Polyethylene Reinforced with Keratin Fibers Obtained from Chicken Feathers
Composites Science and Technology 65 (2005) 173–181 www. elsevier. com/locate/compscitech Polyethylene reinforced with keratin ? bers obtained from chicken feathersq Justin R. Barone *, Walter F. Schmidt USDA/ARS/ANRI/EQL, Bldg. 012, Rm.
1-3, BARC-West, 10300 Baltimore Ave. Beltsville, MD 20705, USA Received 15 January 2004; received in revised form 22 June 2004; accepted 22 June 2004 Available online 24 August 2004 Abstract Polyethylene-based composites were prepared using keratin ? bers obtained from chicken feathers.Fibers of similar diameter but varying aspect ratio were mixed into low-density polyethylene (LDPE) using a Brabender mixing head. From uniaxial tensile testing, an elastic modulus and yield stress increase of the composite over the virgin polymer was observed over a wide range of ? ber loading. Scanning electron microscopy revealed some interaction between the polymer and keratin feather ? ber.
In addition, the keratin ? ber had a density lower than the LDPE used in this study resulting in composite materials of reduced density.The results obtained from mechanical testing are compared to theoretical predictions based on a simple composite material micromechanical model. Published by Elsevier Ltd. Keywords: A. Fibers; A. Polymer-matrix composites; A.
Short-? ber composites; B. Mechanical properties; B. Microstructure 1. Introduction There has been recent interest in developing composites based on short-? bers obtained from agricultural resources.
These ? bers are usually of lower density than inorganic ? bers, environmentally-friendly, and relatively easy to obtain. It is anticipated that the ? ers would not contribute to the wear of polymer processing equipment and may not su? er from size reduction during processing, both of which occur when inorganic ? bers or ? llers are used. Although the absolute property increase when using organic ? bers is not anticipated to be nearly as high as when using inorganic ? bers, the speci? c proper- q Mention of trade names or commercial products in this article is solely for the purpose of providing speci? c information and does not imply recommendation or endorsement by the US Department of Agriculture. * Corresponding author. Fax: +1 301 504 5992. E-mail address: [email protected]
rs. usda. gov (J. R. Barone).
ties are anticipated to be high owing to the much lower density of the organic ? bers. In short-? ber reinforced polymer composites, the integrity of the ? ber/matrix interface needs to be high for e? cient load transfer. Ideally, the molten polymer would spread over and adhere to the ? ber, thus creating a strong adhesive bond. Inorganic ? bers like glass and cellulosic ? bers have hydrophilic surfaces that make them incompatible with hydrophobic polymers. Therefore, inorganic and cellulosic ? bers usually require chemical modi? cation to increase ? ber/polymer interactions .The chemical modi? cation, known as a coupling agent, acts as a ‘‘bridge’’ between the inorganic ? ber and the organic polymer matrix.
The ‘‘bridge’’ must adhere or bond to the ? ber and, in turn, strongly interact with the polymer. When using glass ? bers, the coupling agent has a hydrophilic side that is compatible with the ? ber and a hydrophobic side that is compatible with the polymer. In glass ? bers, the coupling agent reacts with the surface of the glass forming covalent bonds. Without the coupling agent, there is simply adhesion of the polymer to the glass through weak bonding, 0266-3538/$ – see front matter.Published by Elsevier Ltd. doi:10.
1016/j. compscitech. 2004. 06. 011 174 J.
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e. , van der Waals or induction interactions. Organic ? bers may o? er the possibility of covalently bonding the matrix polymer to the ? ber either directly or through a similar type of chemical ‘‘bridge’’ and the chemistry may be easier. Covalent bonds are much stronger than induction or van der Waals interactions so a covalently bonded interface would be advantageous . Most studies of naturally occurring organic ? bers concentrate on cellulose-based ? ers obtained from renewable plant resources such as wood [3–11], cotton , ? ax , sisal [14,15], jute , hemp , ramie , and bamboo . Lundquist et al.
 were able to get modulus increases of about 4 times and yield stress increases of about 2 times using cellulose ? bers of ca. 17 lm diameter and aspect (L/D) ratio of ca. 76. However, the increase required a deviation from tradition melt processing techniques. The polymer matrices used to make the composites can be synthetic or naturally-derived and thermoplastic or thermosetting. There is much work reported on making natural ? er reinforced composites from epoxidized soybean oil or soy based polyurethane [12,20].
Not only are those matrices naturally-derived but processing can occur at low temperature so as to avoid ? ber degradation. Synthetic thermosetting epoxies with ? ax ? bers  and polyurethanes with sisal ? bers  that can be molded at low temperature have also been used. These techniques carefully avoid processing at temperatures similar to the degradation temperature of the ? bers. Conventional synthetic thermoplastic polymers, like polyethylene (PE) and polypropylene (PP) have also been used to make natural ? ber composites.Colom et al.  prepare HDPE/wood ? ber composites using a compounding step at 160 °C in a roll mill and a molding step at 150 °C in a compression molder for up to 20 min.
This is a more traditional polymer composite processing method. Colom et al. are only able to get property increases from cellulose ? bers after treatment with a silane coupling agent to increase ? ber/polymer interactions. These researchers observe modulus increases of about 2 times and yield stress increases of about 1. 6 times in HDPE composites over a ? ber loading range of 10–40 weight percent. However, the aspect ratio of the ? bers is short at L/D $ 9.
Eboatu et al.  compound oil palm particles in PP on an extruder at 190 °C. Kuan et al.  incorporate wood ? ber into HDPE on a laboratory extruder while Sameni et al.
 mix wood ? ber into PP on a Brabender mixing head. Ali et al.  prepare sisal ? ber composites on a Haake mixing head and show a property increase with increasing ? ber aspect ratio at constant 20 wt% loadings. The sisal ? bers have diameters of 40–150 lm and aspect ratios of 50–135. There are some studies detailing the incorporation of organic ? bers in plasticized polymers.
Oksman et al.  mix ? ax ? bers into plasticized polylactic acid PLA) to obtain composites with increased mechanical properties over the plasticized PLA alone. At room temperature, PLA is in a glassy state and plasticizing lowers the glass transition temperature to below room temperature, giving plasticized PLA thermal and mechanical properties similar to PP. The composites are compounded on an extruder at temperatures of about 180 °C, which is a typical temperature for polyole? n processing.
Jana and Prieto  plasticize polyphenylene ether (PPE) to reduce the glass transition temperature. PPE is processed at around 260 °C, which is well above its glass transition of 212 °C.Plasticizing allows for processing at 200– 220 °C with wood ? bers so as to minimize ? ber degradation. In the case of Oksman et al. and Jana and Prieto, the composite of plasticized polymer and cellulose ? ber has increased properties over the plasticized polymer alone. However, the modulus and tensile strength of the plasticized polymer/cellulose ? ber composite are comparable to that of the virgin, unplasticized polymer.
The advantage of plasticization then is to lower the processing temperature and to increase ‘‘toughness’’ as manifested in a larger elongation to break and increased impact properties.There are a few studies detailing composites made from protein ? bers obtained from agricultural resources. Madera-Santana et al.  have prepared leather short? ber reinforced PVC composites. Madera-Santana et al.
show yield stress increases of up to 5 times and modulus increases of up to 150 times when using leather ? bers in PVC but only after chemical treatment of the ? bers. Much smaller property increases are observed at high ? ber loadings without chemical modi? cation. The leather ? bers have small diameter, i. e.
, several microns, and aspect ratios of ca. 50.There does not appear to be a shortage of possible sources of naturally occurring ? bers that could be harvested for use in composite materials. One accessible ? ber resource is the over four billion pounds of chicken feather waste generated by the US poultry industry each year . The feathers are made of keratin, which contains ordered a-helix or b-sheet structures and some disordered structures. The feather ? ber fraction has slightly more a-helix over b-sheet structure.
The clear outer quill has much more b-sheet than a-helix structure . This leads to a crystalline melting point of ca. 30 °C for the outer quill compared to ca. 240 °C for the feather ? ber . Feather keratin has a molecular weight of about 10,500 g/mol  and a cysteine/cystine content of 7% in the amino acid sequence . Cysteine is a sulfur containing amino acid responsible for the sulfur–sulfur bonding in the keratin.
The mechanical properties of whole feathers and feather rachis (i. e. , the quill or shaft of the feather) have been measured for several species of birds. Purslow and Vincent  measure the elastic modulus of primary feather shafts from pigeons, with and without the medulla (inner quill) using a bending test.The modulus val- J. R.
Barone, W. F. Schmidt / Composites Science and Technology 65 (2005) 173–181 175 ues obtained are 7. 75–10 GPa on dehydrated feather shafts.
The values are lower after removal of the medulla. Fraser and MacRae report modulus and peak stress values of Laysan albatross feather at 65% relative humidity as 5. 2 GPa and 200 MPa, respectively . The values are lower at 100% humidity, with modulus being 3. 4 GPa and peak stress 100 MPa.
Bonser and Purslow  use tensile testing to obtain the YoungOs modulus of the feather quill or shaft from a variety of birds. All of the values are about 2. GPa at room temperature and humidity. Cameron et al.
show that the modulus of feather rachis is higher in birds capable of ? ight than it is in terrestrial birds . From tensile tests, the modulus of swan and goose feather rachis is 2–4 times higher than for ostrich feather rachis. In addition, these researchers show that the modulus increases along the length of the feather shaft, with the lowest value near the skin and the highest value at the tip. X-ray di? raction measurements show more keratin molecule orientation further out along the rachis, which is the origin of the higher measured moduli.Swan and goose feather rachis have modulus values of 2. 5–5 GPa while ostrich has a modulus value of about 1.
5 GPa at 50% relative humidity. Much less is known about the physical properties of the ? ber portion of the feather. While there is much data on the quill portion, the quill is predominantly b-sheet protein structure and may di? er in properties from the ? ber . Recently, George et al.
 reported the mechanical properties for turkey feather ? ber. Turkey feather ? ber varies in dimension and properties depending on the position on the feather shaft.Fiber located closer to the bird is smaller in diameter and has lower physical properties than ? ber from further out on the rachis. So the ? ber shows the same trend in physical properties as the rachis. George et al. show that the ? bers closer to the bird have an average denier of 55.
2 g/ 9000 m, average tenacity at break of 0. 36 g/denier, average strain at break of 16. 43%, and average modulus of 4. 47 g/denier.
The ? bers further out on the shaft have values of 142 g/9000 m, 0. 83 g/denier, 7. 96%, and 15. 55 g/denier, respectively.
If the density of the turkey feather ? er is known, the denier values could be converted to ? ber diameters and then the physical property data converted to stress units. In addition, George et al. note that the ? bers closer to the bird are not straight but branched and this may a? ect the reported denier values by making them arti? cially high. Turkey feather ? ber is much larger in diameter than poultry feather ? ber.
Also, mechanical properties are measured on ? ber bundles, not individual ? bers, which may a? ect results. Composite materials have been prepared from poultry feather ? ber. Hamoush and El-Hawary prepared feather? er reinforced concrete at 1–3 vol% of ? ber and found that the workability of the mix is low compared to ordinary cement and more plasticizer is needed . The physical properties of the feather ? ber concrete are lower than ordinary concrete because of the large amount of plasticizer used. The origin of the low workability may lie in a viscosity modi? cation due to ? ber addition as well as the absorption of water by the protein ? ber.
Concrete is typically 15–20 vol% water  so the ? ber may be absorbing water and the authors perhaps should have pursued the addition of water rather than plasticizer to increase workability.Water is lost when the concrete dries and physical properties may have been maintained. Bullions et al. [41,42] prepare composites of kenaf bast, wood pulp, and poultry feather ? ber with polypropylene using a wetlay process.
For feather ? ber/PP composites, modifying PP with maleic anhydride to increase ? ber/ polymer interactions increases the physical properties of the composites. Dweib et al.  use a vacuum-assisted resin transfer molding process to mold recycled paper/feather ? ber composites from epoxidized soybean oil. The composites are processed at room temperature to avoid ? ber degradation. Schuster ompounds feather ? ber in polypropylene on an extruder at 200 °C  and observes increased hardness and heat distortion temperature. Concurrently, tensile stress and modulus are maintained but impact strength and ultimate elongation decrease with the addition of ? ber.
Schuster uses 2% bismaleinanhydride as a coupling agent. Schuster also separates out the straight ? bers from the branched ? bers and notices no di? erence in composite solid-state properties but claims the branched ? bers are more di? cult to process in the melt state. In this paper, keratin feather ? ber ranging in length from 0. 0053 to 0. cm is incorporated into low-density polyethylene at percentages of 0–50% wt.
The composites are mixed in a Brabender mixing head. Following mixing, tensile bars are prepared and tested in uniaxial tension to assess elastic modulus, yield stress, and yield strain. Scanning electron micrographs of the fracture surfaces denote ? ber/polymer interactions and ? ber orientation. 2.
Experimental 2. 1. Keratin feather ? ber Keratin feather ? ber is obtained from Feather? berO Corporation (Nixa, MO). The keratin feather ? ber is cleaned and separated from the quill fraction according to a process developed and patented by the USDA .
Currently, Feather? berO Corporation is the only known commercial supplier of feather ? ber material. While there is much raw feather material available from poultry processors, the feathers would have to be cleaned and separated to obtain the ? ber. The possibility of different ? bers possessing di? erent properties has been alluded to previously. So it may be possible to tailor 176 J. R.
Barone, W. F. Schmidt / Composites Science and Technology 65 (2005) 173–181 composite properties by separating out ? bers from different parts of the feather or from di? erent feathers. In practice, the ? ers used are a combination of all of the ? ber obtained from the poultry feathers, which is a more practical way to produce composites from poultry feather ? ber. The feather ? ber is semi-crystalline and has a constant diameter of approximately 5 lm.
The density of feather ? ber is determined by displacing a known volume and weight of ethanol with an equivalent amount of ? ber. A density value of 0. 89 g/cm3 is obtained. The ? ber lengths received range in length from ca. 0. 32 to 1.
3 cm so uniform lengths are obtained through grinding and sieving the as-received fraction.Fibers of 0. 02, 0. 1, and 0. 2 cm lengths are made by grinding feather ? ber using a Retsch ZM 1000 centrifugal grinder. The rotational velocity of the instrument is 15,000 rpm and contains a torque feedback so as to not feed in too much material and overload the motor.
The ? ber is fed in slowly to avoid motor overload and to minimize frictional heating of the instrument and the ? ber. Smaller ? ber lengths are made by grinding the ? ber on a Retsch PM 400 ball mill. Feather ? ber is loaded into 500 ml stainless steel grinding vessels so that it occupies about a quarter of the volume.The grinding media are four 4 cm stainless steel spheres for a total of 1132 g grinding media.
Grinding proceeds at 200 rpm for 30 min. Each ground fraction is sieved on a Retsch AS 2000 vibratory mill. For the longer ? ber lengths, 1 cm diameter glass beads are used as sieving aids to aid the separation process. Sieving occurs at a constant frequency but amplitude and time can be varied.
The material is loaded into the top sieve of the stack. The sieving stack contains eight sieves with hole sizes from 0. 0710 to 0. 0038 cm.
Sieving at an amplitude of 1. 0 (arbitrary instrument scale) for 60 min e? ctively separates the ‘‘? nes’’ from the desired average ? ber length. 2. 2.
Composite preparation The matrix material is a low-density polyethylene (LDPE) commercially available from Dow called LD133A. The LDPE has a melt ? ow index (MFI) of 0. 22 g/10 min at 190 °C and 2. 16 kg and a density of 0.
92 g/cm3 in the solid-state. The melting temperature and percent crystallinity of LD133A is determined from di? erential scanning calorimetry according to ASTM D3417 and ASTM D3418 using a TA Instruments DSC 910S. The melting temperature is 112 °C and the percent crystallinity is 45%.The total sample weight of each composite is 35 g. Composites are prepared by ? rst adding the LDPE into a Brabender mixing head set at 150 °C and rotating at 50 rpm.
Immediately after adding the LDPE, the respective amount of feather ? ber is added into the mixing head. Feather ? ber loadings range from 0 to 50 wt%. Higher loadings take approximately 2 min to load into the mixing head. The melt temperature is monitored independently and ranges from 171 °C for the pure polymer to 183 °C for the highest ? ber loading. The total mixing time for each composite is 15 min. Following mixing, each sample is sandwiched between Te? n-coated aluminum foil and pressed into three thin sheets in a Carver Press Autofour/30 Model 4394 at 160 °C, 133446 N for 18 s.
The ? lm is then removed and cooled under an aluminum block until it reaches room temperature. After pressing, each thin ? lm is inspected to note feather ? ber dispersion. Good dispersion is observed in all cases except the 50 wt% loading, which appears to be overloaded as evidenced by some agglomeration of ? bers. To examine the e? ect of mixing, LDPE samples are prepared without mixing.
No di? erence in physical properties is observed between mixed and unmixed LDPE after testing.To prepare samples for testing, the three thin sheets are cut into quarters, stacked on top of each other, sandwiched between Te? on-coated aluminum foil, and pressed in the Carver Press at 160 °C and 8896 N for 2 min. After pressing, the ? lms are cooled under an aluminum block until they reach room temperature. This results in ? lms approximately 0. 3 cm in thickness.
Type IV dogbone samples for testing according to ASTM D638 are machined from the ? lms. 2. 3. Composite testing Composite samples are allowed to sit at ambient conditions for one week before testing.
Uniaxial tensile testing is performed using a Com-Ten Industries 95 RC Test System. Test speeds, v, from 2. 5 to 22. 9 cm/min (1 to 9 in. /min) are used. Physical properties have a weak dependence on test speed, i.
e. , (log E) $ (log v)0. 09 and (log ry) $ (log v)0. 03. The data for the applied testing speed of 12.
7 cm/min (5 in. /min) is reported. A minimum of three samples of each composite is tested. Elastic modulus, E, is de? ned as the initial linear portion of the stress-strain curve, after correcting for the ‘‘toe’’ region. This portion of the curve is ? t to a ? rst order polynomial and E is obtained from the slope.
The yield stress, ry, and yield strain, ey, are de? ned as the stress and strain at the ‘‘peak’’ of the stress-strain curve. The composites break at this point and the polymer control samples yield signi? cantly. 2. 4. Microscopy The fracture surfaces are excised from the failed tensile bars using a scalpel blade and transferred into a modi? ed specimen carrier. The specimen carrier is known as an ‘‘indium vise’’ because the dissected pieces J.
R. Barone, W. F. Schmidt / Composites Science and Technology 65 (2005) 173–181 177 are clamped between sheets of indium metal and plunge cooled in liquid nitrogen to A196 °C.
The cooled holder is then transferred to an Oxford CT1500 HF cryo-preparation system attached to a Hitachi S-4100 scanning electron microscope (SEM). The sample temperature is raised to A90 °C for 10 min to remove surface water from the sample surface. The sample is then cooled to below A120 °C and coated with approximately 5 nm of platinum metal using a magnetron sputter coater. Coated samples are transferred to the cold stage in the SEM at A170 °C and observed with an electron beam accelerating voltage of 2 KV. 3.
Results Fig. 1 shows the elastic modulus, Ec, and the speci? c modulus, Ec/qc, of the composites made from 0. cm long ? bers. The composite density, qc, is determined from the equation ! A1 wf wp ? ; ? 1? qc ? qf qp where w is weight fraction, f denotes ? ber and p denotes polymer .
Eq. (1) assumes there is no void volume in the composite from processing. Independent measurement of the composite density shows a slightly higher density than that predicted by Eq. (1). This could be the result of compression molding the composites to obtain tensile bars or of experimental error in the measurements.
The modulus is normalized by the composite density with units of kg/m3. The weight fraction of ? ber s converted to volume fraction using the equation /f = (qcA qp)/(qfAqp). Fig. 2 shows the yield stress, ry, and yield strain, ey, behavior for the composites made from 0.
1 cm ? ber lengths (? ber aspect ratio, L/D = 200) over the entire feather ? ber loading range. Yield stress noticeably increases as feather ? ber loading increases. The yield strain decreases as ? ber loading increases. Fig. 3 shows the e? ect of ? ber aspect ratio on elastic modulus and yield stress.
All of the data are from 30. 0 open symbols: ? y filled symbols: ? y 1. 0 Yield Stress, ? y (MPa) 20. 0 0. 5 10. 0 0.
0 0. 0 0. 1 0. 2 0.
3 0. 4 0. 5 0. 0.
6 Volume Fraction Keratin Feather Fiber, ? f Fig. 2. Composite yield stress and strain versus volume fraction of 0. 1 cm keratin feather ? ber. 0.
15 25 0. 2 0. 0002 20 Specific Modulus, Ec/? c (*10 m /s ) Elastic Modulus, Ec(GPa) Elastic Modulus, Ec (GPa) 0. 10 15 0.
1 0. 0001 10 0. 05 open symbols: Ec filled symbols: ? y 5 open symbols: Ec filled symbols: Ec/? c 0. 0 0. 0 0. 1 0.
2 0. 3 0. 4 0. 5 Volume Fraction Keratin Feather Fiber, ? f 0.
0000 0. 6 0. 00 0 100 200 300 400 0 500 Fiber Aspect Ratio, (L/D)f Fig. 3. Composite elastic modulus and yield stress versus ? ber aspect ratio at constant 20 wt% ? ber loading.Fig.
1. Composite elastic modulus and speci? c modulus versus volume fraction of 0. 1 cm keratin feather ? ber. Yield Stress, ? y (MPa) 9 2 2 Yield Strain, ? y (cm/cm) 178 J.
R. Barone, W. F. Schmidt / Composites Science and Technology 65 (2005) 173–181 composites loaded to 20 wt% (20. 6 vol%) feather ? ber.
There is an indication that modulus and yield stress increase as ? ber aspect ratio increases and that beyond a critical ? ber aspect ratio of approximately 50 the physical properties are unchanged. 4. Discussion The results show that reinforcement of a polymer matrix can be achieved with keratin feather ? ber.In Fig. 1, there is an observed increase in elastic modulus of almost 3 times over the keratin feather ? ber loading range. Fig.
2 shows that the yield stress increases by a factor of 2. 5 over a ? ber volume fraction range of 0–0. 51. A micromechanical approach to predict the strength of composites is rc ? k/f rf ? /p rp ; ? 2? where k is a term sometimes referred to as the ‘‘stress e? ciency factor’’ . The term k is a function of the ? ber/matrix adhesion, which governs the transfer of stress from the polymer matrix to the ? ber, ? ber orientation relative to loading direction, ? ber shape, and ? ber aspect ratio.It is assumed that all of the ? bers are of the same circular cross-section and constant aspect ratio of L/ D = 200, which, from Fig.
3, indicates that L > Lc or the ? ber is su? ciently long to maximize ? ber loading. Table 1 shows the k values as a function of ? ber volume fraction as determined from Eq. (2). The value of rp, determined experimentally, is 10. 24 MPa. The rf = 200 MPa data of Fraser and MacRae is also used.
One factor that could a? ect the micromechanical analysis is the lack of physical property data for the keratin feather ? bers used. However, any deviation from the literature values would be re? cted in the adjustable parameters in the micromechanical analysis. The k values obtained are not zero or negative but are not 1, which would indicate perfect adhesion and all ? bers oriented in the deformation direction . The low k values indicate that all of the ? bers are not oriented in the loading direction and/ or ? ber/polymer interactions are not maximized. To further investigate the orientation of the ? bers and the ? ber/polymer interactions, the fracture surfaces Fig.
4. (a) 10 wt% and (b) 40 wt% 0. 1 cm keratin feather ? ber in LD133A LDPE. Scale bar is 300 lm. Table 1 Stress e? iency factor, k, for LD133A composites with L/D = 200 feather ? ber /f (%) 0 10.
3 20. 6 30. 8 40. 9 50. 9 k 0 0. 153 0.
133 0. 161 0. 135 0. 159 of the tensile bars are imaged using SEM.
Figs. 4(a) and (b) are micrographs of the fracture surfaces at 10 and 40 wt% (10. 3 and 40. 9 vol%), respectively. The loading direction 1 (tensile bar gage length direction) is into the paper.
The transverse directions 2 (width of tensile bar) and 3 (thickness of tensile bar) are marked on the micrographs. There is a fair amount of ? bers oriented in the transverse direction 2. Voids and ? bers in direction 1 also are evident. Observation of Fig.
(b) shows that the voids are about the diameter of the ? bers so the voids may represent volumes once occupied by ? bers. It would seem that if the voids are the result of processing anomalies, then the voids would have a wider size distribution. The ? ber length is 0. 1 cm which is 1/3 the thickness value so not much orientation in the 3 direction would be expected. Fiber length reduction is di? cult to assess because the ? bers are only partially exposed. The ? bers are well dispersed, which is important to obtain good physical properties.
Figs. 5(a) and (b) show the same fracture surfaces as in Fig. 4 but at a higher magni? cation.There is some ? ber/ J. R.
Barone, W. F. Schmidt / Composites Science and Technology 65 (2005) 173–181 179 Fig. 5. (a) 10 wt% and (b) 40 wt% 0.
1 cm keratin feather ? ber in LD133A LDPE. Scale bar is 30 lm. polymer interaction as shown by the matrix adhering to the ? bers to some degree. Some matrix deformation occurs along with the ? bers as the ? bers are pulled. There is some ? ber pullout as noticed by the voids left and the exposed ? bers. Some of the ? bers are fractured in the same fracture plane as the polymer, which would indicate strong ? ber/polymer interactions.
This is shown more clearly in Fig. 6, which is a high magni? ation of the 40 Fig. 6. 40 wt% 0. 1 cm keratin feather ? ber in LD133A LDPE. Scale bar is 6 lm.
wt% composite fracture surface. Here, the ? ber is wetted by the polymer and the ? ber fragment length is not very long, indicative of good adhesion [26,27]. Also shown in Fig. 2 is the onset strain for yielding, ey.
The yield strain trend is indicative of a transition of the material behavior from ductile to brittle. Referring to Fig. 4, at 10 wt%, the fracture topography is ductile, with localized drawing of the polymer. At higher ? ber loadings, the fracture topography becomes ? atter with less localized polymer drawing.
The amino acid sequence of feather keratin shows that the protein has 40% hydrophilic and 60% hydrophobic groups . The keratin feather ? ber should be compatible with hydrophobic polymers to some degree. The intrinsic surface roughness of the ? bers increases the surface area by a factor of about 2. 2 over a perfectly smooth ? ber. The extra surface area factor was estimated from length scale observations from SEM micrographs of individual ? bers. The possibility of strong chemical compatibility and lots of available ? ber surface area may increase the ? ber/surface interactions over smooth inorganic ? bers or cellulose-based ? ers.
Keratin contains ‘‘bound water’’  that is strongly hydrogen-bonded in the protein structure and seems to persist to high temperatures as evidenced by DSC studies on the ? bers . The weaker ? ber/polymer interactions, evidenced by the voids representative of ? ber pullout, may be a result of ? bers that contain more bound water during processing. The water-containing ? bers would not be very compatible with the hydrophobic LDPE matrix. The stronger ? ber/polymer interactions may be a result of ? ber drying during processing and strong interaction of the LDPE matrix with the hydrophobic portions of the keratin.Therefore, low k values are the result of less than optimum ? ber orientation and ? ber/polymer interactions because (1) some of the ? bers are not oriented in the deformation or 1-direction, (2) some of the protein ? ber is hydrophilic and therefore incompatible with the hydrophobic polymer, (3) some bound water may make portions of the ? ber incompatible with the polymer. Fig.
3 shows a plot of the e? ect of ? ber aspect ratio, (L/D)f, on composite mechanical properties. Fiber aspect ratio is important for maximum load transfer from the polymer matrix to the ? ber. If there is strong ? er/polymer adhesion, the application of a tensile load in the ? ber direction will cause a shear stress to develop in the polymer near the interface. The shear stress will cause the polymer to plastically ? ow around the ? ber. The maximum interfacial shear stress, smax, would be related to the shear yield stress of the matrix, i. e,, smax $ sy,p.
It may not be exactly the same value if the physical properties of the polymer in a con? ned space di? er from the bulk physical properties. If there is weak adhesion, then the ? ber will slide along the interface and ‘‘pull-out’’. In other 180 J. R. Barone, W. F.
Schmidt / Composites Science and Technology 65 (2005) 173–181 words, if there is weak adhesion, shear deformation will debond the polymer from the ? ber and the shear stress originates from polymer/? ber sliding friction. In either case, the shear stress at the interface, s, is important to describe the load transfer. A force balance across the ? ber/polymer interface yields  Lc ry;f ; ? 3? ? D f 2s where ry,f is the yield stress of the ? ber. From Fig. 3, composite yield stress is maximized at ? ber aspect ratios greater than 50, which corresponds to a critical ? ber length, Lc, of 0. 025 cm.
Assume that the critical ? er length is the length necessary to load the ? ber to its yield stress. Using (Lc/D)f = 50 and ry,f = 200 MPa yields s = 2 MPa. The yield stress of the polymer is ry,p = 10 MPa. Assuming that the composite is incompressible and the ? ber and matrix are strained equally, smax $ ry,p/3 $ 3. 3 MPa. Therefore, the ? ber may be su? ciently long to load the ? ber to its yield stress by load transfer through the polymer matrix plastically ? owing around it.
The micrographs support this conclusion, i. e. , some of the ? bers have been drawn out with the polymer and some have broken in the fracture plane. The 2 MPa value (