―NOTICE: this is the author’s version of a work that was accepted for publication in Progress in Polymer Science. Changes resulting from the publishing process, such as peer review, editing, corrections, structural formatting, and other quality control mechanisms may not be reflected in this document. Changes may have been made to this work since it was submitted for publication. A definitive version was subsequently published in Progress in Polymer Science, Volume 37, Issue 6, June 2012, Pages 767–780, DOI: 10.1016/j.progpolymsci.2011.09.005‖
Development of self-reinforced polymer composites
Chengcheng Gaoa, Long Yua,b*, Hongsheng Liua, Ling Chena a
Centre for Polymers from Renewable Resources, ERCPSP. South China University of
Technology, Guangzhou 510640, China b
CSIRO Materials Science and Engineering, Private Bag 33, Clayton South, Melbourne,
Victoria 3169, Australia
Abstract This paper reviews the development of single-polymer or self-reinforced composites (SRCs), including the fundamental sciences such as design principles and mechanisms, as well as their preparation techniques and potential application areas. The advantages of such SRC systems include the ability to achieve excellent interfaces between components, their pure chemical functionality, and their higher value as recyclable products due to their relative homogeneity compared to composites composed of different classes of components. Single-polymer composites are particularly important in biomaterials applications, since any additives composed of different chemicals could affect biocompatibility and biodegradation. Various techniques used to design and produce SRCs have been investigated and developed, such as hot compaction, overheating, solution, partial dissolving, cool drawing, physical treatment and chemical modification.
*
Corresponding author: Tel.: +61 3 9545 2777; fax: +61 3 9545 2797. E-mail address:
[email protected] (L. Yu)
Keywords Composites, self-reinforced, interface.
Nomenclature AE
acoustic emission
DMTA
dynamic mechanical thermal analysis
DPIM
dynamic packing injection molding
DSC
differential scanning calorimetry
HDPE
high-density polyethylene
iPP
isotactic polyethylene
LDPE
low-density polyethylene
MC
moisture content
MD
flow direction
OPIM
oscillating packing injection moulding
PA
polyamide
PCL
polycaprolactan
PDLA
poly D-lactide
PE
polyethylene
PET
polyethylene teraphthalate
PGA
polyglycolid
PI
polyimide
PLA
poly(lactic acid)
PLLA
poly L-lactide
PMMA
poly(methyl methacrylate)
2
PP
polypropylene
SAXS
small-angle X-ray scattering
SEM
scanning electron microscopy
SRC
self-reinforced composite
SRP
self-reinforced polymer
TEM
transmission electron microscopy
TD
transverse direction
TPS
thermoplastic starch
UHMWPE
ultra-high-molecular-weight polyethylene
WAXD
wide-angle X-ray diffraction
XRD
X-ray diffraction
3
Contents
1. Introduction 2. Polyolefin-based SRCs 2.1. PE-based composites 2.2. PP-based composites 3. Polyester-based SRCs 3.1.
PLA-based composites
3.2. PET-based composites 3.3.
PMMA-based composites
4. SRCs from natural polymers 4.1.
Protein-based composites
4.2.
Cellulose-based composites
4.3.
Starch-based composites
5. Other kinds of SRCs 6. Summary and future works References
4
1. Introduction
Self-reinforced composites (SRCs), or single-polymer composites, in which a polymer matrix is reinforced with oriented fibers and tapes, or particles of the same polymer, are an interesting concept. The advantages of SRC systems include the ability to achieve perfect interfaces between composite components, their pure chemical functionality, and their higher value as recyclable products due to their relative homogeneity compared to composites composed of different classes of components [1–4]. This homogeneity is particularly important in biomaterials applications, since any additives composed of different chemicals could affect biocompatibility and biodegradability. The concept of single-polymer composites was first described in 1975 by Capiati and Porter [5], who used oriented polyethylene (PE) filaments and PE powder with different melting points. Since then, a variety of techniques for combining oriented fibers with a separate homogeneous phase as the matrix have been reported in the literature. The most reported studies are based on the use of polyolefin and polyester materials, and the most frequently used technique is film stacking, in which the chosen matrix film generally has a lower melting point than the fibers, so that only the interleaved film melts. To date, various techniques used to design and produce SRCs have been investigated and developed, such as overheating, solution, partial dissolving, cool drawing, physical treatment and chemical modification. The concept of ―overheating‖ is an established method for manufacturing SRCs, and it has been validated for two categories of semi-crystalline polymers—the drawable apolar polymers, and the less drawable polar polymers. The interchain interactions in apolar polymers are relatively weak and therefore a high degree of drawability can be obtained. Polar polymers, on the other hand, have relatively strong interchain interactions and are therefore
5
less drawable. A shift in melting temperature of >20C can be achieved in the case of highly extended isotactic polypropylene (iPP) (draw ratios >14), while ultra-drawn PE can only achieve 10C overheating on constraining, and this is due mainly to a change in PE chain mobility during the hexagonal phase. In the case of polyethylene teraphthalate (PET) and polycaprolactam (PA6), draw ratios of only 4 are attainable; however, temperature shifts of about 10C for constrained fibers compared to unconstrained fibers have been measured. More recently, some SRCs from natural or renewable resources have been developed. It is interesting to note that the best examples of SRCs could be natural materials. For example, a block of wood is a kind of cellulose-based SRC, and animal muscle is a kind of SRC of protein. It is noted that Bárány and colleagues [6] recently reviewed self-reinforced polymeric materials, which are classified according to their constituents (single- or multi-component), their production (one-step or in multi-step procedures) and the spatial alignment of the reinforcing phase in the matrix (in one, two or three dimensions). Matabola [7] also reviewed SRCs mainly reinforced by fibers. This new review focuses on the development of polymer– polymer composites, including the fundamental sciences such as design principles and mechanisms, as well as preparation techniques and application areas, including new areas such as biomaterials and composites from renewable resources. Various techniques used to design and produce SRCs have been investigated and developed, such as hot impaction, overheating, solution, partial dissolving, cool drawing, physical treatment and chemical modification, will be discussed.
6
2. Polyolefin-based SRCs
2.1.
PE-based composites Composites based on high-density PE (HDPE) filaments and powders with different
melting points were first reported in 1975 by Capiati and Porter [5]. Their work was based on the fact that aligned and extended chains provide thermodynamically more stable crystals, and thus higher melting points than conventionally crystallized melts. They observed the growth of transcrystalline regions at the interface of the melt matrix, plus a partial melting between fibers and matrix—indications of a strong and intimate interfacial bond with a gradient in morphologies for the HDPE system studied. The interfacial strength in the PE composites was due mainly to unique epitaxial bonding, rather than radial forces induced by compressive shrinkage. Hine et al. [2] have described in detail the preparation of PE- and PP-based SRCs by combining the processes of hot compaction and film stacking. Assemblies of oriented fibers and tapes were produced with and without an interleaved film of the sample polymer type, over a range of processing temperatures. Figure 1 shows a woven cloth made from multifilament bundles, and a permanganic etched image of compacted woven melt-spun multifilaments from a PE-based SRC. Temperatures were carefully chosen so that in some cases only the interleaved film was melted, and in other cases both the film and the fiber surfaces were melted, allowing a traditional film stacking process, a traditional hot compaction technique, and a combination of the two to be compared. In most cases, the optimum compaction temperature was about 1C below the point at which substantial crystalline melting occurred [8]. At this optimum temperature, differential scanning calorimetry (DSC) melting studies showed that about 30% of the original oriented phase had been lost to bonding the structure together.
7
Deng and Shalaby [9] found that tensile strength, tensile modulus and creep resistance were significantly increased when ultra-high-molecular-weight polyethylene (UHMWPE) fibers were incorporated into a UHMWPE matrix. The longitudinal tensile strength of the resultant SRCs increased with fiber content, according to the law of mixtures, however transverse strength remained unchanged for fiber contents of 40% increase in tensile strength and a 30–40% increase in impact strength in TD. The OPIM moldings showed different stress–strain behaviors in MD and TD. Chen and Shen [21] used SEM and WAXD to also studied the morphology and crystallinity of iPP SRCs produced by conventional injection molding and OPIM, and found that the OPIM moldings exhibited three crystalline forms: , and phases. The WAXD results indicated that the highest -phase orientation and largest proportion of -phase content was in the outer shear region of the OPIM moldings. Loos et al. [22] used optical microscopy and low-voltage SEM techniques to examine the morphology of SRCs prepared by embedding constrained high-modulus iPP fibers into thin
12
matrix films based on the same iPP grade. After isothermal crystallization from the melt, a transcrystalline layer of lamellar crystals formed perpendicular to the fiber axis. This work illustrated that PP SRCs are feasible, and that these composites may fulfil the demands for fully recyclable engineering composites. Li et al. [23] prepared SRCs by introducing iPP fibers into supercooled iPP matrices, and used optical microscopy to study the resultant matrix supermolecular structures as a function of fiber introduction temperature. The results showed that the supermolecular structures of the iPP matrices can be attributed to iPP transcrystallization, which is triggered by strong heterogeneous nucleation rather than the shear stress produced by fiber introduction. The morphological features demonstrated that the interfacial transcrystalline iPP was composed of purely -form if the fibers were introduced at 138C, above which an increase in the -iPP content surrounding the iPP fibers was observed. A transcrystallization layer of mainly modified iPP was observed at a fiber introduction temperature of 173C, which was associated with the melting, or at least partial melting, of the iPP fibers. The formation of rich -iPP transcrystallization layers at even higher fiber introduction temperatures (e.g. 178C), further demonstrated that the melting state of the iPP fibers plays an important role in generating -iPP. Li et al. [24] used optical microscopy to study the supermolecular structures of iPP SRCs as a function of crystallization temperature, and found that although the partial melting of iPP fibers favored the initiation of -iPP crystal growth, the interfacial morphology of the composites strongly depended on the crystallization temperature. They found that transcrystalline structures of negative radial (III)-iPP or banded (IV)-iPP were produced within the crystallization temperature range of 105–137C, while transcrystallization zones of pure negative radial (II)-iPP crystals were observed at higher crystallization temperatures (e.g. 141C). On the other hand, the surrounding iPP spherulites formed from the bulk were
13
composed of -iPP for the whole crystallization temperature range, however the optical character of the spherulites was controlled by the thermal condition. Li et al. [25] also studied the morphologies of iPP SRCs at the lamellar level using field emission SEM. The results showed that the induced supermolecular structure of the iPP strongly depended on the fiber introduction temperature. At temperatures lower than 164C, no iPP fiber melting occurred, and the solid iPP fibers could only initiate growth of parallel aligned -iPP lamellae normal to the fiber axis. At temperatures above 165C, the morphological change from microfibrils to lamellae implied that melting and recrystallization of the iPP fiber had occurred during sample preparation. At the same time, the increment of iPP content with increasing fiber introduction temperature demonstrated unambiguously that -iPP crystallization was associated with the above-mentioned melting and recrystallization process. Combining the structural features of the recrystallized iPP fibers and the induced interfacial layers, the authors suggested that the chain orientation status of the molten iPP fiber plays an important role in subsequent -iPP crystallization. Barkoula et al. [26] consider that the concept of ―overheating‖ is validated on two categories of semi-crystalline polymers: the drawable apolar polymers (e.g. iPP, UHMWPE) and the less drawable polar polymers (e.g. PET and polyamides (PAs)). The interchain interactions in apolar polymers are relatively weak and therefore a high degree of drawability can be obtained. Polar polymers, on the other hand, have relative strong interchain interactions and are therefore less drawable. The authors found that a shift in the melting temperature of more than 20C could be achieved in the case of highly extended iPP (draw ratios >14), while ulra-drawn PE only achieved 10C overheating upon constraining, due mainly to a change in PE chain mobility in the hexagonal phase. In the case of PET and PA6, draw ratios of only 4 were attainable; however, temperature shifts of about 10C were observed for constrained fibers. A proof of principle of the potential of the constraining
14
concept for the manufacture of SRCs was undertaken by preparing single-fiber iPP model composites, and examining the effects of post-drawing conditions on overheating. It was concluded that both post-drawing temperature and ultimate draw ratio have a significant influence on the degree of overheating. Cermák et al. [27] studied the ability of elongational flow to induce chain orientation and, as a consequence, improve the properties of extruded profiles. For the study, a novel extrusion die with a semihyperbolic convergence of the channel that generated a high percentage of elongational flow in the melt, was connected to a conventional single-screw extruder. Selfreinforced rod-like extrudates prepared using this processing line possessed outstanding storage modulus at a wide range of temperatures, and they exhibited considerably higher melting points and fusion temperature than rods produced by conventional extrusion. Khondker et al. [28] studied optimum processing conditions for producing high-quality all-PP SRCs by injection compression molding. They found that the tensile and three-point bending properties of the virgin PP were almost unaffected by the introduction of knitted fabric reinforcing layer(s), due to the fabric’s very low fiber content. The three-point bending properties were also unaffected by the surface of indentation–flexure. The applied impact energy was maintained at 5 J for the homo-PP and 27 J for the block-PP materials, respectively, to ensure penetration during drop-impact tests. It is interesting to note that, when compared with the virgin matrix material, the homo-PP SRCs exhibited superior energy absorption capability, and their corresponding plate bending performance also showed consistent improvement. On the other hand, although the virgin block-PP matrix material exhibited better impact performance than its composites with homo-PP knitted fabric, a notably small increase in the fiber content of the fabric produced a considerable improvement in impact properties. These homo-PP/block-PP SRCs have clearly demonstrated their potential to outperform block-PP materials, via the modification and/or manipulation of the
15
reinforcement’s structural/geometrical parameters and fiber content. Both static and dynamic impact properties were likely to have been affected by the local area properties of the face under indentation testing, thereby contributing to the improved performance of the composite specimens where the knit face was under impact. Cabrera et al. [29] examined the thermoforming behavior of an SRC based on co-extruded PP tapes. In contrast to traditional continuously woven glass fabric-reinforced polypropylene (GF/PP) materials for which the sole mode of deformation is either inter- or intraply shearing, all-PP composites have an additional mode of deformation because the fibers (or in this case tapes) can still be deformed. The importance of this additional deformation mode was investigated by Cabrera et al. in a range of stamping experiments in combination with 3D strain mapping experiments. Non-isothermal thermoforming experiments revealed that all-PP woven fabric laminates based on tapes deformed in a different manner to traditional GF/PP. Although the main mode of deformation of both all-PP and GF/PP for the investigated dome parts was intraply shearing, a much lower energy was required to deform the all-PP laminates. The authors recommended that, whenever possible, tape deformation by drawing should be avoided as it requires higher energy, which may lead to higher residual stresses in the final product. However, tape drawing may prove an essential benefit when complex shapes are involved. Kim et al. [30] systematically investigated the optimum consolidation process conditions for a co-extruded PP SRC. They observed that even very small changes in the processing conditions could significantly influence the mechanical behavior of consolidated SRCs (e.g. peel strength could be improved by approximately 30% within a 10C change in temperature). Their investigation included structural analyses such as DSC and X-ray diffraction (XRD), and mechanical tests such as T-peel, tensile and creep tests. A creep potential with three orthotropic material parameters was utilized to describe the anisotropic and nonlinear time-
16
dependent deformation behavior of the SRCs. Combined with the creep potential, the creep test of an off-axis coupon specimen with a fiber orientation of 22.5 was found to represent the anisotropic and nonlinear creep behavior of other off-axis specimens within a reasonable margin of error. In studying the impact behavior of PP SRCs, Bárány et al. [31] used and polymorphs of homo-iPP and random copolymer (with ethylene) for matrix materials, and a fabric woven from highly stretched split PP yarns as reinforcement. Composite sheets were produced by film stacking and consolidated by hot pressing at 5 and 15C above the melting temperature of the matrix material. The composite sheets were subjected to static tensile, dynamic dropimpact and impact tensile tests at room temperature. Dynamic mechanical thermal analysis (DMTA) was also performed on the related composites and their constituents. The results indicated that the -modification of the homo-PP was more straightforward than that of the PP copolymer. Increasing the temperature of consolidation usually increased stiffness and strength, however toughness (tensile impact strength, perforation impact energy) decreased. Based on fractographic results, this was assigned to differences in the failure mode. Izer et al. [32] have also reported on the development of PP SRC using and crystal forms of homo-iPP and random copolymer (with ethylene) as matrix materials, and highly stretched split PP yarns as reinforcement. Composite sheets were produced by film stacking and compression molding at different processing temperatures, while keeping the holding time and pressure constant. The quality of the composite sheets was assessed by optical microscopy, and density and peel strength measurements. PP SRC specimens were subjected to static tensile and flexural strength, and dynamic drop-impact tests and the results were analyzed as a function of processing temperature and polymorphic composition. Based on the results, the optimum processing temperature was determined to be 20–25C above the related matrix melting temperature. It was established that the -modified homo-PP-based one-
17
component PP SRCs possessed similar mechanical properties to the extensively studied random PP copolymer-based two-component ones. Alcock et al. [33] reported on direct forming routes to manufacture simple geometries of all-PP SRCs, by molding woven co-extruded PP tapes directly into composite products, without the need for a pre-consolidated sheet. High-strength co-extruded PP tapes have potential processing advantages over mono-extruded fibers or tapes, as they allow a larger temperature processing window for consolidation, which makes direct forming routes feasible, without the need for an intermediate pre-consolidated sheet. Thermoforming studies have shown that direct forming is an interesting alternative to stamping pre-consolidated sheets, as it eliminates an expensive belt-pressing step normally required in the manufacturing of semifinished sheet products. Moreover, results from forming studies have shown that directly forming a simple dome geometry from a stack of fabrics required only half the energy needed to stamp the same shape from a pre-consolidated sheet. Alcock et al. [34] also studied the effects of tape drawing parameters such as draw ratio (), drawing temperature and thermal annealing on a tape’s final mechanical properties, density and dimensional thermal stability. PP tapes drawn to = 17 had tensile moduli of 15 GPa and tensile strengths of 450 MPa. PP tapes at lower draw ratios ( > 9.3) showed a decrease in density, a change in appearance from transparent to opaque, and increased dimensional thermal stability with increasing draw ratio. The results of an investigation into the effects of a thermal annealing step, targeted at improving the dimensional thermal stability of these highly oriented PP tapes, were also reported.
3. Polyester-based SRCs
3.1. PLA-based composites
18
Derived from agricultural products, poly(lactic acid) (PLA) has become increasingly popular as an engineering plastic because it is biodegradable and compostable, it has high mechanical strength, and it is easier to process than other biopolymers [35]. However, because of its unsatisfactory impact resistance and low heat distortion temperature in the raw state, PLA has not yet gained full market acceptance as an engineering resin. Nevertheless, reinforcing a PLA matrix by embedding PLA fibers is one possible means of meeting the demands for high strength and stiffness required for many applications in the aerospace, automotive, construction, medical and chemical industries. The development of high-stiffness and high-strength polymeric fibers is essential in imparting superior mechanical properties onto the resulting PLA SRCs [36]. The mechanical properties of fibers can be increased via molecular orientation during spinning and drawing [37], and the most commonly used methods of producing PLA fibers are melt spinning and electrospinning. In the melt spinning method, as used by Li and Yao [38] and Mäkelä et al. [39] produce reinforcing PLA fibers, a polymer melt is extruded through small orifices in a spinneret and drawn into thin fibers by a uniaxial drawing process. In the electrospinning method, a fibrous polymer solution is charged and ejecting through a nozzle onto an oppositely charged grounded target. Tsuji et al. [40] electrospun PLA stereocomplex nanofibers and found that the melting temperature increased to about 220C, compared to 178C for the pure polylactides of poly L-lactide (PLLA) and poly D-lactide (PDLA). A PLA stereocomplex could thus be used to increase the processing temperature window [40]. Significantly improved interfacial bonding can be achieved in materials where both matrix and reinforcing elements have the same chemical structure [41]. For example, SRCs consisting of oriented PLA fibers surrounded by a PLA matrix will have improved strength and rigidity compared to non-reinforced PLA. Tormala et al. [42] developed SRCs consisting of a PLLA matrix reinforced with highly oriented PLLA fibers, and attained much higher
19
strengths than an earlier unreincorced implant for medical applications. Figure 2 shows a schematic representation of a PLA-based SRC manufactured by partially sintering the fibers together at high temperature and pressure. A similar technique has also been used to produce SRCs based on polyglycolid (PGA) [43]. Majola et al. [44] showed that the initial bending and shear strengths of PLLA SRC implants were about 250–271 MPa and 94–98 MPa respectively, compared to 145 MPa and 53 MPa respectively for non-reinforced PLLA implants with the same molecular mass. Similarly, Wright-Charlesworth et al. [45] reported that a PLA SRC fabricated from oriented PLA fibers and formed by the hot compaction method had the highest initial mechanical properties. Li and Yao [38] used film stacking and compression molding to prepare PLA SRCs consisting of amorphous sheets as the matrix and highly crystalline fibers, yarns or fabrics as reinforcement. In the study, two amorphous PLA sheets and a layer of PLA fibers, fabric or yarns were compressed and laminated between two heated platens on a Carver hydraulic press. Successful lamination was achieved at a compression pressure of around 1.5 MPa, at which not only could composite thickness be well controlled, but the resultant SRCs were free from air bubbles. With the slow crystallizing characteristics of PLA, a processing temperature window greater than 30C was found to maintain the integrity of the fibers after consolidation. Moreover, due to strong interfacial bonding, the PLA SRCs exhibited a single point of failure during tensile testing, and the original texture-induced structural elasticity of the reinforcement was restrained. The PLA SRCs exhibited enhanced mechanical properties, in particular the tearing strength of the fabric-reinforced samples was almost an order higher than that of the unreinforced PLA. Significant improvements in tensile strength (58.6 MPa) and Young’s modulus (3.7 GPa) was also reported [35]. While PLA SRCs possess improved mechanical properties, problems with biocompatibility have not been addressed, indicating
20
that these composites have potential applications in the packaging industry. To date, sutures, rods, screws and plates made of self-reinforced PLA have been widely available for the fixation of cancellous bone fractures and osteotomies, since the mechanical properties of these materials are sufficient for their use with load-bearing bones. Recently, Yu and his group [46,47] evaluated various physical treatments, such as cool drawing, annealing, CO2 treatment and nucleation, and their effects on incresasing the difference in melting temperature between the PLA matrix and PLA fibers in PLA-based SRCs. A thermal processing window of approximately 7–9C was achieved after cool drawing and annealing the PLA fibers. Thermal treatment (annealing) has also been used in developing non-polymer SRCs. For example, Kim et al. [48] reported that after annealing, the microstructure of alpha-silicon carbide changed from equiaxed grains to a self-reinforced microstructure consisting of large elongated grains and small equiaxed grains, which resulted in significant improvements in toughness.
3.2. PET-based composites Hine and Ward [49] described the development of a process for producing hot-compacted sheets of woven PET multifilaments, and reported on their mechanical properties. Investigation of the various processing parameters showed that a key aspect was the time spent at compaction temperature, or dwell time. Molecular weight measurements, using intrinsic viscosity, showed that hydrolytic degradation occurred rapidly at the temperatures required for successful compaction, leading to embrittlement of the resulting materials with increasing dwell time. A dwell time of two minutes was found to be optimum, as this produced the required percentage of melted material to bind the structure together, while incurring only a small decrease in molecular weight. The authors used a variety of techniques, including mechanical tests, DSC and SEM, to examine the mechanical properties and
21
morphology of the optimum compacted sheets. The results of these tests reinforced the conclusions of previous studies of hot-compacted PP SRCs, whose behavior is a combination of the properties of the two components, i.e. the original oriented multifilaments and the melted and recrystallized matrix. Other key findings from the research included a confirmation of the importance of high ductility in the melted and recrystallized phases (promoted by high molecular weight or by suppressing crystallinity during processing), and the proportionately high impact performance of hot-compacted sheets compared to that of other materials. Rojanapitayakorn et al. [50] fabricated PET SRCs by compacting oriented PET fibers under pressure at temperatures near to but below their melting point. Figure 3 shows the processing scheme for the hot compaction processing. The originally white fiber bundles, which were about 40% crystalline, showed increased crystallinity (55%) but optical translucency after processing. Crystalline orientation, gauged using the Hermans orientation parameter from WAXD data, indicated that no significant loss in orientation of the crystalline fraction occurred during compaction. Mechanical characterization revealed that increasing the compaction temperature from 255 to 259C produced a stepwise decrease in flexural modulus (9.4 to 8.1 GPa) and a concomitant increase in transverse modulus and strength. This transition in behavior was accompanied by a loss of optical transparency and a change in the distribution of the amorphous fraction from fine to coarse interfibrillar domains, as observed by SEM. The mechanical properties of PET compactions were influenced more by orientation of the amorphous phase than by that of the crystalline phase. Characterization using an unnotched Charpy test method showed remarkable impact resistance after compaction, however impact toughness decreased with increasing compaction temperature. Zhang et al. [51] reported on all-PET SRCs prepared by film stacking oriented PET tapes and co-PET films. A processing temperature window was determined through a series of tests
22
(including DSC and T-peel tests) on the tapes and films. The tensile properties of the PET tapes, co-PET films and all-PET composites were reported and compared with a commercial co-extruded pure-PP tape. The effects of compaction temperature and pressure on the tensile properties of the all-PET SRCs were investigated to explore the optimum processing parameters required to balance good interfacial adhesion and the residual tensile properties of the PET tapes. Duhovic et al. [52] recently reported on nanofibrillar PET SRCs prepared by first selectively extracting PP from a knitted nanofibrillar PP/PET (80:20 by wt) textile, and then sandwiching the remaining PET textile (as reinforcement) between lower melting point PET films before compression molding at 120C. The resultant SRCs showed improvements in tensile strength and modulus of 37–100% and 40–140% respectively (depending on the annealing temperature after compression molding and the test direction), compared to those of the starting isotropic matrix film.
3.3. PMMA-based composites Gilbert et al. [53] reported on poly(methyl methacrylate) (PMMA) SRCs consisting of high-strength, high-ductility PMMA fibers embedded in a PMMA matrix. Two types of PMMA fibers with different mechanical properties and of different sizes (40 and 120 m diameter) were used. The results showed that the tensile strength, tensile modulus and tensile strain-to-failure of the SRCs were significantly greater than commercial PMMA (P < 0.05). The flexural strength of the SRCs was not increased compared with PMMA alone, but it was greater than that of bone cement (P < 0.05). There was no difference in flexural modulus between any groups. The flexural strain-to-failure (30–35% for SRCs) was about three times greater compared to PMMA and bone cement, and fracture toughness was also significantly greater (P < 0.001). The SRCs incorporating the 40 m fibers achieved fracture toughness
23
values of 3.2 MPa m1/2, compared to 2.3 MPa m1/2 for the 120 m samples and 1.3 MPa m1/2 for PMMA and bone cement. The fatigue strength of both fiber diameter samples was significantly greater (P < 0.001) at 80 MPa (106 cycles), compared to about 18 MPa for PMMA and bone cement. Fatigue damage in the form of fiber splitting and fiber–matrix interfacial failure was observed in the SRC samples, while the PMMA and bone cement exhibited only smooth fractures. The hysteresis damage energy to failure was about 25 times greater in the SRC samples (2000 J at 106 cycles, compared to 80 J at 106 cycles for PMMA or bone cement), indicating much greater tolerance to fatigue damage in these materials. Wright(-Charlesworth) and colleagues [54–58] have extensively studied PMMA SRCs and their applications. Three different weaves of PMMA SRCs have been evaluated for bending and fracture toughness in air, after immersion in saline solution for 30 days at 37C, and after gamma irradiation followed by the same immersion regime [54]. Bending modulus and strength were decreased by gamma irradiation followed by immersion, however the effects on these properties of immersion alone was negligible. Both methods were shown to have little or no effect on the fracture toughness of the SRCs, however the different weaves produced different fracture processes, which were related to the different orientation of fibers to the fracture toughness pre-crack. Optimally consolidated PMMA SRCs absorb the same amount of water as bone cement. When compared with the results of previous work using bone-cement controls, the results of these studies [54–58] have shown that the performance of PMMA SRC is equal to or better than bone cement in all tests performed, and therefore it warrants further consideration as a candidate biomaterial. After studying the strength of the bonds that PMMA SRC formed with simulated prostheses and bone cement, Wright et al. [55] concluded that one potential use for the composite was as a precoat for hip or other stemmed prostheses, as it has a similar chemical composition to bone cement. PMMA SRC was woven around cobalt–chromium (Co-Cr) rods,
24
and push-out tests were performed on samples that were tested in air as processed or after immersion in saline solution for 30 days at 37C. Three different weaves were investigated and compared to bone cement. The bone cement and PMMA SRC formed interfacial bonds with the Co-Cr rods that failed at an average load (stress) of 980 N (2.0 MPa). After immersion, the interfacial bond strength of the bone cement was 642 N (1.23 MPa), and that of the tight-weave PMMA SRC was statistically stronger at 973 N (1.86 MPa). The shear strengths of the bone cement alone, as measured by push-out tests, were an order of magnitude higher at 9210 N (15.2 MPa) in air and 9900 N (15.7 MPa) after immersion. However, the shear strengths of the bonds between the PMMA SRC and the bone cement were 10900 N (17.9 MPa) in air and 9610 N (15.8 MPa) after immersion. Wright et al. [56] systematically examined the effects of processing time and temperature on the thermal properties, fracture toughness and fracture morphology of PMMA SRCs produced by hot compaction. DSC results showed that the composites containing high amounts of retained molecular orientation exhibited both endothermic and exothermic peaks, depending on processing times and temperatures. An exothermic release of energy just above the glass transition temperature is related to the release of retained molecular orientation, and in this study the release of energy decreased linearly with increasing processing temperature or time over the range investigated. A maximum fracture toughness of 3.18 MPa m1/2 was recorded for samples processed for 65 minutes at 128C. The authors found that optimal structure and fracture toughness were obtained in composites with maximum interfiber bonding and minimal loss of molecular orientation. Composite fracture mechanisms were highly dependent on processing conditions: low processing times and temperatures resulting in more interfiber/matrix fracture; higher processing times and temperatures resulting in higher ductility and more transfiber fracture; and excessive processing times resulted in brittle failure.
25
Wright-Charlesworth et al. [57] also reported on the potential of uniaxial PMMA SRC as a precoat material for the femoral component of a total hip replacement. The production of SRCs via hot compaction is largely an empirical process in which the processing parameters of time, temperature and pressure are varied until the desired properties are obtained. The composites produced for this work were processed at times and temperatures that spanned the relaxation process of a single fiber. It was found that molecular orientation, as measured by birefringence, was lost in composites processed at times greater than the relaxation times of the single fibers. Flexural strength was also found to vary with processing conditions, with the highest values of 165 15 MPa and 168 3 MPa found at high and low processing times respectively—significantly stronger than unreinforced PMMA at 127 14 MPa. The authors hypothesized that diffusion between fibers occurs much more quickly than loss of molecular orientation, and demonstrated that PMMA SRC processing conditions can be predicted from the relaxation times and temperatures of single fibers. Wright-Charlesworth et al. [58] further investigated the properties of PMMA SRCs as a function of processing temperature, by performing nano-indentation tests to measure hardness and modulus at the nanoscale. Nanoscratch tests were performed parallel, orthogonal and longitudinal to fiber orientation to measure residual scratch depths. Significant differences were observed in the hardness, modulus and residual scratch depth as a function of processing temperature, when compared to unreinforced PMMA. As processing temperature increased, hardness decreased and residual scratch depth increased. Data also showed that fiber orientation played a critical role in scratch resistance, with orthogonal scratching producing the least residual scratch depth (ranging from 524 nm at 105C to 838 nm at 150C, compared to 842 nm for unreinforced PMMA).
26
4. Self-reinforced composites from natural polymers
More recently, some SRCs from renewable resources have been developed, with a particular emphasis on the importance of these materials for use as biomaterials.
4.1. Protein-based composites Li and Lu [59] directly manufactured SRC fabrics using regenerated silk fibroin, and test demonstrated that the interfacial adhesion between the silk fibers and the fibroin matrix was enhanced by controlling the fiber dissolution through a 6 mol L-1 LiBr aqueous solution. The overall mechanical properties as well as the thermal stability of the silk fiber/fibroin composites were significantly improved, comparing to those of pure fibroin counterparts.
4.2. Cellulose-based composites Cellulose, a linear homopolymer of glucose (C6H10O5)n, where n ranges from500 to 5000, has been widely used to reinforce other polymers, especially biodegradable polymers [35,60]. Cellulose cannot be simply thermally processed due to its unique crystalline structure, and one of the techniques of producing cellulose-based SRCs is to partly dissolve the cellulose, in particular the surface. Lu et al. [61] reported on the development of self-reinforced meltprocessable composites of sisal. Through a mild benzylation treatment, skin layers of sisal fibers were converted into thermoplastic material, while the core of the fiber cells remained unchanged. On the basis of these modified sisal fibers, self-reinforced composites were prepared using hot pressing, in which the plasticized parts of sisal served as thematrix and the unplasticized cores of the fibers acted as reinforcement. The authors found that there was a balance between melt processability and the reinforcing effects of the benzylated sisal fibers. Unlike conventional plant-fiber composites using petro-polymers as matrices, the self-
27
reinforced sisal composites were characterized by inherent interfacial compatibility and full biodegradability. Similarly, Nishino et al. [62] reported on an all-cellulose composite prepared by distinguishing the solubility of the matrix cellulose into the solvent from that of the cellulose fibers through pretreatment. The tensile strength of the uniaxially reinforced all-cellulose composite was 480 MPa at 25°C, and the dynamic storage modulus was as high as 20 GPa at 300°C, which are comparable or even higher values than those of conventional glass-fiberreinforced composites. In addition, a linear thermal expansion coefficient of about 10–7 K–1 was achieved. This all-cellulose composite offers substantial advantages, in that it is composed of a sustainable resource, there is less interface between the fibers and the matrix, it exhibits excellent mechanical and thermal performance during use, and it is biodegradable after service. Gindla and Keckes [63] studied cellulose-based nanocomposite films with different ratios of cellulose I and II, produced by partial dissolution of microcrystalline cellulose powder in lithium chloride/N,N-dimethylacetamide and subsequent film casting. The films were isotropic, transparent to visible light, highly crystalline, and contained different amounts of undissolved cellulose I crystallites in a matrix of regenerated cellulose. The results showed that, by varying the cellulose I and II ratio, the mechanical performance of the nanocomposites could be tuned. Depending on the composition, a tensile strength up to 240 MPa, an elastic modulus of 13.1 GPa, and a failure strain of 8.6% were observed. Moreover, the nanocomposites clearly surpassed the mechanical properties of most comparable cellulosic materials, with their greatest advantages being that they are fully biobased and biodegradable, as well as possessing relatively high strength. Gindl and Keckes [64] also reported on self-reinforced cellulose films prepared by incomplete dissolution of commercial microcrystalline cellulose in LiCl/DMAc solvent, and
28
the subsequent coagulation of regenerated cellulose in the presence of undissolved microcrystalline cellulose. By drawing in wet conditions and subsequent drying, the preferred orientation was introduced into the self-reinforced cellulose films, resulting in significantly improved tensile strength of up to 430 MPa and modulus of elasticity of up to 33 GPa. A linear relationship was observed between the applied draw and the orientation of cellulose in the films, and the measured elastic modulus and tensile strength, respectively. The strength and modulus of elasticity of the optically transparent drawn films significantly surpassed those of existing all-bio-based planar materials. Therefore, self-reinforced cellulose films present a biodegradable alternative to non-bio-based materials with similar performance.
4.3. Starch-based composites Starch is a heterogeneous material with a more sophisticated microstructure than that of conventional polymers. Most native starches are a mixture of amylose (a linear structure of α1,4-linked glucose units) and amylopectin (a highly branched structure of short α-1,4 chains linked by α-1,6 bonds). Thermal processing of starch-based materials is much more complex than for conventional polymers, due to the multi-phase transitions involved [65–71]. Starch is one of the most promising natural polymers because of its inherent biodegradability, overwhelming abundance and its ongoing annual renewal. The development of biodegradable starch-based materials for general applications has been spurred by oil shortages and the growing interest in easing the environmental burden of petrochemically derived polymers [72–76]. Starch-based materials or associated composites have also been developed for more high-value applications such as biomaterials, for instance in tissue engineering [77–79]. However, improving the mechanical properties and processability of starch-based materials is an ongoing challenges.
29
Recently Lan et al. [80] reported on a unique technique of designing and preparing selfreinforced starch-based films using a high amylose cornstarch that was chemically modified in different ways. Hydroxypropylation was used to decrease the gelatinization temperature and to improve processability. Lowering the processing temperature by reducing the gelatinization temperature was due mainly to the presence of the substituted group, which weakened or strained the internal bond structure holding the granules together [81]. The reinforcing component consisted of crosslinked starch granules, where the crosslinking increased granular thermal stability and moisture resistance. Since both hydroxypropylated and crosslinked starches are from broadly the same material, albeit slightly changed through different chemical modifications, it is hard to distinguish them using conventional techniques. The distribution of crosslinked starch has been imaged by Mano et al. [79] using confocal laser scanning microscopy (see Figure 4). It was found that the modulus and tensile properties of the starch films were increased by about 30% and 20% respectively after the addition of rigid crosslinked starch particles (20% wt) with no significant loss in failure strain. A perfect interface between matrix and reinforcing agent was obtained as both materials are derived from the same resource, albeit with different chemical modifications.
5. Other kinds of SRCs As mentioned previously, ―SRC‖ generally refers to a composite comprising polymeric oriented reinforcing elements (usually fibers or tapes) or rigid particles in a matrix of the same polymer. However, there are other kinds of SRCs based on molecular orientation through synthesizing or processing. Fan et al. [82] reported on a novel in-situ self-reinforced polyimide (PI) film material, producred using a PI with mesogenic units, which was synthesized by the reaction between the pendent hydroxyl group of PI and the epoxy group of
30
a liquid crystalline compound in the presence of N,N-dimethylaniline at 120–130C. The macromolecular reaction between the hydroxyl group and epoxy group was investigated using a model reaction, and the results indicated that: (a) the PI with mesogenic units exhibited a smectic phase; (b) the level of mesogenic units influenced side-chain crystallinity, as well as the thermal and mechanical properties of the polymers; and (c) the enhanced PI films showed good solubility, and higher thermal stability, tensile strength and modulus. Some polymers and copolymers are referred to as self-reinforced polymers (SRPs) because of their intrinsic high strength and modulus, without the addition of a reinforcing agent. For example, Morgan et al. [83] reported on a copolymer based on benzoyl-1,4phenylene and 1,3-phenylene, which they evaluated using nanoprobe investigation techniques and compared to the properties obtained at the macroscale. Specimens were prepared by spin casting, solvent casting and compression molding. In nano-indentation studies, molded SRP samples demonstrated one and a half to two times the surface hardness and reduced modulus compared to traditional engineering thermoplastics. These improved nanomechanical properties also contribute to the low friction coefficients of SRPs by providing more resistance to wear, deformation and local penetration.
6. Summary and future works
The advantages of polymer-based self-reinforced composites (SRCs) include the ability to achieve a excellent interface between components, their pure chemical functionality, and their higher value as recyclable products due to their relative homogeneity compared to composites with different classes of components. ―SRC‖ generally refers to a composite comprising polymeric oriented reinforcing elements (usually fibers or tapes) or rigid particles in a matrix of the same polymer. However, there are other kinds of SRCs based on molecular orientation
31
through synthesizing or processing. SRCs are particularly important in biomaterials applications, since any additives composed of different chemicals could affect biocompatibility and biodegradability. This is a new and developing area, which is far from mature in relation to the fundamental issues and processing techniques. There are only a few commercial products on the market. So far, the techniques developed to design and produce SRCs include thermal processing (―overheating‖), hot impaction, cool drawing, solution, partial dissolution, physical treatment and chemical modification. The most reported studies of polymer SRCs are based on the use of polyolefin and polyester materials. The concept of ―overheating‖ is a recognized method for manufacturing SRCs, and it has been validated on two categories of semi-crystalline polymers—the drawable polar polymers, and the less drawable polar polymers. The interchain interactions in polar polymers are relatively weak and therefore a high degree of drawability is possible. Polar polymers, on the other hand, have relative strong interchain interactions and are therefore less drawable. Another technique of developing SRCs is to increase the difference between the thermal processing temperatures of the matrix and the reinforcement (oriented fibers and tapes, or rigid particles), which are essentially the same material. Research has shown that processing is very temperature sensitive, and that in most cases the optimum compaction temperature is about 1C below the point at which substantial crystalline melting occurs. At this optimum temperature, some of the original oriented phase is lost to bonding the structure together. The small difference in melting temperature between the reinforce elements and the matrix poses a big challenge during fabrication, as both constituents have basically the same chemical structure and hence melting temperatures. As there are many polyolefins with different melting temperature available on the market, various PE- and PP-based SRCs have been developed based on the same polymer with different thermal properties. In fact, the
32
concept of single-polymer composites was first developed more than 30 years ago using oriented polyethylene filaments and polyethylene powder with different melting points. Since then, a variety of techniques for combining oriented fibers with a separate matrix phase have been reported in the literature. The most frequently used technique is film stacking, in which the chosen film generally has a lower melting point than the fibers, so that only the interleaved film melts. On the other hand, polyesters have shown multiphase transitions, in particular cool crystallization during thermal treatments, suggesting that there are opportunities to increase the difference in melting temperatures for these materials. While the various techniques for producing SRCs are far from mature and new technologies are still being developed, research so far has demonstrated the great potential of these composites. Recently, some SRCs from renewable resources (e.g. starch-based SRCs) have been developed. It is expected that more and more R&D works on SRCs will be in the areas of biomaterials and environmentally friendly materials. Furthermore, the excellent, if not perfect, interface between the reinforcement and matrix in SRCs provides a practical guide for designing and predicating the performance of composites, and a baseline for computer simulation and modeling.
33
References
[1]
Amer MS, Ganapathiraju S. Effects of processing parameters on axial stiffness of selfreinforced polyethylene composites. J Appl Polym Sci 2001;81:1136–41.
[2]
Hine PJ, Olley RH, Ward IM. The use of interleaved films for optimising the production and properties of hot compacted, self reinforced polymer composites. Compos Sci Technol 2008;68:1413–21.
[3]
Manninen MJ, Päivärinta U, Pätiälä H, Rokkanen P, Taurio R, Tamminmäki M, Törmälä P. Shear-strength of cancellous bone after osteotomy fixed with absorbable self-reinforced polyglycolic acid and poly-l-lacid acid rods. J Mater Sci—Mater Med 1992;3:245–51.
[4]
Suuronen R, Wessman L, Mero M, Tormala P, Vasenius J, Partio E, Vihtonen K, Vainionpaa S. Comparison of shear strength of osteotomies fixed with absorbable selfreinforced poly-l-lactide and metallic screws. J Mater Sci—Mater Med 1992;3:288–92.
[5]
Capiati NJ, Porter RS. The concept of one polymer composites modelled with high density polyethylene. J Mater Sci 1975;10:1671–7.
[6]
Kmetty Á, Bárány T, Karger-Kocsis J. Self-reinforced polymeric materials: a review. Prog Polym Sci 2010;35;1288–310.
[7]
Matabola PK, De Vries RA, Moolman SF, Luyt SA. Single polymer composites: a review, J Mater Sci 2009;44:6213–22.
[8]
Hine PJ, Ward IM, Jordan DN, Olley R. A comparison of the hot-compaction behavior of oriented, high- modulus, polyethylene fibers and tapes. J Macromol Sci–Phys 2001;40:959–89.
[9]
Deng M, Shalaby SW. Properties of self-reinforced ultra-high-molecular-weight polyethylene composites. Biomaterials 1997;18:645–55.
34
[10] Guan Q, Lai FS, Mccarthy SP, Chiu D, Zhu XQ, Shen KZ. Morphology and properties of self-reinforced high density polyethylene in oscillating stress field. Polymer 1997;38:5251–3. [11] Huang HX. Continuous extrusion of self-reinforced high density polyethylene. Polym Eng Sci 1998;38:1805–11. [12] Zhang G, Jiang L, Shen KZ, Guan Q. Self-reinforcement of high-density polyethylene low density polyethylene prepared by oscillating packing injection molding under low pressure. J Appl Polym Sci 1999;71:799–804. [13] Lei J, Jiang CD, Shen KZ. Biaxially self-reinforced high-density polyethylene prepared by dynamic packing injection molding. II: Microstructure investigation. J Appl Polym Sci 2004;93:1591–6. [14] Lei J, Jiang CD, Shen KZ. Biaxially self-reinforced high-density polyethylene prepared by dynamic packing injection molding. I: Processing parameters and mechanical properties. J Appl Polym Sci 2004;93:1584–90. [15] Zhuang XM, Yan X. Investigation of damage mechanisms in self-reinforced polyethylene composites by acoustic emission. Compos Sci Technol 2006;66:444–9. [16] Hine PJ, Ward IM, Jordan DN, Olley R, Bassett CD. The hot compaction behaviour of woven oriented polypropylene fibres and tapes. I. Mechanical properties. Polymer 2003;44:1117–31. [17] Jordan DN, Bassett CD, Olley R, Hine PJ, Ward IM. The hot compaction of woven oriented polypropylene fibres and tapes. II. Morphology of cloths before and after compaction. Polymer 2003;44:1133–43. [18] Huang HX. Self-reinforcement of polypropylene by flow-induced crystallization during continuous extrusion. J Appl Polym Sci 1998;67:2111–8.
35
[19] Guan Q, Zhu XG, Chiu D, Shen KZ, Lai FS, Mccarthy SP. Self-reinforcement of polypropylene by oscillating packing injection molding under low pressure. J Appl Polym Sci 1996;62:755–62. [20] Chen LM, Shen KZ. Biaxial self-reinforcement of isotactic polypropylene prepared in uniaxial oscillating stress field by injection molding. I: Processing conditions and mechanical properties. J Appl Polym Sci 2000;78:1906–10. [21] Chen LM, Shen KZ. Biaxial self-reinforcement of isotactic polypropylene prepared in uniaxial oscillating stress field by injection molding. II: Morphology. J Appl Polym Sci 2000;78:1911–7. [22] Loos J, Schimanski T, Hofman J, Peijs T, Lemstra PJ. Morphological investigations of polypropylene single-fibre reinforced polypropylene model composites. Polymer 2001;42:3827–34. [23] Li H, Jiang S, Wang J, Wang D, Yan S. Optical microscopic study on the morphologies of isotactic polypropylene induced by its homogeneity fibers. Macromolecules 2003;36:2802–7. [24] Li H, Zhang X, Duan Y, Wang D, Li L, Yan S. Influence of crystallization temperature on the morphologies of isotactic polypropylene single-polymer composite. Polymer 2004;45:8059–65. [25] Li H, Zhang X, Kuang X, Wang J, Wang D, Li L, Yan S. A scanning electron microscopy study on the morphologies of isotactic polypropylene induced by its own fibers. Macromolecules 2004;37:2847–53. [26] Barkoula NM, Peijs T, Schimanski T, Loos J. Processing of single polymer composites using the concept of constrained fibers. Polym Compos 2005;26:114–20.
36
[27] Cermák R, Obadal M, Habrová V, Stoklasa K, Verney V, Commereuc S, Fraisse F. Selfreinforcement of polymers as a consequence of elongational flow. Rheol Acta 2006;45:366–73. [28] Khondker OA, Yang X, Usui N, Hamada H. Mechanical properties of textile-inserted PP/PP knitted composites using injection-compression molding. Compos Part A—Appl S 2006;37:2285–99. [29] Cabrera NO, Reynolds CT, Alcock B, Peijs T. Non-isothermal stamp forming of continuous tape reinforced all-polypropylene composite sheet. Compos Part A—Appl S 2008;39:1455–66. [30] Kim KJ, Yu WR, Harrison P. Optimum consolidation of self-reinforced polypropylene composite and its time-dependent deformation behavior. Compos Part A—Appl S 2008;39:1597–605. [31] Bárány T, Izer A, Karger-Kocsis J. Impact resistance of all-polypropylene composites composed of alpha and beta modifications. Polym Test 2009;28:176–82. [32] Izer A, Bárány T, Varga J. Development of woven fabric reinforced all-polypropylene composites with beta nucleated homo- and copolymer matrices. Compos Sci Technol 2009;69:2185–92. [33] Alcock B, Cabrera NO, Barkoula NM, Peijs T. The effect of processing conditions on the mechanical properties and thermal stability of highly oriented PP tapes. Eur Polym J 2009;45:2878–94. [34] Alcock B, Cabrera N, Barkoula N, Peijs T. Direct forming of all-polypropylene composites products from fabrics made of co-extruded tapes. Appl Compos Mater 2009;16:117–34. [35] Yu L, Dean K, Li L. Polymer blends and composites from renewable resources. Prog Polym Sci 2006;31:576–602.
37
[36] Matabola KP, De Vries AR, Moolman FS, Luyt AS. Single polymer composites: a review. J Mater Sci 2009;44:6213–22. [37] Alcock B, Cabrera NO, Barkoula NM, Loos J, Peijs T. The mechanical properties of unidirectional all-polypropylene composites. Compos Part A—Appl S 2006;37:716–26. [38] Li R, Yao D. Preparation of single poly(lactic acid) composites. J Appl Polym Sci 2008;107:2909–16. [39] Mäkelä P, Pohjonen T, Törmälä P, Waris T, Ashammakhi N. Strength retention properties of self-reinforced poly L-lactide (SR-PLLA) sutures compared with polyglyconate (Maxon(R)) and polydioxanone (PDS) sutures: an in vitro study. Biomaterials 2002;23:2587–92. [40] Tsuji H, Nakano M, Hashimoto M, Takashima K, Katsura S, Mizuno A. Electrospinning of poly(lactic acid) stereocomplex nanofibers. Biomacromolecules 2006;7:3316–20. [41] Tormala P. Biodegradable self-reinforced composite materials; manufacturing structure and mechanical properties. Clin Mater 1992;10:29–34. [42] Tormala P, Pokkanen P, Vainionpaa S, Laiho J, Heponen VP, Pohjonen T. Material for osteosynthesis devices. US Patent 4,743,257; 1988. [43] Vainionpää S, Kilpikari J, Laiho J, Helevirta P, Rokkanen P, Törmälä P. Strength and strength retention in vitro, of absorbable, self-reinforced polyglycolide (PGA) rods for fracture fixation. Biomaterials 1987;8:46–8. [44] Majola A, Vainionpää S, Pollanen P, Millola HM, Törmälä P. Absorbable self-reinforced polylactide (SR-PLA) composite rods for fracture fixation: strength and strength retention in the bone and subcutaneous tissue of rabbits. J Mater Sci—Mater Med 1992;3:43–7.
38
[45] Wright-Charlesworth D, Miller DM, Miskioglu I, A. KJ. Nanoindentation of injection molded PLA and self-reinforced composite PLA after in vitro conditioning for three months. J Biomed Mater Res A 2005;74A:388–96. [46] Yu L, Liu H, Xie F, Chen L, Li X. Effect of annealing and orientation on microstructure and mechanical properties of PLA. Polym Eng Sci 2008;48:634–41. [47] Yu L, Liu H, Dean K, Chen L. Cold crystallization and post-melting crystallization of PLA plasticized by compressed carbon dioxide. J Polym Sci Part B, Polym Phys 2008;46:2630–6. [48] Kim JY, Kim YW, Lee JG, Cho KS. Effect of annealing on mechanical properties of self-reinforced alpha-silicon carbide. J Mater Sci 1999;34:2325–30. [49] Hine PJ, Ward IM. Hot compaction of woven poly(ethylene terephthalate) multifilaments. J Appl Polym Sci 2004;91:2223–33. [50] Rojanapitayakorn P, Mather PT, Goldberg AJ, Weiss RA. Optically transparent selfreinforced poly(ethylene terephthalate) composites: molecular orientation and mechanical properties. Polymer 2005;46:761–73. [51] Zhang JM, Reynolds CT, Peijs T. All-poly(ethylene terephthalate) composites by film stacking of oriented tapes. Compos Part A—Appl S 2009;40:1747–55. [52] Duhovic M, Bhattacharyya D, Fakirov S. Nanofibrillar single polymer composites of poly(ethylene terephthalate). Macromol Mater Eng 2010;295:95–9. [53] Gilbert JL, Ney DS, Lautenschlager EP. Self-reinforced composite poly(methyl methacrylate): static and fatigue properties. Biomaterials 1995;16:1043–55. [54] Wright DD, Lautenschlager EP, Gilbert JL. Bending and fracture toughness of woven self-reinforced composite poly(methyl methacrylate). J Biomed Mater Res A 1997;36:441–53.
39
[55] Wright DD, Lautenschlager EP, Gilbert JL. Interfacial properties of self-reinforced composite poly(methyl methacrylate). J Biomed Mater Res B 1998;43:153–61. [56] Wright DD, Gilbert JL, Lautenschlager EP. The effect of processing temperature and time on the structure and fracture characteristics of self-reinforced composite poly(methyl methacrylate). J Mater Sci—Mater Med 1999;10:503–12. [57] Wright-Charlesworth DD, Lautenschlager EP, Gilbert JL. Hot compaction of poly(methyl methacrylate) composites based on fiber shrinkage results. J Mater Sci— Mater Med 2005;16:967–75. [58] Wright-Charlesworth DD, Peers WJ, Miskioglu I, Loo LL. Nanomechanical properties of self-reinforced composite poly(methyl methacrylate) as a function of processing temperature. J Biomed Mater Res A 2005;74A:306–14. [59] Li M, Lu S. Material for repairing defective tissue of body and its preparing process. PRC Patent 02138127.5; 2002. [60] Wollerdorfer M, Bader H. Influence of natural fibres on the mechanical properties of biodegradable polymers. Ind Crop Prod 1998;8:105–12. [61] Lu X, Zhang MQ, Rong MZ, Shi G, Yang GC. Self-reinforced melt processable composites of sisal. Compos Sci Technol 2003;63:177–86. [62] Nishino T, Matsuda I, Hirao K. All-cellulose composite. Macromolecules 2004;37:7683–87. [63] Gindl W, Keckes J. All-cellulose nanocomposite. Polymer 2005;46:10221–25. [64] Gindl W, Keckes J. Drawing of self-reinforced cellulose films. J Appl Polym Sci 2007;103:2703–8. [65] Liu H, Yu L, Xie F, Chen L. Gelatinization of cornstarch with different amylose/amylopectin content. Carbohydrate Polymer 2006;65:357–62.
40
[66] Liu H, Yu L, Simon G, Zhang X, Dean K, Chen L. Effect of annealing and pressuring on microstructure of corn starches with different amylose/amylopectin ratios. Carbohyd Research 2009;344:350–4. [67] Chen P, Yu L, Simon G, Petinakis E, Dean K, Chen L. Morphologies and microstructure of starch studied by CLSM. J Cereal Sci 2009;50:241–7. [68] Xie F, Yu L, Su B, Liu P, Wang J, Liu H, Chen L. Rheological properties of starches with different amylose/amylopectin ratios. J Cereal Sci 2009;49:371–7. [69] Liu H, Yu L, Simon G, Dean K, Chen L. Effect of annealing on gelatinization and microstructure of corn starches with different amylose/amylopectin ratios. Carbohyd Polym 2009;77:662–9. [70] Liu P, Yu L, Wang X, Li D, Chen L, Li X. Glass transition temperature of starches with different amylose/amylopectin ratios. J Cereal Sci 2010;51:388–91. [71] Liu H, Li M, Chen P, Yu L, Chen L, Zhen T. Morphologies and thermal behaviours of hydroxypropylated higher amylose starch. Cereal Chem 2010;87:144–9. [72] Arvanitoyannis I. Totally and partially biodegradable polymer blends based on natural and synthetic macromolecules: preparation, physical properties, and potential as food packaging materials. Polym Rev 1999;39:205–71. [73] Arvanitoyannis I, Nakayama A, Aiba S. Edible films made from hydroxypropyl starch and gelatin and plasticized by polyols and water. Carbohyd Polym 1998;36:105–19. [74] Yu L, Christie G. Microstructure and mechanical properties of orientated thermoplastic starches. J Mater Sci 2005;40:111–6. [75] Dean K, Yu L, Wu DY. Preparation and characterization of melt-extruded thermoplastic starch/clay nanocomposites. Compos Sci Technol 2007;67:413–21. [76] Liu HS, Xie FW, Yu L, Chen L, Li L. Thermal processing of starch-based polymers. Prog Polym Sci 2009;34:1348–68.
41
[77] Reddy N, Yang YQ. Preparation and properties of starch acetate fibers for potential tissue engineering applications. Biotechnol Bioeng 2009;103:1016–22. [78] Pashkuleva I, Marques AP, Vaz F, Reis RL. Surface modification of starch based blends using potassium permanganate-nitric acid system and its effect on the adhesion and proliferation of osteoblast-like cells. J Mater Sci—Mater Med 2005;16:81–92. [79] Mano JF, Koniarova D, Reis RL. Thermal properties of thermoplastic starch/synthetic polymer blends with potential biomedical applicability. J Mater Sci—Mater Med 2003;14:127–35. [80] Lan C, Yu L, Chen P, Chen L, Zou W, Simon G, Zhang XQ. Design, preparation and characterization of self-reinforced starch films through chemical modification. Macromol Mater Eng 2010;295:1025–30. [81] Rindlav-Westling A, Stading M, Hermansson AM, Gatenholm P. Structure, mechanical and barrier properties of amylose and amylopectin films. Carbohydr Polym 1998;36:217–24. [82] Fan HJ, Gu Y, Xie ML. In-situ self reinforced polyimide for film materials. Polym Int 2001;50:1331–7. [83] Morgan SE, Misra R, Jones P. Nanomechanical and surface frictional characteristics of a copolymer based on benzoyl-1,4-phenylene and 1,3-phenylene. Polymer 2006;47;2865–73.
42
Figures captions
Fig. 1. Typical PE-based SRC: (a) a woven cloth made from multifilament bundles; (b) permanganic etched image of compacted woven melt-spun multifilaments. Reproduced from [2] with permission, Elsevier, copyright 2008.
Fig. 2. Schematic representation of PLA-based SRC produced by partially sintering the fibers together at high temperature and pressure.
Fig. 3. Processing scheme for hot compaction processing of PET-based SRC: (a) commercial PET fiber yarn; (b) fiber wound on open square frame; (c) compaction of wound fibers; (d) schematic processing profile. Reproduced from [45] with permission, Elsevier, copyright 2005.
Fig. 4. Oxidized starch particles in (a) water suspension and (b) in hydroxypropylation starch matrix observed by confocal scanning laser microscopy. Reproduced from [75] with permission, Wiley, copyright 2010.
43
Fig. 1. Typical PE-based SRC: (a) a woven cloth made from multifilament bundles; (b) permanganic etched image of compacted woven melt-spun multifilaments. Reproduced from [2] with permission, Elsevier, copyright 2008.
44
Matrix (m) Fibre (f) m f
Fig. 2. Schematic representation of PLA-based SRC produced by partially sintering the fibers together at high temperature and pressure.
45
Fig. 3. Processing scheme for hot compaction processing of PET-based SRC: (a) commercial PET fiber yarn; (b) fiber wound on open square frame; (c) compaction of wound fibers; (d) schematic processing profile. Reproduced from [45] with permission, Elsevier, copyright 2005.
46
(a)
(b)
Fig. 4. Oxidized starch particles in (a) water suspension and (b) in hydroxypropylation starch matrix observed by confocal scanning laser microscopy. Reproduced from [75] with permission, Wiley, copyright 2010.
47
