 Accepted 31 March 2008 Available online 11 April 2008 Keywords: Cellulose Silk fibroin Blends fibers NMMO Wet spinning
a b s t r a c t Fibers made of cellulose and silk fibroin at different composition were wet spun from solutions by using N-methylmorpholine N-oxide hydrates (NMMO/H2 O) as solvent and ethanol as coagulant. Different spinning conditions were used. The fibers were characterized by different techniques: FTIR-Raman, scanning electron microscopy, wide-angle x-ray diffraction, DSC analysis. The results evidence a phase separation in the whole blends compositions. The tensile characterization, however, illustrates that the properties of the blends fibers are higher respect to a linear behaviour between the pure polymers, confirming a good compatibility between cellulose and silk fibroin. The fibers containing 75% of cellulose show better mechanical properties than pure cellulose fibers: modulus of about 23 GPa and strength to break of 307 MPa. © 2008 Elsevier B.V. All rights reserved.
1. Introduction The preparation of fibers with a variety of functions and properties has attracted the interest of scientists for many years. In particular spinning of polymers blends is followed to get highperformance fibers and composite materials [1–3]. Silk fibroin (SF), a fibrous protein consisting of Glycine, Alanine and Serine as main amino acid residues, is one of the most extensively studied materials due to its good biological compatibility as well as biodegradability [4,5]. Moreover, the natural silk fibers are one of the strongest and toughest materials mainly because of the dominance of well orientated -sheet structures of protein chains [6,7]. Nevertheless until now the obtained regenerated SF materials, in which the random coil conformation is predominant, are brittle. So the possibility of engineering silk-based biomaterials by blending with other polymers [8], is a subject of great interest and of intensive investigations due to the simplicity and effectiveness of mixing two different polymers to obtain new materials. Especially blending silk with cellulose (CE) is an important way to exploit for realizing new environmentally friendly artificial materials. The usefulness of cellulose is due to its abundance, low cost and peculiar properties. In the textile field cellulose, for instance, is widely used
∗ Corresponding author. Tel.: +39 0103538727; fax: +39 0103538727. E-mail address: [email protected] (E. Marsano). 0141-8130/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ijbiomac.2008.03.009
thanks to its superior moisture absorbing property and biodegradability [9]. Nowadays in modern industry blends based on synthetic polymers are widely used in order to develop new polymeric manufacts, able to yield property profiles superior to those of the individual components. This method is usually cheaper and less time-consuming for the creation of a novel polymeric materials with new properties, which may depend on different factors, such as molecular structure, miscibility or compatibility between the polymers, blend composition, morphology, and processing conditions [10–12]. A great advantage of polymer blends is that the materials’ properties can be tailored by combining component polymers and changing blend composition. The blending process must be performed in a fluid state that can be reached at temperature higher than the melting or glass transition temperature for semicrystalline or amorphous polymers. When polymers decompose before they undergo melt flow, the blending must be performed by the dissolution of the polymers in a common solvents followed by elimination of the solvent through evaporation or coagulation in a non-solvent. In the case of natural polymer as silk or cellulose the melting temperature is higher than the decomposition one; hence for these polymers it is necessary to find a solvent able to dissolve both. Studies regarding the phase behaviour of blend materials, often in form of film, based on SF and natural or synthetic polymers such as: cellulose [13–15], chitosan [16,17], sodium polyglutamate [18], sodium alginate [19], poly(vinyl alcohol) [20–22], nylon [23]
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have been reported. Most samples showed a good compatibility and an improvement of the mechanical properties [16–19,23]. Not only blend films based on SF were reported in the literature, but also fibers obtained by a wet spinning process blending SF with: cellulose [8,24,25], chitin [26], chitosan [27]. In our previous work [8] blend fibers composed by 70% CE and 30% SF were prepared by wet spinning solutions of the two polymers in DMAc-LiCl, using ethanol as coagulant. The fibers showed good mechanical properties, modulus and breaking strength in comparison to the pure CE fibers obtained using the same operative conditions. Once verified the spinnability of the dopes and the good compatibility between CE and SF, we studied the possibility to use an industrial, renewable and friendly solvent: Nmethylmorpholine-N-oxide (NMMO), a compound belonging to the organic cyclic amine oxide family, able to dissolve both SF [28] and CE [29]. Cellulosic fibers produced by using NMMO as solvent in the spinning process are commercialized with the trade name Lyocell® [30]. The process uses NMMO – or more generally a mixture of NMMO and water – to directly dissolve pulp without prior derivatization of the cellulose chains. This process is environmentally benign because the nontoxic NMMO solvent is used and almost all the solvent used is totally recycled [31]. Owing to these advantages, lots of technical developments of Lyocell® fibre have been reported by many manufacturers [32–35]. Lyocell® fibers have proved commercially successful because of better mechanical properties when compared with viscose rayon [36–39]. More recently reports on the swelling and dissolution of silk fibroin polymers in NMMO appeared [28,40,41] and regenerated SF were obtained by spinning these solutions at high polymer concentration [42,43]. In our previous papers [42,43] we described the results obtained by using a dry-jet wet spinning line, starting from SF–NMMO solutions. Regenerated SF fibers with good performance were obtained and it was observed that the tensile properties were strongly dependent on the draw ratio, which directly affects the molecular orientation of both amorphous and crystalline domains. The aim of the present work is to prepare regenerated CE/SF blend fibers, at different composition, by using NMMO/H2 O as solvent and to determine whether the polymers are compatible and if the resulting fibers have structural and physical-mechanical properties comparable to those of conventional regenerated cellulose fibers. 2. Experimental 2.1. Materials SF fibers were purified from silk cocoons degumming with water in autoclave at 120 ◦ C for 20 min, followed by extensive rinsing with warm distilled water to completely remove sericin, a gum protein surrounding SF filaments. The molecular weight of starting degummed SF fibers (Mv ) ranged from 220 to 240 kDa. To prepare SF for dissolution in NMMO, fibers were dissolved in saturated aqueous LiBr (9–10 M) at 60 ◦ C for 3 h. The solution was then filtrated to remove insoluble foreign matters, dialysed until complete removal of salt, and freeze-dried. The resulting SF in form of porous, sponge-like material was used for the dissolution tests. A commercial aqueous solution, NMMO/H2 O 50/50 (w/w), and the antioxidant n-propyl gallate (PG) were purchased from Aldrich Co. A regenerated Cellulose II sample having Mv = 63,000 (DP = 390) was supplied by Lenzing AG, Austria.
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2.2. Dope preparation Singles and mixtures of the two polymers were dispersed into the commercial aqueous solution of NMMO containing 0.7% (w/w) n-propyl gallate with respect to NMMO monohydrate. The dispersion was, then, concentrated under vacuum at 55 ◦ C to obtain NMMO hydrates with the desired water content (the maximum value was about 13.1% w/w). The water content was determined by Karl Fischer method. The overall concentration (Cp ) was fixed to 17 % (w/w) with respect to the solution. Blend solutions were prepared mixing silk and cellulose in the following ratio: 25/75, 50/50 and 75/25. In order to prepare a homogeneous solution in NMMO monohydrate, the dispersions were heated at 110 ◦ C for 20 min, under a nitrogen atmosphere, with a fast energetic stirring, controlling time by time the progress of dissolution with an optical microscope. Air bubbles were trapped in the solutions, due to the fast energetic stirring, and caused flow instability in the dry-jet wetspinning line. So before spinning the centrifugation of the dopes at high temperature was used in order to disengage air bubbles. After the centrifugation step, no phase separation was observed. 2.3. Spinning The dry-jet wet-spinning line is composed by an extrusion unit, a coagulation bath, ethanol, and two sets of spools, take up and roller. Take-up rate (V1 ) and roller rate (V2 ) can be varied so that fibers are collected with different drawn ratios. DR1 is defined as the ratio between V1 and Vo (DR1 = V1 /Vo ) where Vo is fiber’s rate at the spinneret hole. The post-spinning ratio, DR2 , is defined as the ratio between V2 and Vo (DR2 = V2 /Vo ). An extrusion rate Vo = 4 m/min, a 100 m spinneret and an air gap of 100 mm were used to extrude CE/SF-NMMO solutions. Collected fibers were further washed in ethanol for about 72 h in order to extract the residual NMMO and then they were dried and stored over CaCl2 until testing. 2.4. Fibre characterization 2.4.1. SEM analysis SEM characterisation was performed by using a Jeol JSM-6380LV microscope, at an acceleration voltage of 15 kV and 15 mm working distance. Films based on pure polymers and blends were fractured in liquid nitrogen and then mounted onto aluminium specimen stubs by means of double-sided adhesive tape and sputter-coated with a thin gold layer under rarefied Argon atmosphere, using a Polaron SC7620 Sputter Coater, with a current of 30 mA for 180 s. While the fibers were simply mounted onto aluminium specimen and coated with a thin gold layer. 2.4.2. Optical microscopy Fibers diameter were determined using an optical microscope. To measure the diameters, a small piece, 1 cm long adjacent to one side of each sample were observed. The average diameter was the result of at least six measurements on each adjacent piece. 2.4.3. Mechanical properties Mechanical properties of single fibers were measured using an Instron dynamometer mod. 5500, with a 50 mm gauge length at a crossbar rate of 30 m/min, corresponding to strain rate of 1.0 × 10−2 s−1 if elongation at break of specimen was up to 8%, and at a crossbar rate of 15 mm/min, relative strain rate of 6.7 × 10−3 s−1 if elongation at break was under 8% according to ASTM D3822-012004 standard method. Elastic modulus (E), stress at break ( b ), deformation to break (εb ) were calculated as the average of at least 15 measurements from stress–strain curves.
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2.4.4. Raman analysis FT-Raman spectra were recorded with 4 cm−1 resolution using a Bruker RFS100 instrument and NdYAG laser excitation line (1084 nm). 2.4.5. Thermal analysis Differential scanning calorimetry (DSC) was performed with a Mettler DSC 30 instrument, calibrated by an Indium standard. The calorimeter cell was flushed with 20 ml/min of N2 . Sample weight was about 3 mg and Al crucibles were used. The temperature range scanned was 25–500 ◦ C, at a heating rate of 10 ◦ C/min. Samples were pre-heated until 120 ◦ C, at 10 ◦ C/min, to remove moisture, then temperature was decreased to 25 ◦ C and the scan started. 2.4.6. X-ray data Wide-angle x-ray diffraction (WAXD) data were obtained at 20 ◦ C using a Siemens D-500 diffractometer equipped with a Siemens FK 60–10 2000 W tube (Cu K␣ radiation, = 0.154 nm). The operating voltage and current were 40 kV and 40 mA, respectively. The data were collected from 5 to 35 2 ◦ at 0.02 2 ◦ intervals, by using a silicon multi-cathode detector Vortex-EX (SII). The fibers were cut, mixed and pressed for a random panel to be submitted to WAXD measurements. The degree of crystallinity, Xc , was calculated from diffracted intensity data in the range 5–35 2 ◦ by using the area integration method [44]. 3. Results and discussion Before showing the results and discussing them, it is necessary to clarify the terminology that we use regarding miscibility and compatibility. As previously reported by Olabisi et al. [11], we use the term “miscibility” to describe polymer–polymer blends with behaviour similar to that expected of a single phase system. The term compatibility that has been used by many investigators involving various studies on polymer–polymer blend behaviour is a more general term. In a strict technological sense it can be used to describe whether a desired result occurs when two polymers are combined together and this can be correlated, for immiscible system, to good adhesion between two or more phases. For thermodynamic reasons most polymer pairs are immiscible and, as a result, phase separation could occur. Accordingly the dispersed phases could exist in different forms of spheres, ellipsoids, fibrils and lamellas, depending on the compositions, chemical, and physical structures of polymer blends, as well as processing conditions. Nevertheless, the degree of compatibility may vary widely. It is well known that the mechanical properties of immiscible blends are strongly dependent on their morphologies and, therefore, the control of the morphology of an immiscible polymer blend is extremely important for the tailoring of the final product properties [45,46]. Among different morphologies of the dispersed phase, the microfibrillar one is preferred because it can improve mechanical properties, as it is previously shown by different authors [47,48]. The morphology usually results from the complex history experienced by the different constituents during processing. During the wet-spun processing the final size, shape and distribution of the dispersed phase are determined by a wider variety of parameters such as the composition, miscibility in solution, coagulation rate and phase separation rate in the coagulation bath, shear rate/shear stress, interfacial tension among the polymers and also processing conditions, such as extrusion and take-up rate. In our previous work [8] blend fibers composed of 70% CE and 30% SF were prepared by wet spinning solutions in DMAc-LiCl and using ethanol as coagulant. The CE/SF fibers showed good mechanical properties, Young’s modulus, breaking strength and elongation
to break were quite similar, differences not more than 15% to those of CE fibers obtained in the same conditions were observed. Furthermore, small-angle x-rays analysis showed that CE/SF fibers were essentially amorphous and that there was a homogeneous dispersion of SF domains into a CE matrix. The cross-sectional size of the SF domain accounts for only few nanometres, equivalent to the dimension of 6–8 protein chains in an ordered conformation; the nano-dispersion of the two phases was the responsible of the good compatibility between the two polymers.
3.1. Influence of draw ratio on preparation of CE/SF fibers Fibers based on SF and CE blends obtained from NMMO solutions by using a laboratory dry-jet wet spinning line were prepared and characterized. All spinning processes are characterized by two drawing ratios: a coagulation ratio, DR1 , and a post-spinning ratio, DR2 . The influence of both was initially investigated by increasing the stretching of the filaments (DR1 and DR2 ) until a maximum value that can be used without the break up of the filaments. Our wet-spinning equipment allows a maximum DR1 = 6 and DR2 = 15 for an extrusion velocity (Vo ) equal to 4 m/min. The observed field of spinnability evidenced that CE could not undergo a second post-spinning stretching, thus DR2 was always equal to 1, while SF fibers could not be stretched more than three times in the first coagulation bath without observing the break-up of the filament at the extrusion unit. This behaviour could be due to different aspects: the molecular weight, the stiffness of the macromolecules, the viscoelasticity of the dope and the coagulation behaviour in solution. It is well know that cellulose and cellulose derivatives show mesophase formation in solution and that this correlates to the worm-like aspect of the macromolecules with persistence length ranging from 6 to 11 nm [49]. For instance, theoretical estimate for the persistence length of the cellulose chains in polar solvent, water, yields about 14 nm [50]. For this system a coagulation ratio higher than 1 is allowed, but once the filament is coagulated the high orientation of the stiff macromolecules does not allow further stretching. On the other hand, SF most probably has a coil conformation in NMMO solution and it maintains a coiled conformation also in the coagulated filament so that the fibers can be stretched in the post-spinning step. Another important factor in the stretching of fibers is the different coagulation behaviour of CE and SF filaments in ethanol. We observed that SF filaments after coagulation maintains a swollen behaviour with very high degree of elasticity; this can be correlated to the diffusion rate of ethanol into the SF/NMMO fibers, which was higher than the diffusion rate of NMMO into the ethanol coagulation bath. For coagulated CE filament we observed a slight decrease of the filament’s diameters in the coagulation bath and a formation of a stiff skin due to a high coagulation rate of CE in ethanol. The obtained filament is compact and cannot be further stretched. Another significant difference in the stretching behaviour can be surely correlated to the difference in the viscoelasticity of the dopes. Indeed, both systems present entanglements and they behave as synthetic polymers in a melt state. However, at the same concentration, the CE/NMMO system showed a higher viscosity with respect to the SF/NMMO one [51]. Moreover, we observed that all solutions based on CE/SF blends were not able to be stretched in the post-spinning step, even those with a very low content of cellulose, as show in Fig. 1. So fibers based on CE/SF blends were collected with different coagulation ratios (DR1 = 1–3–6), but a unitary post-spinning ratio (DR2 = 1).
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Fig. 1. The blocks show the field of spinnability of solutions of SF (position 1), CE (position 5) and their blends at different compositions (CE/SF 25–75 position 2; CE/SF 50/50 position 3; CE/SF 75/25 position 4) in function of the draw ratios DR1 and DR2 .
3.2. Raman Vibrational spectra are especially sensitive to the geometry of molecules and to the system of intramolecular and intermolecular interactions. For this reason FTIR and Raman have often been used as a useful tool to determine the nature of mixing between two polymers. The presence of a peak at a definite wavenumber would indicate the presence of a specific chemical bond, if interactions took place between the two polymers, the most obvious and significant difference would be the appearance of new peaks or shift of existing peaks. Fig. 2 shows the Raman spectra of CE, SF and CE/SF blend fibers collected applying only a draw ratio during coagulation (DR2 = 1). CE Raman spectrum presents a peak at 2894 cm−1 related to the stretching of the methylene group, while in the finger print region different peaks assigned to the structure of cellulose II can be detected: 1465 cm−1 due to bending COH, 1375 cm−1 and 1266 cm−1 due to bending of the methylene groups in the plane and out of the plane, respectively. Moreover, symmetric and asymmetric COC stretching of the glycoside groups are evident at 1114 cm−1 and 1097 cm−1 [52]. All these peaks clarify the presence of cellulosic chains in the structure of cellulose II [53]. On the other hand SF Raman spectrum, reported in Fig. 2, presents a peak at 3062 cm−1 due to the CH stretching in the C CH group, and asymmetric and symmetric stretching of methyl groups at 2985 cm−1 and 2935 cm−1 , respectively. In Fig. 2 SF Raman spectrum shows different typical peaks of silk fibroin. In particular Amide I vibrations indicative of antiparallel -sheet of silk occurred at 1663 cm−1 , while at 1229 cm−1 is evident the amide III, as well as CH2 bending scissoring and the CC stretching of the structure -sheet can be observed respectively at 1448 cm−1 and 1084 cm−1 [54–56]. The shape and intensity of all these bands can be associated with a -sheet molecular structure. In any of CE/SF blend fibers spectra neither bands shift superior to instrument resolution (4 cm−1 ), nor the appearance of new peaks occurred. As a consequence, no specific interaction between silk and cellulose could be detected. Therefore, the presence of a phase
Fig. 2. Raman spectra of the pure and blends fibers collected applying only a draw ratio during coagulation (DR1 = 3; DR2 = 1): (a) CE/SF 100/0; (b) CE/SF 75/25; (c) CE/SF 50/50; (d) CE/SF 25/75; (e) CE/SF 0/100. Above the 3200–2700 cm−1 region Raman spectra are shown, while below the finger print regions are reported.
separation in the fibers based on CE/SF blends at every prepared composition can be hypothesized. 3.3. Wide-angle X-ray diffraction characterization The typical WAXD pattern of CF/SF samples collected with drawn ratios DR1 equal to 3 are shown in Fig. 3. The profile of the CF/SF 100/0 sample is typical of regenerated cellulose fibers spun by organic solution. The reflection peaks can be indexed referring to the known cellulose II crystal structure [57,58]. The X-ray diffractograms of regenerated SF fibers shows the peak at 2 ∼ = 21◦ characteristic of the -sheet crystals of silk fibroin [59]. Intermediate profiles were collected for the CF/SF fibers prepared with different blend ratios. Analogous trend was observed for the fibers prepared with a value of DR1 equal 6. As reported in Table 1 the overall crystallinity (Xc ) decreases with the increase of the silk component in the blend. Slightly higher Xc values were calculated for pure CE and CF/SF blend fibers collected with DR1 equal to 6, even if the differences between the values are included in the error limits of the calculation method and thus they must be considered not significant. In Fig. 4, the experimental WAXD profiles of the blend fibers (DR1 = 3) are compared to the smoothed profile constructed by summing together the relative fraction contributions of the experimental pure components, CF/SF 0/100 and CF/SF 100/0. A good fitting was observed for all compositions confirming the presence
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Fig. 3. WAXD profiles of sample fibers at the different compositions collected with DR1 = 3 and DR2 = 1. (a) CE/SF 100/0; (b) CE/SF 75/25; (c) CE/SF 50/50; (d) CE/SF 25/75; (e) CE/SF 0/100. Table 1 Crystallinity, Xc , of pure and blends fibers prepared with different draw ratios, DR1 Samples Pure CE fibre
Blend fibre
Polymer weight ratio (CE/SF)
DR1
Xc (%)
100/0
3 6
32 36
75/25
3 6 3 6 3 6
24 26 23 25 21 22
1 3
21 21
50/50 25/75
Pure SF fibre
0/100
in the blend fibers of both components separately crystallized, confirming the hypothesis done previously. 3.4. Thermal analysis The properties of CE/SF blend fibers were further investigated by thermal analysis. Fig. 5 shows the DSC profiles of pure and blend fibers wet-spun from NMMO solutions and collected applying only a draw ratio during coagulation (DR2 = 1). The DSC thermogram of
Fig. 5. DSC curves of regenerated CE (a), SF (e) and CE/SF blend fibers with different blend compositions: (c) CE/SF 75/25; (d) CE/SF 50/50; (e) CE/SF 25/75. All the samples are collected applying only a draw ratio during coagulation (DR1 = 3; DR2 = 1).
CE fibers showed a weak and progressive endo shift until about 200 ◦ C. Heating at above 200 ◦ C caused a change of the trend from endo to exo. It has been reported that the thermal transitions occurring in this temperature range are associated with the onset of thermal degradation of cellulose, which consists of a sequence of events starting with rearrangement of chains in the amorphous regions, followed by main chain fission by cleavage of glucosidic bonds, volatilization of pyrolysis products (mainly levoglucosan), and dehydration, accompanied by increase of crystallinity and intermolecular cross-linking [60–62]. Beyond 300 ◦ C fast thermal degradation with extensive weight loss has been reported to occur [60]. Pure SF fibers showed an intense endothermic peak at 293 ◦ C representing the thermal decomposition of NMMO-regenerated SF fibers with oriented -sheet structure [42]. This peak clearly appeared also in the blend fibers (SF content ≥50 wt%), falling at a temperature very close to that of pure SF fibers (290.2 ◦ C and 288.4 ◦ C for the blends containing 75% and 50% SF, respectively). These features seem to indicate that the thermal behaviour of SF was only slightly influenced by blending with cellulose. With reference to WADX results we may suggest that well organized and crystalline SF domains were present in the blend fibers, especially at high SF content, and responded to the thermal stresses almost independently of the presence of cellulose. The situation was significantly different when the SF content decreased to 25 wt%. This fiber sample showed a CE-like profile with only a small SF contribution falling at 281 ◦ C in form of weak and broad endothermic peak. The contribution of CE to the DSC thermograms of blend fibers was identifiable as a high temperature shoulder of the main decomposition peak of SF. Interestingly; the shoulder resulted in an exothermic peak falling at about 350 ◦ C in the 50 wt% fiber blends, where it appeared with the highest intensity. Since this peak is present only in the blend fibers it can be attributed to some chemical and/or physical interactions between the degradation products of SF and CE, both of which are undergoing very fast and extensive thermal degradation in this high temperature range which leads to almost complete pyrolysis of the polymers. For this reason any interpretation of this thermal event can be hardly suggested. 3.5. SEM morphology
Fig. 4. Experimental and constructed WAXD profiles of blend fibers (DR1 = 3): (a) CF/SF 75/25; (b) CF/SF 50/50; (c) CF/SF 25/75.
The miscibility of CE and SF in the blends was systematically investigated by analysing the fracture morphology of films pre-
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Fig. 6. SEM cross-section morphology of SF, CE, and CE/SF blend films. (a) CE/SF 0/100; (b) CE/SF 100/0; (c and f) SF/CE 75/25; (d and g) SF/CE 50/50; (e and h) SF/CE 25/75. Mag. bars: (a–e) 10 m; (f–h) 2 m.
pared from the spinning dopes by smearing the polymers-NMMO hydrates onto a glass plate, followed by coagulation and solvent removal in an ethanol bath. As shown in Fig. 6, the fracture morphology of SF films was smooth and glassy (Fig. 6a) while that of CE films showed a granular and colloidal appearance (Fig. 6b). CE/SF blends displayed a gradual transition from one morphology to the other as a function of the blending ratio (Fig. 6c–e). Evidences of phase separation were detected in all the blends, irrespective of blend composition (Fig. 6f–h). Morphologically different domains randomly appeared in the fractured surfaces as areas with prevalently smooth or granular morphology sharply separated one from the other. These results correlate with the above spectroscopic and thermal analysis, and can be taken as indirect evidence that the two polymers tend to form separate domains when mixed in NMMO hydrates solvent. Regenerated SF, CE and CE/SF fibers were also characterized by SEM (Fig. 7). The whole fibers showed a circular shape with fairly regular size along the fiber axis. The surface morphology changed from granular to smooth for pure SF and CE fibers, respectively. Interestingly, a fine longitudinal striation was evidenced by CE/SF blend fiber. This feature appeared in all blends and was almost independent on the draw ratio applied. The morphological analysis of fiber cross-sections did not provide any clear evidence of phase separation observed in films, although the presence of cross-sectional features with a prevalently fibrous or granular structure was observed (pictures not shown).
3.6. Mechanical properties Generally mechanical properties, above all stress and elongation at break, are very important to evaluate fiber performances for proper application. Regenerated SF, CE and SF/CE blend fibers were tested for their cross-sectional dimension and tensile properties. The average mechanical properties were calculated from at least 10 curves for samples at different compositions and draw ratios, as it is shown in Table 2. Fig. 8 shows the diameters and the standard deviation of CE, SF and blend fibers as a function of the draw ratio. It can be noticed that for all the fibers a decreasing of the diameters with respect the drawing and not meaningful differences between samples with different composition were observed.
3.7. Tensile properties of pure polymers Fig. 9 illustrates the typical stress–strain plots of CE and SF samples. A different behaviour can be noticed, elongation at break of CE fibers decreased from 15 to 7%, and breaking strength increased from 180 to 360 MPa as the draw ratio during the coagulation (DR1 ) increased. For SF fibers a change from a fragile fiber to a ductile one was observed with increasing DR1 , as indicated by the values of both elongation at break and the breaking strength, which increased from 2 to 10% and from 52 to 116 MPa, respectively.
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Fig. 7. SEM photographs of SF (a), CE (b), and SF/CE blend fibers with different blend composition: (c) SF/CE 75/25; (d) SF/CE 50/50; (e) SF/CE 25/75. Mag. bars: (a–e) 5 m. All the samples are collected applying only a draw ratio during coagulation (DR1 = 3; DR2 = 1).
The mechanical behaviour of regenerated SF fibers can be understood from the microstructural analysis. As previously shown by the WADX analysis the existence of -sheet microcrystallites in all regenerated SF fibers is evident and the degree of crystallinity is approximately comparable, independently on the draw ratio in the coagulation. The change in tensile properties could be correlated with the alignment of protein chains in the amorphous regions that depends on the ability of the chains to slip at entanglements [63], which, in turn, depends on two microstructural features: the force required for the slippage of the chains at an entanglement and the density of the entanglements. In this context, the ductile behaviour of regenerated SF fibers subjected to high stretching suggests a reduction in the number of chain entanglements, facilitating chain slippage and plastic behaviour. In CE fibers, which are made by a more rigid polymer with a lower entanglement density, the increase of the draw ratio allows
an increase of the orientation of the amorphous phase that causing an increase of modulus, breaking strength and fragility of the fibers. It should be noted that the mechanical properties of our pure cellulose fibers were found to be significantly lower than those reported for the conventional regenerated cellulose fibers or films [64,65]. This can be correlated to our laboratory-scale dry-jet wet spinning equipment. 3.8. Tensile properties of blend fibers Typical stress–strain curves of wet-spun CE/SF blend fibers with different SF content are shown in Fig. 10 where four representative profiles per sample are plotted. The presence of SF in the blends should affect the orientation and mechanical strength of the fibers because a rigid polymer, cellulose, is mixed with a random coil polymer SF. According to Flory theory [66] a phase separation occurs upon blending rigid and coil polymers, beyond a critical concentration, in the absence of strong
Table 2 Polymer compositions (polymer weight ratio), average mechanical properties (Young’s modulus E, tensile strength b and elongation at break εb ) and diameters of the fibers obtained by using Vo = 4 m/min and different draw ratios DR1 in the coagulation bath Samples
Polymer weight ratio (CE/SF)
DR1
Diameter (m)
E (GPa)
b (MPa)
εb (%)
100/0
1 3 6
34.3 ± 5.2 16.1 ± 2.5 14.2 ± 1.0
11.3 ± 0.8 23.0 ± 1.5 24.2 ± 1.9
186 ± 9 260 ± 39 364 ± 50
15 ± 3 5±1 7±2
75/25
1 3 6 1 3 6 1 3 6
41.9 20.4 16.3 38.2 23.2 16.4 32.3 23.2 17.1
1 3
53.6 ± 6.9 24.5 ± 1.1
Pure CE fibers
50/50 Blend fibers 25/75
Pure SF fibers
0/100
± ± ± ± ± ± ± ± ±
1.6 6.0 2.6 0.6 0.2 0.2 1.0 0.8 0.2
11.7 22.7 21.7 14.4 16.2 18.1 11.2 11.7 14.0
± ± ± ± ± ± ± ± ±
0.8 2.9 1.0 0.7 0.7 0.9 0.6 0.4 0.4
5.4 ± 0.5 7.9 ± 0.4
159 307 261 193 236 315 148 155 169
± ± ± ± ± ± ± ± ±
10 20 18 4 16 39 5 7 8
52 ± 6 116 ± 5
8 3 3 8 7 6 4 5 7
± ± ± ± ± ± ± ± ±
2 1 1 1 1 2 1 3 1
2±1 10 ± 5
E. Marsano et al. / International Journal of Biological Macromolecules 43 (2008) 106–114
Fig. 8. Diameters versus the draw ratio during coagulation, DR1 , of pure and blends fibers.
intermolecular interactions. The phenomenon is due to the interference of the random coil polymer with the mutual orientation of the molecules of the rod-like polymer. A qualitative agreement between theory and experimental results is reported in literature for system based on a rigid polymer and a flexible one [67,68]. The CE/SF 25/75 fibers showed behaviour similar to that of pure SF fibers. A slight increase of the elongation at break with the DR1 was observed (Table 2). Fibers composed of 50/50 CE/SF displayed a practically constant elongation at break irrespective of the draw ratio. Fibers composed of 75/25 CE/SF fibers showed a behaviour similar to pure CE, the fragility of the fibers increasing with the draw ratio. Fig. 11 shows the trend of Young’s modulus and stress at break for fibers obtained at DR1 = 3 with different blends compositions. The decrease in mechanical properties seems for all the samples to be less than the linear decreases with the amount of SF added. In particular the blend containing more than 75% of CE show higher modulus and strength than those of pure cellulose fiber, while these properties had lower values in the fiber with a CE/SF weight ratio of 50/50 and 25/75. In compatible blends, a linear relationship between blend composition and such mechanical properties as elastic modulus and strength at break is usually observed. In general, a negative deviation from the linear relationship is considered an indication of
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Fig. 10. Stress–strain plot of pure and blend fibers at different compositions obtained using a draw ratio during coagulation, DR1 = 3: (a) CE/SF 100/0; (b) CE/SF 75/25; (c) CE/SF 50/50; (d) CE/SF 25/75; (e) CE/SF 0/100.
Fig. 11. Variation of Young’s modulus () and stress at break () with CE content for CE/SF fibers at DR1 = 3.
poor compatibility between blend components, whereas a positive deviation is considered an indication of improved compatibility. The results reported in Fig. 11 are an indication of good compatibility between CE and SF fiber and put also in evidence the role of the DR in the formation of morphologies of the dispersed phase that can enhance the compatibility between the two polymers. The
Fig. 9. Stress–strain plot of the pure polymer fibers obtained using different draw ratios, DR1 , in the coagulation bath: (a) CE fibers; (b) SF fibers.
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results of the mechanical behaviours seem to indicate that with the increase of CE content in the fibers and for a moderate draw ratio (DR1 = 3) good interaction between the two phases occur. This increase of interactions could be due to the formation of morphologies that enhanced the amount of amine groups of SF that can be incorporated into the CE amorphous matrix and interacted with the OH groups in the fibers. 4. Conclusions CE/SF blends monofilaments, with diameters ranging between 14 and 41 m, were successfully prepared by extruding solutions in NMMO·H2 O in a coagulation bath containing ethanol, through a dry-jet wet spinning line. Blends fibers were obtained in different drawing conditions and the amount of SF did not influence the spinnability regions of solutions based on CE/SF respect to cellulose one. FTIR, DSC and WAXD analyses of CE/SF blends fibers evidenced that both polymers underwent crystallization upon solidification and two phase systems were obtained. The phase separation was observed in blends films with SEM; the microphotographs showed that the fracture of CE/SF blends films displayed a gradual transition with blend compositions, from a smooth and glassy morphology, typical of SF films, to a granular one characteristic of CE films. This behaviour was detected in all the blends, irrespective of composition and, therefore, the presence of a phase separation in the fibers at every prepared composition could be hypothesized. Respect to the linear behaviour of the main mechanical properties between the pure polymers, the tensile properties of the blends fibers showed higher values, although the two polymers were not miscible and showed phase separation in all samples. This is a clear indication of the existence of some kind of interactions between the two phases. In particular CE/SF blend fibers containing more than 75% of CE exhibit higher tensile properties than those of pure cellulose, obtained by using the same spinning condition (DR1 = 3). This result seems to indicate that good interactions between the two phases occur suggesting a good compatibility of these two natural polymers in the fibers. To this end, CE/SF blend monofilaments have great potential as a new regenerated biodegradable fiber with structural and physic-mechanical properties comparable to those of conventional regenerated cellulose fibers. Acknowledgments The authors wish to thank Mr. Mario Traverso for his helpful work on dry-jet wet spinning line. References [1] S. Arcidiacono, C.M. Mello, M. Butler, E.A. Welsh, J.W. Soares, A. Allen, D. Ziegler, T. Laue, S. Chase, Macromolecules 35 (2002) 1262–1266. [2] F. Buehler, V. Baron, E. Schmid, P. Meier, H.J. Schltze, USA USPatent 5516815 (1996). [3] R.D. Gilbert, X.C. Hu, R.E. Fornes, J. Appl. Polym. Sci. 58 (1995) 1365–1370. [4] A. Chiarini, P. Petrini, I. Bozzini, I. Dal Pra, U. Armato, Biomaterials 24 (2003) 789–799. [5] D.L. Kaplan, G.H. Altman, F. Diaz, C. Jakuba, T. Calabro, R.L. Horan, J. Chen, H. Lu, J. Richmond, Biomaterials 24 (2003) 401–416. [6] Z. Shao, F. Vollrath, R.J. Sirichaisit, R.J. Young, Polymer 40 (1999) 2493–2500. [7] Z. Shao, F. Vollrath, Nature 418 (2002) 741. [8] E. Marsano, M. Canetti, G. Conio, P. Corsini, G. Freddi, J. Appl. Polym. Sci. 104 (2007) 2187–2196. [9] C.H. Park, Y.K. Kang, S.S. Im, J. Appl. Polym. Sci. 94 (2004) 248–253. [10] D.R. Paul, S. Newman, Polymer Blends, Academic Press Inc., New York, 1978. [11] O. Olabisi, L.M. Robeson, M.T. Shaw, Polymer-Polymer Miscibility, Academic Press, New York, 1979.
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