Progress in Polymer Science 35 (2010) 1288–1310

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Progress in Polymer Science

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Self-reinforced polymeric materials: A review

Ákos Kmettya, Tamás Báránya,∗, József Karger-Kocsisa,b

a Department of Polymer Engineering, Budapest University of Technology and Economics, H-1111 Budapest, Muegyetem rkp. 3, Hungaryb Polymer Technology, Faculty of Engineering and Built Environment, Tshwane University of Technology, Pretoria 0001, Republic of South Africa

a r t i c l e i n f o

Article history:Received 4 December 2009Received in revised form 19 April 2010Accepted 21 June 2010Available online 7 August 2010

a b s t r a c t

The preparation, properties and applications of self-reinforced polymeric materials(SRPMs), representing an emerging family of composite materials, are surveyed. SRPMsare classified according to their constituents (single- or multi-component), their produc-tion (in one-step (in situ) or in multi-step (ex situ) procedures) and the spatial alignment ofthe reinforcing phase in the matrix (in one-, two- or three-dimensions; 1D, 2D or 3D, respec-tively). The pros and cons of the related processes and products are introduced, and further

Keywords:PolymersSelf-reinforcedProcessingMicrostructureStructure–property relationshipsSemicrystalline polymers

developments with SRPMs are highlighted. The driving forces in the development of SRPMsinclude possibilities for recycling (i.e., reprocessing via remelting) and the fabrication oflightweight structures.

© 2010 Elsevier Ltd. All rights reserved.

Application

Contents

1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12882. Self-reinforced polymeric materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1291

2.1. Single-component SRPMs. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12912.1.1. One-step (in situ) production . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12912.1.2. Multi-step (ex situ) production . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1294

2.2. Multi-component SRPMs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13012.2.1. Single-step production . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13012.2.2. Multi-step production . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1301

3. Outlook and future works . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1307Acknowledgements . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1308References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1308

∗ Corresponding author. Tel.: +36 1 463 3740; fax: +36 1 463 1527.E-mail address: barany@eik.bme.hu (T. Bárány).

0079-6700/$ – see front matter © 2010 Elsevier Ltd. All rights reserved.doi:10.1016/j.progpolymsci.2010.07.002

1. Introduction

Presently, considerable research activities and accom-

panying commercial interest are devoted to all-polymericmaterials, especially to self-reinforced versions. In all-polymeric materials, both the reinforcing and matrixphases are given by suitable polymers. In self-reinforced

Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310 1289

Nomenclature

ac,n impact strength (notched) (kJ/m2)ac,n,L longitudinal impact strength (kJ/m2)ac,n,T transversal impact strength (kJ/m2)D extruder screw diameter (mm)EB E-modulus (flexural) (GPa)Et E-modulus (tensile) (GPa)Et,L E-modulus (tensile and longitudinal) (GPa)Et,T E-modulus (tensile, transversal) (GPa)k thermal conductivity (W/mK)L extruder screw length (mm)p processing pressure (MPa)pA pressure amplitude (MPa)TD drawing (stretching) temperature (◦C)Tg glass transition temperature (◦C)Tproc processing temperature (◦C)Tproc,opt optimal processing temperature (◦C)Tm melting temperature (◦C)Tmelt melt temperature (◦C)Tmold mold temperature (◦C)V extrusion velocity (mm/min)� drawing ratio�B tensile strength (MPa)�B,L longitudinal tensile strength (MPa)�B,T transversal tensile strength (MPa)�F flexural strength (MPa)�Y yield strength (MPa)BP based pressure for VIM (MPa)CBT cyclic butylene terephthalate oligomerCM conventional injection moldingCNF carbon nanofiberCP cross-ply structureDMA dynamical mechanical analysisDSC differential scanning calorimetryEP ethylene-propylene copolymerEPR ethylene-propylene rubberGF glass fiberHDPE high-density polyethylenehPP polypropylene homopolymeriPP isotatic polypropyleneLCP liquid crystalline polyesterMD machine directionMFC microfibrillar compositeOPIM oscillating packing injection moldingPA polyamidePA-12 polyamide-12PA-6 poliamid-6PA-6.6 poliamid-6.6PBA polybutyl acrylatePCTG polycyclohexane-terephthalate glycolPE polyethylenePEEK polyether-ether-ketonePEN polyethylene-naphthalatePET polyethylene-terephthalatePETG polyethylene-terepthalate glycolPMMA polymethyl-methacrylatePOM polyoxymethylene or polyacetalPP polypropylene

PPS polyphenylene-sulfidePS polystyrenePTFE polytetrafluoro-ethylenePVC polyvinyl-chloridePVDF polyvinylidene-fluoriderPP random polypropylene copolymerSCORIM shear-controlled orientation injection

moldingSEM scanning electron microscopySRPM self-reinforced polymeric materialSRPP self-reinforced polypropyleneTD transverse (to machine) directionTEM transmission electron microscopyTMA thermo mechanical analysisUD unidirectional alignment, structureUHMPE ultra high modulus polyethyleneUHMWPE ultra high molecular weight polyethyleneVF vibration frequency (Hz)VIM vibration injection moldingVPA vibration pressure amplitude (MPa)�-PP isotactic PP (alpha form)�-rPP random polypropylene copolymer (alpha

form)�-PP isotactic PP (beta form)

polymeric materials (SRPMs), the same polymer formsboth the reinforcing and matrix phases. SRPMs are alsoreferred to as single-phase or homocomposites. Moreover,in the open literature, such polymer composites in whichthe reinforcement and matrix polymers are different butbelong to the same family of polymers (see below) are alsotermed as SRPMs.

SRPMs may compete with traditional composites in var-ious application fields based on their performance/costbalance. With respect to their performance, the ease ofrecycling has to be emphasized because they representlikely the best recycling option when reprocessing viaremelting is targeted. Accordingly, SRPMs can be consid-ered to be environmentally benign materials. The conceptsused to produce SRPMs can also be adapted to biodegrad-able polymers to improve their property profiles, wherebyeven the degradation properties can be tailored uponrequest.

A further driving force for SRPMs is the possibil-ity of manufacturing lightweight parts and structuresbecause the density of SRPMs is well below thoseof traditional filled polymers. The density of the cor-responding composite is usually higher than that ofan SRPM because the former contain reinforcementssuch as glass (density: 2.5–2.9 g cm−3), carbon (density:1.7–1.9 g cm−3), basalt (density: 2.7–3.0 g cm−3), aramid(density: 1.38–1.44 g cm−3) and/or fillers like talc (density:2.7–2.8 g cm−3), chalk (density: 1.1–2.5 g cm−3) and silica

−3

(density: 2.1–2.6 g cm ).The basic concept of self-reinforcement is the creation

of a one-, two- or three-dimensional alignment (1D, 2Dor 3D, respectively) within the matrix to fulfill the roleof matrix reinforcement. Reinforcing action requires that

1290Á

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Polymer

Science35

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Fig. 1. Classification of self-reinforced polymeric materials (SRPMs), *not yet explored.

olymer S

tsttfadasss

Smsi1fo“attlbfcbfiSrfseowwpc

rPta(tri

bawiepmfibsc

Á. Kmetty et al. / Progress in P

he generated structure possesses a higher stiffness andtrength than the matrix and, in addition, is well “bonded”o the matrix polymer. As a consequence, the stress can beransferred from the “weak” matrix to the “strong” rein-orcing structure, according to the “working principle” ofll composites. The reinforcing structure can be produceduring one (in situ) or more processing steps (ex situ). This,long with the spatial alignment of the reinforcement, mayerve for their straightforward classification, with one suchcheme given in Fig. 1 to serve as the guide in this articleurveying SRPMs.

From a historical point of view, the development ofRPMs started with the in situ production of 1D-reinforcedaterials. This occurred mostly with solid-phase extru-

ion forming and techniques exploiting melt shearingn solidifying melts. The related operations resulted inD, aligned supermolecular structures acting as the rein-orcement in SRPMs, whereby the covalent bond strengthf the macromolecules was indirectly utilized. The termsupermolecular structure” already suggests that SRPMsre almost exclusively semicrystalline polymer-based sys-ems. Compared to the matrix, the reinforcing structure inhem either has different crystalline and/or supermolecu-ar (also referred to as higher-order) structures or is giveny a preform, a prefabricate (e.g., fiber, tape and their dif-erent textile architectures) with higher (and different)rystallinity. The 1D reinforcement can also be generatedy multi-step stretching. Related technologies, practicedor example in fiber spinning operations, are groupednto multi-step (ex situ) productions of single-componentRPMs (cf., Fig. 1). According to our terminology, self-einforced polypropylene (PP) composites, produced onlyrom fibers or fabrics as preforms, are classified as multi-tep products of single-component SRPMs, which mayxhibit 1D (unidirectional fiber alignment), 2D (fabric plies)r 3D (e.g., braided structure) reinforcements. In otherords, when an SRPM is produced solely from preforms,ith prefabricates instead of primary granules of a givenolymer, it is classified as the product of multi-step pro-essing.

A commercial break-through with SRPMs occurredecently. Self-reinforced PP composites (also called all-P composites) are now available on the market underhe trade names Curv®, Pure® and Armordon®. Curv® issingle-component, multi-step product usually with 2D

fabric) reinforcement, whereas Pure® and Armordon® arewo-component, multi-step versions originally with 1Deinforcement (as stretched tapes with different PP gradesn the core and surface layers).

As mentioned before, the majority of SRPMs areased on semicrystalline polymers. On the other hand,morphous matrix-based SRPMs can also be createdhen interpreting “self-reinforcement” in a broader sense,

.e., extending it for a given family of polymers. Forxample, amorphous copolyesters can be reinforced byolyethylene-terephthalate (PET) fibers, thereby retaining

elt reprocessability. It is noteworthy that transesteri-

cation reactions in the melt (which can be triggeredy additional additives) guarantee the necessary adhe-ion between the reinforcing PET and the amorphousopolyester matrix. In other SRPM versions, the differ-

cience 35 (2010) 1288–1310 1291

ence in the melt temperature of semicrystalline polymersbelonging to the same family is exploited. One can findreports on such all-PA or self-reinforced PA materials forwhich the constituents are different PAs; for example, thematrix is given by PA-12, with a low melting temperature,whereas the reinforcement is from a higher melting PA,such as PA-6 or PA-6.6.

2. Self-reinforced polymeric materials

The grouping outlined in Fig. 1 will be followed inthis chapter. Accordingly, the single- and multi-componentSRPMs will be treated separately by considering their pro-duction (i.e., in situ or ex situ) and spatial reinforcingstructure (i.e., 1D, 2D or 3D).

2.1. Single-component SRPMs

2.1.1. One-step (in situ) productionA 1D self-reinforcing structure can be produced by

extrusion molding whereby the extruder is equipped with adie having a convergent section (cf., Fig. 2). The convergentsection (with an angle of 45◦ or higher) is foreseen to gen-erate the molecular orientation via extensional flow that is“frozen” in the subsequent sections of the die (calibrationzone). Pornnimit and Ehrenstein [1] used this techniqueto manufacture self-reinforced HDPE. It was shown thatthe oriented molecules act as (self) row nuclei and triggerthe formation of cylindrical and shish-kebab-type super-molecular structures (Fig. 2). As a controlling parameterof the formation of the self-reinforcement, the tempera-ture program of the die (affecting the pressure build-upwithin) was identified. Upon cooling the outcoming zoneof the die, a high extrusion pressure could be reached,which supported the formation of the shish-kebab crys-tals. The self-reinforced HDPE rod exhibited considerablyhigher stiffness and strength and highly reduced thermalshrinkage when measured in the reinforcing direction.

Although the shish-kebab structure has been knownsince the mid 1960s, the mechanism of its formation is stilldebated. Kornfield et al. [2] in their recent work proposedthat long chains are not the dominant species of the shishformation as thought before. Nevertheless, the presenceof long macromolecules strongly favors the propagation ofshish.

The basic prerequisites of the extrusion procedureyielding 1D self-reinforcement were identified as fol-lows [3]: molecular orientation in the melt via forcedextensional flow; processing close to the crystallizationtemperature of the polymer; and “fixing” of the result-ing structure in the final section of the die by raising thepressure. DSC investigations showed that the melting peakof the self-reinforced HDPE was shifted towards highertemperatures (by ca. 4 ◦C). The thermal stability of the ori-ented crystals could be well detected by polarized opticalmicroscopy, during which different fusion temperatures

were set prior to the subsequent crystallization steps.

Farah and Bretas [4] developed shear-induced crystal-lization layers in iPP via a slit die attached to a twin-screwextruder. The output rate was below 10 kg/h. The die tem-peratures were set between 169 and 230 ◦C. Rheological

1292 Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310

tion in a

Fig. 2. Scheme of the 1D supermolecular structure forma

studies revealed that the induction time, at a given crystal-lization temperature, decreased as the shear rate increased.At a given shear rate, higher crystallization temperaturesgave longer induction times. It was observed that at a givenoutput rate, the thickness of the shear-induced crystallinelayer decreased with the increase of die temperature. Threelayers were found by SEM and TEM. Two layers werespherulitic while one layer was composed of highly ori-ented lamellae.

Huang et al. [5,6] produced self-reinforced HDPE byusing a convergent die (angle 60◦) and an extrusion pres-sure ranging from 30 to 60 MPa. Similar to the methods in[7], the authors cooled the melt before leaving the die at128 ◦C. The strength of the resulting 1.5-mm-thick sheetswas eight times higher than that of the conventionallyextruded sheet. The anisotropy in the sheets was detectedin mechanical and tribological tests and was also demon-strated by microhardness results.

Parallel to the works on PEs, PP was also discovered as asuitable candidate for SRPM [8,9]. Song et al. [10] producedself-reinforced PP by a conventional single-screw extruderwith pressure regulation (L/D = 30, maximum pressure:100 MPa), equipped with a convergent die (entrance angle45◦) with two or more calibration (cooling) sections. Theproperties of the extrudate were superior to counterpartsproduced by the conventional extrusion molding.

Self-reinforced structures can also be generated byinjection molding. The related techniques differ from oneanother whether the oriented structure is created outsideor within the mold. Prox and Ehrenstein [11] produced self-reinforced material using the technique of converging dieinjection molding. They injected the low temperature melt

into the cavity just after the melt passed a convergent diesection. Note that this concept requires a careful mold con-struction and well-defined processing conditions to avoidrelaxation phenomena reducing the molecular orientation.

die with a convergent section during extrusion molding.

Those injection molding techniques that generate theself-reinforcement in the mold have become far more pop-ular than the above-mentioned variant. They are knownunder shear-controlled orientation in injection molding(SCORIM)) [12,13] or oscillating packing injection mold-ing (OPIM). The common characteristic of these techniquesis that the molecular orientation is set in the mold byshearing/oscillation of the solidifying melt via a suitablearrangement of pistons. The pistons start to work whenthe cavity is already filled. The related mold constructionmay be very different [14], although in SCORIM three basicoperation modes exist (cf., Fig. 3).

Guan et al. [16] used the OPIM to produce self-reinforcedHDPE. An injection pressure of ca. 41 MPa was superim-posed by an oscillating packing pressure (varied between32 and 48 MPa) with a frequency of 0.3 Hz. An operationmode “A” in Fig. 3 was chosen, and 220 and 42 ◦C were setfor the temperatures of the melt and mold, respectively.The molded parts were subjected to mechanical and mor-phological tests. The stiffness and strength of the OPIMmoldings were superior to the conventional ones. Mor-phological studies revealed the presence of a microfibrillarstructure. The TEM study showed that the microfibrillarstructure was composed of shish-kebab formations. Basedon DSC measurements, the authors concluded that themicrospherulitic structure melts at 132 ◦C, whereas theshish-kebab crystals melt at 137 ◦C. In a follow-up work,Guan et al. [17] adapted the OPIM on PP. Studying theeffects of processing conditions, the authors concluded thatthe mechanical properties of the moldings strongly dependon the operation mode and to a lesser extent depend on theoscillation frequency, frequency/mode and frequency/time

combinations.

Chen and Shen [18] produced biaxial self-reinforced(i.e., 2D) PP by OPIM. An operation mode “A” in Fig. 3 wasselected, and 195 and the range of 20–80 ◦C were chosen

Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310 1293

F three bao n a conS

fTspdttm4

ttmbetito

Fr

ig. 3. Scheme of the function of the SCORIM procedure along with thef phase; Mode B: pistons are activated in phase; the pistons are held dowons.

or the temperatures of the melt and mold, respectively.he products exhibited quite balanced (i.e., less anisotropy)tatic mechanical properties (strength improvements com-ared to conventional injection molding in the melt flowirection and transverse to it at 55–70 and 40%, respec-ively), but further on a pronounced anisotropy in respecto impact strength was seen (improvement to conventional

olding in the melt flow direction and transverse to it at00 and 30–40%, respectively).

Kalay et al. [15,19] investigated the influence of PPypes on the corresponding SCORIM products and deducedhe basic rules on how to prepare products with opti-

um properties. It is important to emphasize that theasic advantage of SCORIM/OPIM is the pronounced ori-

ntation of the molecules in the whole cross-section ofhe molded parts. This is because of the repeated shear-ng/oscillation movements in the melt that are acting untilhe melt solidifies. This suppresses the relaxation of theriented molecules.

ig. 4. Working principle of the vibration injection molding [20]. Copy-ight 2008 with permission from Taylor & Francis.

sic operations (A, B and C) – Mode A: the pistons are activated 180◦ outstant pressure [15]. Copyright 1997 with permission from John Wiley &

A further variant of the injection molding resulting inself-reinforcement is vibration injection molding (VIM),which was pioneered by Li et al. [20]. The working principleof VIM is depicted in Fig. 4. The ram itself is a part of both theinjection and vibration systems. Without vibration, the set-up works as a conventional injection molding (CM) unit.However, working in the VIM mode, pulsations occur inthe injection and holding pressure stages. This causes aneffective compression and decompression of the melt andshearing at the melt–solid interface. Note that solidificationprogressed from the surface to the core of the molding inthe cavity. For this VIM device, the main processing param-eters are vibration frequency (VF) and vibration pressureamplitude (VPA).

In the cited study, the authors used a single screwextruder as the plastification unit. The PP melt was vibratedfor 25 s, and the cooling time was fixed at 20 s. The injec-tion pressure for CM and the base pressure (BP) for VIMwas 49.4 MPa. In the latter case, the pressure amplitude wasfixed at 19.8 MPa. The mechanical properties and morphol-ogy of the specimens were determined. It was found thatthe mechanical properties of the VIM-produced parts wereenhanced compared to conventional injection molding.The yield strength steeply rose with the vibration fre-quency in the range of 0–1 Hz. Afterwards, a constant valuewas noticed for the range of 1–2.5 Hz. The tensile strengthincreased with increasing vibration frequency. The impactstrength of PP was doubled compared to the conventionalmolding using VIM at 2.33 Hz. The crystalline structure ofthe VIM-produced PP showed the simultaneous presenceof the crystalline alpha, beta- and gamma-modificationsof PP. In a companion study [21] using HDPE and setting

the vibration frequency at 2.33 Hz and the pressure ampli-tude at 19.8 MPa, the authors observed the formation ofa shish-kebab along with row-nucleated crystalline lamel-lae. Their presence resulted in an upgrade of the mechanicalproperties of HDPE.

olymer S

1294 Á. Kmetty et al. / Progress in P

Attention should be paid to a widely practiced designmethod in injection molded items, to the film or “plas-tic” hinge. It was recognized early-on that the convergent(hinge) section of the molded parts of both semicrys-talline and amorphous thermoplastics has a peculiarperformance: it withstands multiple bending movements.Now, this design principle has been incorporated intomany products of everyday life, especially for dispens-ing packages. Morphological studies on such hinges [22]demonstrated the presence of strongly oriented (1 or 2D)supermolecular structures, including shish-kebab types.The hinges consist of two highly oriented surface layersand one almost isotropic core in between. The core exhibitsa small-sized spherulitic structure whereas the orientedsurface layers contain shish-kebab structures. The mechan-ical behavior of the oriented layers is similar to that of“hard elastic fibers”, which show a high stiffness and a highstrain recovery. So, products with film or ‘plastic’ hingesrepresent nice examples of the one-step- (in situ) pro-duced SRPCs, although only a given section of them is reallyself-reinforced. Some results of the previously presentedmethods are summarized in Table 1 – cf. Fig. 1.

2.1.2. Multi-step (ex situ) productionSingle-component SRPCs can also be produced by multi-

step production methods, such as die- and zone-drawing,ram extrusion, hydrostatic extrusion, rolling (using vari-ous solid “preforms” that are eventually produced on-line),gel drawing or spinning (where the “preform” is a dilutepolymer solution). As the reader will see, in many casesthe preparation of the “preform” and the generation of thereinforcing structure within occur on-line, but in differentstages or steps. This is the reason why they are listed amongthe multi-step production methods. When the orientationand thus the creation of the reinforcing structure take placein the solid state of the polymer (i.e., below its melting tem-perature), the related methods are referred to as solid-stateprocesses [10].

2.1.2.1. Solid phase extrusion. Developed in the early1970s, ram extrusion involves the pressing of a solid pre-form through a metallic die of conical (convergent) shape.This technique was successfully adapted to many ther-moplastics, covering not only semicrystalline (PE, PTFE,PP, PET, and PA) but also amorphous versions (PS) [24].Major problems with the ram extrusion include: a verylow output rate due to the very high friction betweenthe solid polymer and the die surface and the coexis-tence of different morphological superstructures throughthe cross-sections of the extrudate [25]. Legros et al. [26]studied the effects of the processing conditions (additionaluse of lubricant, variation in the extrusion speed, use ofa take-up device) of the ram extrusion on the propertiesof HDPE and PP rods. The experiments were performed ata barrel area/die exit area ratio of 6. The maximum drawratio, � ∼ 6, was obtained with a low extrusion speed of

0.1 mm/s. At higher speeds, like at high extrusion tem-peratures, � was markedly reduced for PP. For HDPE,the decrease in the draw ratio as a function of experi-mental conditions was less pronounced than for PP. Anincreasing draw ratio was accompanied with enhanced

cience 35 (2010) 1288–1310

crystallinity, as expected. By the take-up device, the relax-ation phenomena in the rod, after leaving the die, couldbe reduced. Note that this technique is nowadays wellestablished for the manufacturing of various PTFE-basedproducts.

Using hydrostatic extrusion [24], some drawbacks ofram extrusion can be circumvented. For example, theextrudate has a homogeneous reinforcing structure. In thisprocess the polymer preform is pressed with the help ofa hydraulic fluid through a conical die and the outcomingextrudate is pulled away (cf., Fig. 5).

The hydrostatic extrusion was successfully adapted tomanufacture high-modulus tapes and fibers even fromfilled (hydroxyapatite/PE) and reinforced polymers (dis-continuous glass fiber-reinforced POM). Disadvantages ofthis process include discontinuous operation and the veryhigh flow stress at the exit of the die. The polymer has thehighest strain rates at the exit of the conical die, wherethe plastic strain is the greatest. The strain-rate sensitiv-ity of flow stress in solid-state extrusion increases rapidlywith plastic strain. As this situation incurs very high flowstresses as the polymer reaches the die exit, high extrusionpressures are therefore required [24].

The die-drawing, credited to Ward et al. [24], is a furtherdevelopment in this field. The change in the morphol-ogy due to the die/drawing is depicted schematically inFig. 6.

The advent of the die-drawing is that the draw ratio canbe set accordingly. This technique was used for differentpolymers, like PE [27], PP [28,29], PVC [24], PET [24], PEEK[24], PVDF [24] and POM [30]. Owing to the high molecu-lar orientation, the related products exhibited pronouncedimprovements in the E-modulus, strength, barrier and sol-vent resistance. In addition, the extrudates were less proneto creep than the conventionally produced counterparts.This method is used to produce PE (gas, water) and PVCpipes (drainage) and PET containers (food storage) [24].

2.1.2.2. Super drawing. A two-stage drawing techniquewas applied to the super-drawing of PTFE virgin pow-der by Endo and Kanamoto [31]. In the first-stage,the compression-molded PTFE film was solid-state co-extruded (extrusion draw ratio (EDR) between 6 and 20)at 10 ◦C below the Tm. The second-stage draw was made byapplying a pin-draw technique in the temperature rangecovering the static Tm of PTFE. The maximum achievedtotal draw ratio was 160. The maximum tensile modulusand strength at 24 ◦C reached 102 ± 5 and 1.4 ± 0.2 GPa,respectively.

2.1.2.3. Rolling. Rolling processes can induce a permanentdeformation in the morphology by transforming the ini-tial spherulitic structure to a fibrillar structure. This canbe achieved by series of pairs of rolls (heated or not)and temperature-conditioning steps. Rolling is usuallypreferred for semicrystalline instead of amorphous poly-

mers because the latter show more pronounced relaxationbehavior [24,32]. PE and PP are used for room tempera-ture rolling, whereby a thickness reduction ratio of up to∼5:1 can be reached. At high speeds (as high as 20 m/min),rolling occurs adiabatically. As a consequence, the chemical

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Table 1Production method, conditions and product characteristics of single-component SRPMs produced in one-step (in situ).

Num. Processing Materials Processing conditions Results Comment Ref.

1 Self-reinforced extrusionHDPE

Entrance semi-angel of 45◦ ,p = 60–100 MPa, Tmold = 180 ◦C

�B = 160 MPa, Et = 2–17 GPa Fibrillar structure (next to the die wall),Shish-kebab structure

[1]

Entrance semi-angel of 60◦ ,p = 30–60 MPa, Tmold = 128 ◦C

�B = 130–192 MPa, Fibrillar structure, Micro hardness and lightpermeability are increasing

[5,6,23]

PP p = 30–70 MPa, Tmold = 160 ◦Cv = 160–200 mm/min

�B = 60 MPa, Et = 3.3 GPa – [10]

2 Self-reinforced injection molding hPP Tmelt = 160–190 ◦C, Tmold = 25–80 ◦C, �B = 62–95 MPa, Et = 2–3 GPa – [10,11]

3 SCORIM/OPIM

HDPE

Tmelt = 140–180 ◦C,pproc = 25–28 MPaTmold = 40–60◦C

�B,L = 24–29 MPa, ac,N,L = 5–6 kJ/m2,�B,T = 25–30 MPa, ac,N,T = 5–6 kJ/m2

Static packing[14]

�B,L = 36–56 MPa, ac,N,L = 8–14 kJ/m2

�B,T = 23–36 MPa, ac,N,T = 2–3 kJ/m2Dynamic packing (f = 0.2–0.5 Hz)

pproc = 32–48 MPa, Tmelt = 220 ◦C,Tmold = 42 ◦C

�B = 70–90 MPa, Et = 3–6 GPa – [16]

PPTmold = 42 ◦C,Tmelt = 210–240 ◦C,

�B = 34–35 MPa Static packing[17]

�B = 47–53 MPa Dynamic packing (f = 0.3–1 Hz)iPP Tmold = 20–80 ◦C, pproc = 40 MPa, �B,L = 44–55 MPa, ac,N,L = 7–13 kJ/m2

�B,T = 35–45 MPa, ac,N,T = 2–3 kJ/m2Dynamic packing f = 0.125–0.5 Hz [18]

PP copolymer pproc = 100 MPa, Tmold = 60–110 ◦C,Tmelt = 220–250 ◦C,

�y = 50–59 MPa, Et = 2–3 GPa Different packing mode[15,19]

iPP pproc = 100–160 MPa, Tmold = 60 ◦C,Tmelt = 205–250 ◦C

�y = 38–77 MPa, Et = 2–3 GPa Different packing mode

4 VIMiPP pproc = 100 MPa, Tmold = 40 ◦C,

Tmelt = 210–230 ◦C, f = 0–2.5 Hz,pA = 0–73.5 MPa

�y = 32–38 MPa, ac,n = 11–32 kJ/m2 – [20]

HDPE pproc = 39.5 MPa, Tmold = 40 ◦C,Tmelt = 180–200 ◦C, VF = 0–2.33 Hz,pA = 0–59.4 MPa

Crystallinity = 60–70% Laminar and shish-kebab structure [21]

1296 Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310

schema

Fig. 5. Working principle of the hydrostatic extrusion process

and thermal stability of the polymer should be considered.The rolling process increases the crystalline and amor-phous molecular orientations and thus enhances both thestrength and E-modulus of the polymer [33].

It is well known that the plastic deformation of crys-talline polymers, especially upon drawing, is associatedwith cavitation. Cavitation, however, can be suppressed byapplying compressive stress during orientational drawing.This was demonstrated by Polish researchers, whom devel-oped a method called rolling with side constraints [34–37].The materials used were mostly HDPE and PP.

Galeski [38] reviewed the structure–property rela-tionships in isotactic PP and HDPE produced by rollingwith side constraints. Rolling was done in a speciallyconstructed apparatus at various speeds (0.5–4 m/minfor iPP and 200 mm/min for HDPE) and at differenttemperatures. Both the E-modulus and ultimate ten-sile strength increased with increasing deformation ratio.The maximum strength/deformation ratio values were340 MPa/10.4 and 188 MPa/8.3 for iPP and HDPE, respec-tively.

Mohanraj et al. [39] prepared highly oriented polyacetal(POM) bars via a constrained rolling process. In this process,the heated polymer billet is deformed in a channel given bythe circumference of the bottom roll, which provides lat-eral constraint to the material when it deforms. POM wasrolled below the crystalline melting temperature. The mod-ulus and strength parallel to the rolling direction increased

almost linearly with the compression ratio.

2.1.2.4. Gel drawing. Via gel drawing (spinning), films andfibers can be produced from dilute polymer solutions. Thisrequires, however, a polymer with a high mean molecular

Fig. 6. Working principle of the die-drawing a

tically [24]. Copyright 2000 with permission from the authors.

weight and suitable molecular weight distribution char-acteristics. If the molecules are less entangled in the gel,this guarantees drawing to high degrees [40–42]. Orientedsynthetic fibers of UHMWPE (Dyneema (www.dsm.com)and Spectra (www51.honeywell.com)) can be formed bygel spinning (gel drawing process) to have tensile strengthsas high as 2.8 GPa. These fibers are mostly used to produceballistic vests covers, safety helmets, cut resistant gloves,bow strings, climbing ropes, fishing lines, spear lines forspear guns, high-performance sails, suspension lines inparachutes, etc. (tensile strength of the ballistic materials∼3.5 GPa).

2.1.2.5. Orientation drawing. Elyashevich and coworkers[43,44] prepared high-modulus and high-strength PE fibersvia orientation drawing. Drawing took place between theglass transition (Tg) and melting temperature (Tm) of thegiven polymer. During orientation, the folded chain crystallamellae rotate, beak-up, defold and finally form alignedchain crystals (cf., Fig. 7).

Fibers with very high orientation (draw ratio) were pro-duced in one or more drawing steps. In the latter case, theisothermal drawing temperature was increased from oneto the next drawing step. Elyashevich et al. [43,44] man-ufactured (with one-step orientation) PE fibers having anE-modulus and tensile strength of 35 and 1.2 GPa, respec-tively.

Baranov and Prut [45] produced ultra high modulus

PP tapes by a two-step isothermal drawing process. Theisothermal drawing of the parent film was done in a tensiletesting machine equipped with a thermostatic chamber.The first drawing occurred at 163–164 ◦C, while the secondone was at 165 ◦C. The E-modulus and strength of the tapes

: unoriented phase, b: oriented phase.

Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310 1297

Fig. 7. Scheme of chain orientation.

the exa

wssatact

mMtr

bwaaSaPttdTotisaptd

2np

Fig. 8. Principle sketch of hot compaction on

ere 30–35 and 1.1 GPa, respectively. PP and PET tapes andtrips are widely used for packaging purposes. Their ten-ile strength ranges are 220–350 and 430–570 MPa for PPnd PET, respectively. Morawiec et al. [46] demonstratedhat the strength of PET, even from scrap (recycled bever-ge bottles), may reach 700 MPa when suitable orientationonditions prevail. This was demonstrated using an on-line,wo-step extrusion drawing unit.

The structural “basis” of high-strength and high-odulus polymers is well reviewed by Marikhin andyasnikova [47]. This chapter helps the interested reader

o also trace pioneering activities of researchers in theelated fields.

Alcock et al. [48] produced highly oriented PP tapesy extrusion and drawing steps. The tensile deformationas achieved by pulling a tape from one set of rollers

t 60 ◦C through a hot air oven to a second set of rollerst 160–190 ◦C. The tapes were classified into two series;eries A describes PP tapes drawn to varying draw ratiost the same drawing temperatures, while Series B coversP tapes drawn to � = 13 at a range of drawing tempera-ures in the second drawing stage. The results showed thathe density was approximately constant with an increasingraw ratio up to � = 9.3, above which it sharply dropped.he decrease in density was associated with a change inpacity of the tape due to the onset of microvoiding withinhe tape. Karger-Kocsis et al. [49] noticed that microvoid-ng in stretched iPP tapes takes place even at � = 10. In thetudy of Alcock et al. [48], the density reached 0.73 g cm−3

t � = 17, which indicates an almost 20% reduction com-ared to the undrawn tape. PP tapes possess an ∼15 GPaensile modulus and an ∼450 MPa tensile strength by a high

rawn ratio (� = 17).

.1.2.6. Hot compaction. Ward et al. [50,51] developed aew method to produce SRPCs that they called “hot com-action”. The related research started with highly oriented

mple of unidirectional (UD) arranged fibers.

PE fibers and tapes. When these preforms were put underpressure and the temperature was increased, their surfaceand core showed different melting behaviors. This find-ing was exploited to melt the outer layer of the fibers andtapes, which after solidification (crystallization) becamethe matrix. The residual part of the fibers and tapes (i.e.,their core section) acted as the reinforcement in the result-ing SRPC (cf., Fig. 8).

It was found that hot compaction works well forsemicrystalline, liquid crystalline and amorphous ther-moplastics as well [52]. By hot compaction, differenthigh-strength SRPMs were produced from PET [53,54],PE [55,56], PEN [52], PA-6.6 [57], PPS [52], POM [58], PP[59], PMMA [60] and PEEK [52]. It is intuitive that theprocessing window during the hot compaction of single-component polymeric systems is very narrow. When thecompaction temperature approaches the melting temper-ature of the fiber, the transverse strength of compositeswith UD-aligned (i.e., 1D) reinforcement increases, albeit ata cost to the stiffness and strength measurable in the lon-gitudinal direction [61] (cf., Fig. 9). Fig. 9 also displays thenarrowness of the temperature range for the productionsof SRPMs.

It was also reported that in order to set optimummechanical properties, a given amount of the fiber shouldmelt and work later as the matrix. This was given by ca. 10%of the cross-section (i.e., outer shell) of the fiber. This valueis very closely matched with the amount that is requiredto fill the spatial voids between those fibers that adapt ahexagonal-like cross-section owing to the acting pressure.The hexagonal shaping of the initially spherical fibers alongwith the formation of a transcrystalline layer between the

residual fiber (core) and formed matrix have been proven[61]. Ratner et al. [62] experienced an additional surfacecrosslinking during hot compaction of UHMWPE fibers.The surface of the fibers was coated by a solution contain-ing a peroxide prior to the hot compaction (T = 140–150 ◦C,

1298 Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310

vs. comp

Fig. 9. Longitudinal flexural modulus (�) and transverse strength (�)

pressure: 31 MPa, time: 30 min). In this way, the stresstransfer between the residual fiber (reinforcement) andthe newly formed matrix has been improved comparedto non-treated versions. Here it is appropriate to drawattention to the effect of the transcrystalline layer, whichis controversial from the point of view of fiber/matrixadhesion. Though the development of the transcrystallinelayer is necessary, its internal build-up may be of greatrelevance, as outlined by Karger-Kocsis [63]. Ratner etal. [64] found that the crosslinked interphase betweenfiber and matrix is more beneficial than the usual tran-scrystalline one, especially when long-term propertieslike fatigue are considered. Hine et al. [55] producedSRPMs using fabrics (i.e., 2D reinforcement) composedof high-modulus PE fibers (E-modulus: 42 GPa). Withincreasing hot compaction temperature, it was establishedthat the molten proportion of the fibers increased and thecrystallinity of the formed matrix became markedly lessthan that of the initial fibers. A further important findingwas that the processing temperature for 2D fabrics washigher than for UD- (1D-) aligned fibers. This is becausean assembly of woven fabrics has more interstitial spaceto be filled with the matrix than a parallelized 1D fiberone. The quality of the related SRPM was measured byinterlayer T-peel tests. The T-peel strength increasedsteeply with the matrix fraction (up to 30%) and reacheda constant value afterwards. Based on tensile tests anddetailed morphological studies, the authors quoted thatthe final matrix content should be between 20 and 30% inorder to set optimum properties for SRPMs from wovenfabric layers. It was also emphasized that the processingwindow for 2D fabrics is even smaller than that for 1D

fibers or tapes. However, UHMWPE loses its stiffness andstrength and becomes prone toward creep with increasingtemperature. To overcome this problem, the UHMWPEfibers were exposed to �-irradiation to trigger their cross-linking [52]. Orench et al. [65] performed a comparative

action temperature of melt spun polyethylene fibers (based on [61]).

study on SRPMs produced from commercially availablehigh-strength fibers and tapes (Spectra®, Dyneema®).

Due to the low temperature resistance of PE, the hotcompaction research shifted to PP [59]. This directionyielded new insights, such as that PP should be kept underhigh pressure during heating to the compaction temper-ature to prevent its thermal shrinkage. Hine et al. [66]compacted PP tapes from fibrillated woven PP in both openand closed molds. Based on flexural tests and morpholog-ical inspection, the optimum processing conditions weredefined. Teckoe et al. [67] manufactured 2 mm thick sheetsfrom woven fabrics consisting of high-strength PP fibers.The fabric layers were subjected to a 2.8 MPa pressureuntil the compaction temperature (varying between 166and 190 ◦C) was reached. This temperature was kept for10 min before raising the pressure suddenly to 7 MPa andmaintaining this during cooling to 100 ◦C, when demold-ing took place. At low compaction temperatures, the voidswithin the woven structure were not completely filled,while at high temperatures too much matrix was producedand thus the reinforcement content diminished. It wasclaimed that the final matrix content should be between20 and 30% for good quality products. It is worth notingthat the heating of the related preform to the compactiontemperature is accompanied by the release of its internalstress state. Due to the high pressure applied, the materialmelts under constraint conditions, so its melting occurs ata higher temperature than under normal conditions. This isthe reason why the optimum hot compaction temperatureis close, and even above, the usual melting under uncon-strained conditions. Jordan and coworkers [68–70] studiedthe effects of hot compaction on the performance of PE

and PP tapes and fabrics. The latter differed in their meanmolecular weights, which influenced the consolidationquality assessed by tear tests. Bozec et al. [71] investigatedthe thermal expansion of self-reinforced PE and PP con-taining 2D (i.e., woven fabric) reinforcements. Good quality

Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310 1299

F lf-reinfs ence ofC

pPsfiwt[oi“dadtisvInisti6ACsb

sTfcsieiraor

hf

ig. 10. Effects of testing temperature on the stress-strain behavior of setress (�) and failure stress (�) as a function of temperature; (b) dependopyright (2000) with permission from Maney Publishing.

roducts were received under the following conditions:E: p = 0.75 MPa, T = 139 ◦C; PP: p = 3 MPa, T = 183 ◦C. Thehrinkage, E-modulus and linear thermal expansion coef-cient of the corresponding SRPMs were determined. Itas reported that especially the PP systems were sensi-

ive to changes in the compaction conditions. Hine et al.72] devoted a study to determine whether the insertionf film layers between the fabrics to be compacted resultsn improved consolidation quality, as well as whether thisinterleaving concept” can widen the temperature win-ow of the processing. Note that this method is basicallycombination of hot compaction and film stacking (to beiscussed later). This strategy yielded the expected results:he consolidation quality was improved (well reflectedn the mechanical property profile), the interlayer teartrength enhanced, and the processing temperature inter-al enlarged. This approach was also followed for PP fibers.n a further study, Hine et al. [73] incorporated carbonanofibers (CNF up to 20 wt.%) to improve the reinforc-

ng activity of the PP preform after hot compaction. Themall amounts of CNF significantly improved the proper-ies of isotropic PP. For example, adding CNF at a 5 vol.%ncreased the Young’s modulus at room temperature by0% and reduced the thermal expansion coefficient by 35%.ttempts were also made to improve the bonding betweenNF and PP via oxygen plasma treatment of the CNF (lessuccessful) and using a maleic anhydride grafted compati-ilizer for PP (more successful).

McKown and Cantwell [74] studied the strain-rate sen-itivity of a hot-compacted, self-reinforced PP composite.he SRPP specimens were subjected to strain rates rangingrom 10−4 to 10 s−1. The SRPP composite showed similarharacteristics to the neat PP material in respect to thetress–strain behavior with increasing strain rate. Stiffen-ng of the material in the elastic region was followed bynhanced yield stress and maximum stress with increas-ng strain rate. Parallel to that, the strain-to-failure waseduced. The failure mode of the SRPP composite was char-cterized by longitudinal fiber fracture with varying degree

f inter-ply delamination over the dynamic tensile loadingange studied.

Prosser et al. [75] investigated the thermoformability ofot compacted PP sheets with 2D reinforcement (woven

abric). It was reported that the self-reinforced PP sheets

orced PP with 2D reinforcement, schematically. (a) Dependence of yieldyield strain (�) and failure strain (�) as a function of temperature [75].

experienced considerable work hardening, according to in-plane tensile tests performed at high temperatures (cf.,Fig. 10). The authors observed that the optimum thermo-forming temperature is very close to that of the melting ofthe matrix formed by recrystallization of the molten partof the parent fiber/tape.

Romhány et al. [76] studied the fracture and failurebehavior of woven fabric-reinforced self-reinforced PP(Curv®), making use of mechanical fracture concepts andrecording the acoustic emission during the loading of thespecimens. The latter technique proved to be well suitedto characterize the consolidation quality. Jenkins et al. [77]prepared a range of flat hot-compacted single-polymercomposite panels from oriented PP and PE. The panelsdiffered in their dynamic modulus and damping capacityvalues. SRPMs were subjected to mechanical excitation,allowing their acoustic frequency responses over the audiobandwidth to be measured. The results showed the cor-relation of mechanical and acoustic frequency responsefunctions with the dynamic modulus, damping and spe-cific modulus of the panel materials. The ideal combinationof material properties to maximize the acoustic output ofthe panels was given by: high stiffness and low density toreduce the impedance of the panel and low damping toenhance the efficiency.

One major goal of the hot compaction technology was tooffer lightweight and easily recyclable thermoplastic com-posites to the transportation sector. As further applicationfields, sporting goods, safety helmets, covers and shells(also for luggage) were identified. Hot compacted PP sheetsfrom woven PP fabrics are marketed under the trade nameof Curv® (www.curvonline.com).

As mentioned before, the hot compaction method wassuccessfully transferred to many other polymers, like mul-tifilament assemblies of PET and PEN [53,78], PA-6.6 [57],POM and PPS [52], PEEK [52] and even PMMA [60]. Needlessto say, the optimum compaction conditions are stronglymaterial-dependent.

2.1.2.7. Production by film stacking. During film stacking,the reinforcing layers are sandwiched in-between thematrix-giving film layers before the whole “package” issubjected to hot pressing. Under heat and pressure, thematrix-giving material, which has a lower melting tem-

1300 Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310

posite p

exploiting the difference in the melting behavior of alpha-and beta-polymorphs. The alpha-PP tapes were arrangedin UD and cross-ply (CP) manners by winding, puttingbeta-nucleated PP films in-between the related reinforc-

Fig. 11. Scheme of the com

perature than the reinforcement, melts and infiltratesthe reinforcing structure accordingly. Recall that both thematrix and the reinforcement are given by the same poly-mer or polymer family. The film stacking procedure ishighlighted in Fig. 11.

The necessary difference in the melting temperaturesbetween the matrix and the reinforcement can be set byusing different polymer grades (e.g., copolymers for thematrix and homopolymers for the reinforcement, whichper definition belongs to the multi-component SRPMs)or polymorphs (e.g., lower melting modification for thematrix and higher melting one for the reinforcement; thisconcept yields a single-component SRPM). It is of greatimportance to have a large enough difference betweenthe melting temperatures of the composite constituents.Accordingly, the matrix-forming grade melts and wets outof the reinforcing structure without causing a temperature-induced degradation in the stiffness and strength of thereinforcement, or at least keeping it at an acceptable level.Those thermoplastic systems, which can be used to producesingle- and multi-component SRPMs via film stacking, aresummarized in Table 2.

In the follow-up section, we shall treat only the single-component SRPM versions. Bárány et al. [79–82] produceddifferent PP-based SRPMs. For reinforcement, highly ori-ented fibers in different textile architectures (carded mat,carded and needle-punched mat, in-laid fibers in knittedfabrics) were used, whereas for matrices either PP fibersof lower orientation (the same textile assemblies as indi-cated above) or beta-nucleated PP films were selected.Note that some of the above preforms do not even containinterleaving films and thus do not fall strictly under theheading of film stacking. The matrix-giving phase in themis either a discontinuous fiber or a knitted fabric. Never-theless, their consolidation occurs by hot pressing as in thecase of film stacking. One consideration is that the melt-ing temperature of the beta-modification of isotactic PPis >20 ◦C lower than the usual alpha-form [83]. The beta-modification can be achieved by incorporating a selectivebeta nucleating agent in the PP through melt compound-ing [84]. The concept of this alpha (reinforcement)/beta(matrix) combination should be credited to Karger-Kocsis

[85].

The consolidation quality of the all-PP composites pro-duced by Bárány et al. [86] was mostly studied as a functionof processing conditions, viz. temperature. With increas-ing temperatures, the stiffness and strength increased and

rocessing via film stacking.

the resistance to the out-of-plane-type perforation impactdecreases. The consolidation quality of the layered compos-ite laminates could be well qualified by the interlaminartear strength. Bárány et al. [80,81] later used PP fabric(woven type from split yarns) as the reinforcement andbeta-nucleated PP film as the matrix-giving material. Asmentioned above, the benefit of the beta-modification isthe widening of the melting temperature range betweenthe reinforcement and the matrix [87]. With increasingprocessing (pressing) temperature, the consolidation qual-ity was improved. Parallel to that, the density, the tensileand flexural stiffness and the strength increased, whereasthe penetration impact resistance diminished. The authorsproved by polarized light microscopy the presence of atranscrystalline layer between the PP reinforcement andPP matrix (cf., Fig. 12).

Izer and Bárány [82] manufactured all-PP compositesby direct hot pressing of suitable textile assemblies. Asindicated above, these assemblies contained both the rein-forcement and matrix-giving phases in form of fibers withdifferent orientations (draw ratios). Recall that the latter isthe guarantee for a small difference in the melting temper-atures, which was used in this case. Abraham et al. [88]produced all-PP composites with tape reinforcement by

Fig. 12. Transcrystalline layer of PP fiber and �-rPP matrix.

Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310 1301

Table 2Possible polymer pairs to produce SRPMs.

Composite Matrix Reinforcement Processing temperature range (�T)

• PE LDPE UHMWPE fiber 20–40 ◦CHDPE UHMWPE fiber 20–40 ◦C

• PP �-PP* highly oriented iPP fiber 20 ◦Crandom PP copolymer highly oriented iPP fiber 25 ◦CiPP* highly oriented iPP fiber 8–10 ◦C

• Polyester PETG PET fiber 40–60 ◦CPETG PEN fiber 15–20 ◦C

× PBT ◦

LCP (Vec

itm

iylahbmWtma3mdSi

2

mTomSsiN

22samtC(sF

va

CBT• LCP LCP

* Single component SRPM.× Production occurs via liquid composite molding.

ng tape layers. The stiffness as a function of temperature ofhe corresponding composites was determined by dynamic

echanical thermal analysis (DMTA).Bhattacharyya et al. [89] prepared an SRPM by combin-

ng hot compaction and film stacking. High tenacity PA-6arn was used as a reinforcement, and PA-6 film from pel-ets was used as the matrix. The yarn was subjected tonnealing in a vacuum (3 h at 150 ◦C) in order to get aigher melting point. Two yarn layers were sandwiched in-etween two matrix films and subjected to compressionolding at 200 ◦C for 5 min under a pressure of 15 MPa.ith the combination of these two techniques, good wet-

ing properties were achieved and materials with excellentechanical properties were produced. The tensile modulus

nd strength of the composites were improved by 200 and00–400%, respectively, compared to the initial isotropicatrix film. An overview on the production methods, con-

itions and product characteristics of single-componentRPMs produced in multi-step (ex situ) processing is givenn Table 3 – cf. Fig. 1.

.2. Multi-component SRPMs

SRPMs can also be produced by the combination of poly-ers that belong to the same family of polymeric materials.

he major goal during their preparation is the achievementf good adhesion (bonding) between the reinforcing andatrix-giving polymer phases. Like the single-component

RPMs, the reinforcing structure may be generated iningle- (in situ) or multi-step (ex situ) operations. Accord-ngly, a similar grouping as before can also be followed here.ext, the different variants will be briefly introduced.

.2.1. Single-step production

.2.1.1. Multi-component extrusion yielding self-reinforcedtructures. The extrusion die with a convergent sectionllowed us to set a unidirectional (1D) molecular align-ent in situ, which will work as the reinforcement owing to

he supermolecular structure formed by the crystallization.hen et al. [90] solved the problem of biaxial orientation2D), however, by using a specially designed fish-tail-

haped (bi-cuneal shape) extrusion die, as depicted inig. 13.

Composites with planar reinforcement were producedia this die from HDPE and HDPE/UHMWPE blends using

single-screw extruder. The mold was oil-tempered

60–80 Ctran® M) 25 ◦C

(T = 126–137◦C), and the optimum processing pressurewas between 15 and 30 MPa. Under conventional extru-sion conditions, the tensile strength of the extruded sheetwas comparable to conventionally molded HDPE sam-ples. The tensile strength was almost the same in boththe machine (MD) and the transverse directions (TD). Thetensile strengths of the HDPE/UHMWPE in the extrusionand transverse directions were six and three times higher,respectively, than those of the related traditionally pro-duced sheet (HDPE).

2.2.1.2. Multi-component SCORIM/OPIM. Zhang et al.[91,92] processed LDPE/HDPE and HDPE/PP blends by theearlier introduced OPIM technique (oscillation frequency:0.3 Hz). It was established that with increasing LDPEcontent the tensile strength diminishes, whereas thetoughness increases for the LDPE/HDPE blends. Morpho-logical studies confirmed the onset of a shish-kebab-typesupermolecular structure. The tensile strength of theHDPE/PP blends could also be markedly increased (five-fold) when the PP content remained below 10 wt.%. Zhanget al. [93] investigated the performance of HDPE/UHMWPEwhen processed by the SCORIM technique. Tribologicaltests showed that the wear resistance of the related systemwas ca. 50% better than that of traditionally molded speci-mens. Table 4 displays the production methods, conditionsand product characteristics of multi-component SRPMsproduced in single-step (in situ) processing - cf. Fig. 1.

2.2.2. Multi-step productionThe first publication of this processing version should

be credited to Capiati and Porter [94]. They combinedHDPEs with different melting characteristics. The high-modulus fibers (reinforcement) melted at 140, while thematrix-giving HDPE melted at 131 ◦C. The HDPE fiber wasembedded in the molten HDPE using a special rheome-ter. After cooling/solidification, the fiber in this single-fiberreinforced composite was subjected to a pull-out test.It was reported that the interfacial shear strength wascomparable with that of the glass fiber/epoxy system.Moreover, the presence of a transcrystalline layer was

detected at the fiber/matrix surface.

2.2.2.1. Consolidation of coextruded tapes. The develop-ment of SRPMs is best reflected by searching for optionsthat amplify the difference between the melting of the

1302Á

.Km

ettyet

al./Progressin

Polymer

Science35

(2010)1288–1310

Table 3Production methods, conditions and product characteristics of single-component SRPMs produced in multi-step (ex situ).

Num. Process Materials Processing conditions Results Comment Ref.

1 Ram extrusionHDPE Entrance semi-angel of 30◦ ,

p = 40–110 MPa, T = 90–120 ◦CEt = 0.2–0.9 GPa high friction,

low extrusion rates[24]

HDPE p = 0–100 MPa, � = 3.75–6, T = 105 ◦C Et = 1.1 GPa, k = 0.4–0.5 W/mK thermal relaxation after leavingthe die

[26]

PP p = 0–100 MPa, � = 3,75–6, T = 130 ◦C Et = 1.7 GPa, k = 0.1 W/mK lubrication applied [26]

2 Hydrostatic extrusion UHMPE v = 1–50 cm/min, p = 0–250 MPa Et = 10–60 GPa batch process [24]

3 Die-drawingHMWPE � = 1–13, T = 115 ◦C � = 1–7 elongated spherulites

� = 7–12 lamella structuredecreased creep [27]

PP � = 6–11, T = 200 ◦C Et = 1–12 GPa increased tenacity [28,29]POM Entrance semi-angel of 5–20◦ , � = 1–16 EB = 3–26 GPa increased crystallinity [30]

4 Super drawing PTFE � = 6–160 Et = 38–102 GPa, �B = 0.7–1.4 GPa increased crystallinity [31]

5 RollingHDPE � = 2–9 Et = 3–15 GPa, �B = 30–290 MPa increased crystallinity [24,34–36]PP � = 2–9 Et = 4–10 GPa, �B = 30–350 MPa increased crystallinity [24,37]

6 Gel drawing UHMWPE � = 3–22 Et = 10–90 GPa, �B = 0.3–3.5 GPa increased crystallinity [40–42]

7 Orientation drawingHDPE TD = 110–120 ◦C, � = 5.8–30 Et = 6–35 GPa, �B = 250–1200 MPa micro-cracking [43,44]iPP TD = 163–164 ◦C, � = 21.5–35.5 Et = 18.5–24.7 GPa, �B = 0,6–1.1 GPa physical aging [45]

8 Hot compaction

UHMPE* Tproc = 134–154 ◦C, p = 0.7–21 MPa Et = 9.9–85 GPa, �B = 200 MPa unidirectional, woven structure [52,55,56,61,65,68]UHMWPE Tproc = 145 ◦C, p = 31–49 MPa Et = 0.4 GPa, �B = 3.7 MPa crosslinked [62,64]PP* Tproc = 164–195 ◦C, p = 1.1–14 MPa Et = 1.6–3.9 GPa, �B = 25–168 MPa high acoustic output [59,66,69–72,74–77]PET Tproc = 249–256 ◦C, p = 1.9–32.4 MPa Et = 11.5–13 GPa, �B,T = 15–35 MPa good adhesion [53]PEN Tproc = 268–276 ◦C, Et = 2–9.6 GPa, �B = 22–207 MPa 0/90 multifilament [52,78]PA-6.6 p = 2.8 MPa Et = 4.1 GPa, �B = 50 MPa woven structure [57]PPS Tproc,opt = 288 ◦C, Et = 5.2 GPa, �B = 80 MPa chemical resistance [52]PP/CNF Tproc = 230 ◦C, Tmold = 190 ◦C Et = 1.5–2.7 GPa 2, 4, 6, 10, 20 wt% of CNF [73]POM Tproc,opt = 182 ◦C Et = 1.9–5.3 GPa, POM powder [58]PMMA Tproc = 100–125 ◦C �B = 65–165 MPa unidirectional structure [60]PEEK Tproc,opt = 347 ◦C Et = 3.65 GPa, �B = 100 MPa woven structure [52]

9 Film-stacking

�-PP/�-iPP Tproc = 150–170 ◦C, p = 7 MPa Et,L = 2.4–2.7 GPa, �B,L = 20–100 MPaEt,T = 1.6–2.4 GPa, �B,T = 20–43 MPa

carded needle-punched mat [79]

�-PP/PP Tproc = 156–186 ◦C, p = 7 MPa Et = 2.5–3 GPa, �Y = 90–100 MPa PP woven textile [80,81]iPP/iPP Tproc = 160–170 ◦C, p = 6 MPa Et = 2.1–2.5 GPa, �B = 29–140 MPa knitted, carded and needle

punched mats[82]

�-PP/�-PP Tproc = 160 ◦C, p = 7 MPa Et = 2.3 GPa, �B = 60 MPa winding [88]

* Summarized results.

Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310 1303

F ie: (1) tes yright (

rhcd2tc

FS

ig. 13. Schematic representation of the self-reinforcing sheet extrusion dection, and (4) double functional temperature–pressure sensor [90]. Cop

einforcement and the matrix. Recall that this range wasighly limited for hot compaction. Peijs [95] developed aoextrusion technique for which the melting temperature

ifference between the composite constituents reached0–30 ◦C. The invention was to “coat” a PP homopolymerape from both sides by a copolymer through a continuousoextrusion process. Note that a copolymer melts always

ig. 14. Coextrusion technology with additional stretching to produce high-streons.

Fig. 15. Production of composites with UD ta

mperature controlling oil bath, (2) the straight section, (3), the convergent2007) with permission from Elsevier.

at lower temperatures than the corresponding homopoly-mer, owing to its less regular molecular structure. Thecoextruded tape was stretched additionally in two-steps

(cf., Fig. 14). This resulted in high-modulus, high-strengthtapes.

The primary tapes could be assembled in different ways:as in composite laminates (ply-by-ply structures with dif-

ngth tapes [96]. Copyright (2007) with permission from John Wiley and

pe alignment from coextruded tapes.

1304 Á. Kmetty et al. / Progress in Polymer S

Tab

le4

Prod

uct

ion

met

hod

s,co

nd

itio

ns

and

pro

du

ctch

arac

teri

stic

sof

mu

lti-

com

pon

ent

SRPM

sp

rod

uce

din

sin

gle-

step

(in

situ

).

Nu

m.

Proc

essi

ng

Mat

eria

lsPr

oces

sin

gco

nd

itio

ns

Res

ult

sC

omm

ent

Ref

.

1Se

lf-r

ein

forc

edex

tru

sion

HD

PE/U

HM

WPE

T pro

c=

125–

180

◦ C,

p=

2–30

MPa

�B

=20

–170

MPa

spec

iall

yd

esig

ned

fish

-tai

lsh

aped

extr

usi

ond

ie[9

0]

2SC

OR

IM/O

PIM

HD

PE/L

DPE

T mel

t=

200

◦ Cp p

roc

=90

MPa

T mol

d=

25◦ C

�B

=21

–109

MPa

shis

h-k

ebab

stru

ctu

re(W

AX

D)

[91]

HD

PE/P

PT m

elt=

220

◦ Cp p

roc

=32

–48

MPa

T mol

d=

42◦ C

�B

=25

–90

MPa

(<10

wt%

PP)

[92]

HD

PE/U

HM

WPE

–�

=0.

3–0.

5(0

.3)

Load

-car

ryin

gca

pac

ity

=70

N(3

5N

)

mic

ro-c

rack

sp

aral

lelt

oth

esu

rfac

e(c

onve

nti

onal

pac

kin

g)[9

3]

cience 35 (2010) 1288–1310

ferent tape orientations, such as UD (cf., Fig. 15) and CP)or integrated in various textile structures (e.g., woven fab-rics). The consolidation of the related assemblies occurredby hot pressing.

The advantage of this method is that the reinforce-ment (core) content of the tape may be as high as ca.80%. This, along with the high draw ratio, yielded tapes ofexcellent mechanical properties (E-modulus > 6 GPa, ten-sile strength > 200 MPa). Cabrera et al. [97] prepared all-PPcomposites from UD and woven fabric assemblies of coex-truded tapes. For the consolidation of the UD composites, a17 MPa pressure was used and the temperature covered therange between 140 and 170 ◦C. The time was kept constant(15 min) during hot pressing. The E-modulus of the lami-nates, measured both in the tape direction and transverseto it, was not much affected by the processing tem-perature. In contrast, the interlaminar tear strength wasimproved by increasing the temperature, well reflectingthe improvement in the consolidation quality. The wovenfabric-reinforced composites were subjected to falling dart(perforation impact) tests. Based on the related specific (i.e.,thickness-related) perforation impact energy data, the all-PP composites outperformed both the glass fiber (GF) mat-(three times higher) and flax mat-reinforced counterparts(six times higher). Alcock et al. [98] manufactured UD com-posite sheets by winding the coextruded tapes on a metallicframe that was later put in-between the plates of a pressoperated in the temperature interval of T = 140–160 ◦C. Theproperties of the composites were determined in mechan-ical investigations, whereas the reinforcement content(reaching 90 wt.%) was determined via microscopic inves-tigations. As usual, for all UD-reinforced composites, boththe tensile E-modulus and strength decreased with increas-ing angle between the reinforcing and loading directions(off-set) during their testing. The transverse compres-sive strength (10 MPa) was not affected by the pressingtemperature. The results received were compared withthose measured on 50 wt.% UD GF-reinforced PP com-posites. Although the UD-GF PP composite preformedbetter than the all-PP material, the latter took the leadwith respect to the related specific (i.e., density-related)properties. In follow-up studies, Alcock et al. [96,99–102]investigated the structure–property relationships in all-PP composites produced from woven fabrics composedof coextruded tapes. When the consolidation took placeat low temperatures (T = 125 ◦C) and under low pressures(p = 0.1 MPa), the sheets exhibited excellent resistance tothe perforation impact. This was traced to an intensivedelamination between the fabric layers that was triggeredduring this high-speed perforation process. Up to a 2 mmsheet thickness, the perforation energy increased linearlywith the sheet thickness. Ballistic test results confirmedthat the performance of composite sheets from Pure®

tape is comparable with that of the state-of-art ballisticmaterials. The authors draw attention to the fact that themechanical performance of the all-PP composites, which

contain fabrics of coextruded tapes, can be optimizedupon request by selecting suitable textile architecturesand hot pressing/consolidation parameters (pressure, tem-perature). Barkoula et al. [103] investigated the fatigueperformance of PP tapes and woven tape fabric-reinforced

olymer S

a(pdpcic

fwwidflttistaoiofpw[fiiRtc

2SmPitrIpttmhStaiToithaidt

Á. Kmetty et al. / Progress in P

ll-PP composites. They found that the endurance limitor fatigue threshold, below which no fatigue-inducedroperty reduction occurs), controlled by the onset ofelimitation, is strongly affected by the processing tem-erature. The fatigue threshold of the optimum processedomposite was at 65% of the static tensile strength. Thiss markedly higher than that of GF mat-reinforced PP-omposites, which show a range of 30–40% [104].

Banik et al. [105,106] studied the short-term creep per-ormance of coextruded tape-reinforced PP compositesith both UD- and CP-type tape lay-ups. The related sheetsere produced by vacuum bagging in an autoclave (which

s almost exclusively used for thermoset composite pro-uction) under a 2.4-MPa pressure and at T = 138 ◦C. Theexural creep tests were performed in a DMTA device inhe temperature range of 20–80 ◦C. It was reported thathe creep depends on the composite lay-up. By adopt-ng the temperature-time superposition principle to thehort-term creep results, a master curve was constructedhat predicted the long-term creep at a given temper-ture. Kim et al. [107] also studied the creep responsef all-PP composites and emphasized that small changesn the processing conditions have a pronounced effectn the creep behavior. It is noteworthy that compositesrom coextruded PP tapes in different assemblies wereroduced by various techniques, such as hot pressing, tape-inding [108], stamp forming and vacuum bag/autoclaving

109,110]. Moreover, the related sheets were used for theace-covering of different sandwich structures with coresncluding honeycomb structures and foams. The face sheet-ng occurred with or without additional primer [111].ecall that the coextruded PP tapes are known underhe trade names of Pure® and Armordon® (www.pure-omposites.com; www.armordon.com).

.2.2.2. Film stacking. This technique is usually used forRPMs in which the constituents are from the same poly-er family. Shalom et al. [112] produced high-strength

E fiber- (Spectra®) reinforced HDPE composites by wind-ng the fiber in a unidirectional manner and sandwichinghe HDPE films in-between the wound fiber layers. Theeinforcing fiber content in the UD assembly was 80 wt.%.ts consolidation occurred by hot pressing (T = 137 ◦C,= 16.5 MPa). Coupon samples were subjected to tensile

ests with variation of the loading direction in respect tohe UD fiber alignment (off-axis tests). As expected, the E-

odulus, yield strength and resistance to fracture were alligher when the off-axis angle was smaller. Houshyar andhanks [113] used a mat from PP homopolymer fibers ashe reinforcement (fixed at 50 wt.%) and PP copolymer films the matrix-giving material. The difference in their melt-ng temperatures was ca. 16 ◦C according to DSC results.he fiber diameter in the mat varied. The hot consolidationccurred between 155 and 160 ◦C. It was found that with anncreasing diameter of the mat fibers, both the stiffness andhe strength of the composites increased. The surface of the

omopolymer PP fiber acted as a heterogeneous nucleatornd initiated transcrystalline growth. In follow-up stud-es [114,115], it was demonstrated that with increasingiameter of the reinforcing PP fiber, the void content inhe composite can be reduced. The maximum strength was

cience 35 (2010) 1288–1310 1305

reached when the diameter of the fiber was ca. 50 �m.The creep results of the related composites, which werealso modeled by the Burgers model, demonstrated that anincreasing reinforcing fiber content was accompanied byan increasing resistance to creep. The objective of furtherstudies of the Shanks’ group [116,117] was to deduce possi-ble effects of different textile architectures (covering bothnon-woven and woven ones) on the mechanical propertiesof the related all-PP composites. During the consolidation,they were subjected to a low pressure (ca. 0.01 MPa) atT = 158 ◦C for 15 min. The mechanical results showed thatthe properties of the woven composites strongly dependon the woven geometry. The composite with the satin clothdelivered the best properties. This was due to the advan-tages of the satin parameters, such as the long float length,high fiber count, few interlace points and loose pattern.

It is noteworthy that the authors used the term “com-paction” even though this is reserved for those techniquesin which a part of the reinforcing phase becomes moltenand thus overtakes the role of the matrix after cooling.This is not the case in film stacking, where the melt-ing temperature of the reinforcing fiber or tape is usuallynot surpassed. In order to improve the energy absorp-tion capability of the resulting composites, Houshyar andShanks [118] modified the matrix. This was done by extru-sion melt compounding of the matrix-giving PP copolymerwith ethylene-propylene rubber (EPR; up to 30 wt.%) withfollow-up sheeting.

Houshyar et al. [119] modeled the PP fiber-matrixcomposites with finite element analysis. This modeldemonstrated that the stress concentration at the fiber-matrix interface increased with decreasing fiber content.The ratio between matrix and fiber stiffness was signifi-cant, and the interfacial stress carried by both constituentsacted to reduce the risk of premature interfacial failure andincrease the mechanical properties of the composite. Thefinite element model showed that at a low fiber content, thefiber was not able to share a larger portion of the appliedstress. The matrix carried the main portion of stress andyielded on a large scale when the applied stress reachedthe matrix strength.

Bárány et al. [120] prepared composites using ran-dom PP copolymer films and carded and needled punchedmats as matrix and reinforcing phases, respectively. Thenominal reinforcement content was 50 wt.%. The consol-idation was performed at different temperatures in therange of T = 150–165 ◦C. Consolidation at 150 ◦C resultedin poor performance, whereas that above T = 165 ◦C didnot yield additional property improvement. Bárány et al.[80,81] also studied the perforation impact resistance ofall-PP composites containing woven fabrics from split flatyarns as reinforcement and films composed of both alphaand beta-phase random PP copolymers as matrix-givingmaterials. The beta-modification was produced by usinga selective beta-nucleant. The perforation impact resis-tance of the composite with beta-nucleated random PP

copolymer was higher that the alpha variant. Izer andBárány [121] estimated the long-term flexural creep ofself-reinforced polypropylene composites based on short-term creep tests performed at different temperatures. AnArrhenius-type relationship was used to shift the related

1306 Á. Kmetty et al. / Progress in Polymer Science 35 (2010) 1288–1310

Fig. 16. Lamellar structure within the transcrystalline layer as a functionof the consolidation temperature [122]. Copyright (2003) with permissionfrom John Wiley and Sons.

creep data along the time axis. It was found that withimproved consolidation (increasing processing tempera-ture) the creep compliance decreased. Moreover, goodcorrelations were found between the creep compliance anddensity and between the creep compliance and interlayerpeel strength.

Abraham et al. [88] produced high-strength alpha PPhomopolymer tapes by a single-step hot stretching tech-nique and used this as UD or CP reinforcement in alpha- andbeta-phase random PP copolymer matrices. The interphasebetween the reinforcement and matrix was composed of atranscrystalline layer that was larger in the beta- than in thealpha-phase random PP copolymer matrix. This finding wastraced to the fact that the composite with beta-nucleatedmatrix performed better than the alpha version.

Kitayama et al. [122] produced SRPMs from PPhomopolymer fibers (reinforcement) and a random PPcopolymer (matrix) by film stacking and studied the inter-phase formed. Here, a transcrystalline layer was resolved,the structure of which changed with the consolidation tem-perature (cf., Fig. 16).

The lamellar structure, depicted in Fig. 16, can bestretched upon loading without detaching from the surfaceof the reinforcing PP homopolymer fiber, which is very ben-eficial in composites. Recall that the lamellar structuring inthe transcrystalline layer for optimum stress transfer fromthe matrix toward the fiber, proposed by Karger-Kocsis etal. [123], is very similar to that in Fig. 16 (cf., Fig. 17).

The recycling via melt processing of one- and two-component all-PP composites has already been topic ofinvestigations [124].

Ruan et al. [125] manufactured nanoparticle-filled self-reinforced PP composites, for which they used fumed SiO2as the nanoparticle. The nanoparticles were preheated at140 ◦C under vacuum for 5 h. Then, the mixture of monomer(butyl acrylate) and the nanoparticles and a certain amount

of solvent was irradiated by 60Co �-ray in air at a dose rateof 80 kGy. The resultant, poly(butyl acrylate) (PBA) graftednano-SiO2 (SiO2-g-PBA) with a percent grafting of 3.35%,was used for the subsequent composite manufacturing.Untreated or treated nano-silica was melt compounded

Fig. 17. Hypothesized interphase with lamellar interlocking and amor-phous phase as adherent for the transcrystaline layer initiated by flat-ontype lamellar growth on fiber surface [123]. Copyright (2000) with per-mission from Elsevier.

with iPP at 180 ◦C in the mixer. The content of nano-silica in all of the composites is 1.36 vol.%. The sheets ofSiO2/PP were produced by hot pressing, and then the sheetswere hot drawn under a temperature slightly lower thanthe melt temperature of PP (150 ◦C) at a constant veloc-ity. A film with a thickness of 50 �m was blown fromthe random PP copolymer by film blowing. Finally, thestretched sheets were film-stacked with copolymer filmsby a specially designed mold and were hot pressed at differ-ent processing temperatures (T = 150–175 ◦C) and holdingpressures (2.0–5.0 MPa) for a constant holding time of10 min. According to the mechanical properties reported,the incorporation of nanoparticles into the polymer matriximproved the mechanical properties of the self-reinforcedcomposites.

Pegoretti et al. [126–128] used thermoplastic liquidcrystalline polyester fibers for both the reinforcement(Vectran® HS, Tmelt = 330 ◦C) and the matrix (Vectran® M,Tmelt = 276 ◦C). Unidirectional composites were preparedin a two-stage process. At first, both Vectran® M and HSas-received fibers were wound on an open metal frame,and after winding the LCP was consolidated in a hot press.An optimum processing temperature of T = 275 ◦C wasdeduced and was associated with the lowest void contentand highest mechanical strength. Table 5 lists the produc-tion methods, conditions and product characteristics ofmulti-component SRPMs produced in multi-step (ex situ)processing - cf. Fig. 1.

2.2.2.3. Microfibrillar reinforced composites (MFCs). MFCsare polymer-polymer composites containing constituentsthat are incompatible with each other and possess differentmelting temperatures. In MFCs, the reinforcing microfib-rils are given by “flexible” macromolecules that have beenaligned during the production. Essential stages of MFCpreparation are as follows: blending, extrusion, drawing

and annealing. The latter occurs at a constant strain abovethe Tm of the component that melts at a lower tempera-ture. MFCs are usually made of condensation (PET, PA-6.6)polymers working as reinforcements and polyolefins (PP,PE) acting as matrices. MFCs from PP/PE, PET/PE, PET/PP,

Á. Kmetty et al. / Progress in Polymer STa

ble

5Pr

odu

ctio

nm

eth

ods,

con

dit

ion

san

dp

rod

uct

char

acte

rist

ics

ofm

ult

i-co

mp

onen

tSR

PMs

pro

du

ced

inm

ult

i-st

ep(e

xsi

tu).

Nu

m.

Proc

essi

ng

Mat

eria

lsPr

oces

sin

gco

nd

itio

ns

Res

ult

sC

omm

ent

Ref

.

1C

oext

rusi

on

rPP/

PPT p

roc

=14

0,16

0◦ C

,p=

2.4

MPa

,�=

17�

B,L

=38

5M

Pa,E

t,L

=13

GPa

�B

,T=

5M

Pa,E

t,T

=1.

5G

Pau

nid

irec

tion

alst

ruct

ure

[98]

rPP/

PPT p

roc

=14

0–17

0◦ C

,p=

17M

Pa,�

=16

�F

=20

5M

Pa,E

B=

6G

Paw

oven

stru

ctu

re[9

7]rP

P/PP

T pro

c=

140

◦ C,p

=1–

15M

Pa,�

=17

�B

=20

–230

MPa

,Et=

2–7

GPa

wov

enst

ruct

ure

,in

crea

sin

gim

pac

t-en

ergy

abso

rpti

on[9

9–10

2]

rPP/

PPT p

roc

=14

5◦ C

,p=

2.4

MPa

�F

=90

–160

MPa

,EB

=3–

12G

Paco

nso

lid

atio

nin

vacu

um

bag

(UD

,0/9

0st

ruct

ure

)[1

05,1

06]

2Fi

lm-s

tack

ing

HD

PE/P

ET p

roc

=13

7◦ C

,p=

16.5

MPa

�Y

=34

–190

MPa

,Et=

2–7

GPa

win

din

gan

gels

(30–

45◦ )

[112

]rP

Pco

pol

ymer

/PP

T pro

c=

155–

160

◦ C,p

=16

.5M

Pa�

Y=

15–3

0M

Pa,

E t=

1.6–

1.8

GPa

tran

scry

stal

lin

est

ruct

ure

(lig

ht

mic

rosc

opy)

[113

]

rPP

cop

olym

er/P

PT p

roc

=15

8◦ C

,p=

11–1

4kP

a,�

B=

17–2

6M

Pa,E

t=

2–3

GPa

dif

fere

nt

wov

enst

ruct

ure

[116

]rP

Pco

pol

ymer

-EP/

PPT p

roc

=15

2◦ C

�B

=17

–26

MPa

,E t

=0.

3–0.

7G

Pa0–

30w

t%EP

[118

]

rPP/

PPT p

roc

=14

7–17

7◦ C

,p=

7M

Pa�

Y=

110–

120

MPa

,E t

=2–

2.5

GPa

hig

hly

stre

tch

edsp

lit

PPta

pes

[80,

81]

�-r

PP/P

PT p

roc

=13

6–16

6◦ C

,p=

7M

Pa�

Y=

100–

110

MPa

,E t

=1.

5–2

GPa

hig

hly

stre

tch

edsp

lit

PPta

pes

[80,

81]

rPP/

hPP

T pro

c=

150–

163

◦ C,p

=1.

9M

Pa�

B=

14–4

2M

Pa,

E t=

0.4–

0.5

GPa

tran

scry

stal

lin

est

ruct

ure

(TEM

)[1

22]

LCP/

LCP

T pro

c=

260–

285

◦ C,p

=4.

4M

Pa�

F=

24–3

6M

Pa,E

B=

36–4

2G

Pau

nid

irec

tion

alst

ruct

ure

[126

,127

]

cience 35 (2010) 1288–1310 1307

PA-6.6/PP and PA-6.6/PE blends (note that in the above listsome reinforcement/matrix combinations are highlighted)exhibited pronounced improvements in the mechanicalproperties compared to those of the respective isotropicmatrix. The mechanical properties of MFC are similar tothose of short glass fiber-reinforced composites containingthe same matrix [129].

3. Outlook and future works

Self-reinforced polymeric materials (SRPMs) will con-tinue to remain of interest. This prediction is based onthe fact that SRPMs are lightweight (their densities arelower than most of the traditional composites) and envi-ronmentally benign (especially due to their easy recyclingvia reprocessing in the melt).

The R&D works with single-component SRPMs willprobably focus on multi-step production methods, whichallow greater freedom in shaping and design. For thatpurpose, angular pressing, which is well established formetals, will be explored next [130]. However, the dreamof researchers is still the single-step production of single-component SRPMs via injection molding.

For multi-component SRPMs, one can expect the fastdevelopment of one-step extrusion blow molding opera-tions. This will be fuelled by the need for hollow containerswith improved barrier properties. Among the multi-stepproduction methods of multi-component SRPMs, thosewith loose textile assemblies that can be consolidated andshaped simultaneously are in a favored position. Thereis ongoing research to develop injection-moldable multi-component SRPMs.

With respect to the matrix/reinforcement combi-nations, amorphous/semicrystalline and semicrystalline/semicrystalline (by exploiting the melting temperature dif-ference between polymorphs) combinations will be furtherinvestigated. The reinforcing phase will be often modi-fied by nanofillers, especially by those having high aspectratios (carbon nanotubes, carbon nanofibers, graphene lay-ers, layered silicates), to increase its stiffness, strength andthermal stability.

The feasibilities of some combinations, listed in Fig. 1,have yet to be checked. For example, it is of great impor-tance whether such a polymer part can be produced forwhich a 3D self-reinforcing structure (for example, by ori-ented crystallization) is simultaneously generated. In thisrespect, novel processing techniques allowing the achieve-ment of very high heating and cooling rates (in hundredsof ◦C/min range) may be of great help. It can be prophesizedthat in the near future, SRPMs with an amorphous poly-mer matrix and a semicrystalline reinforcing phase (usingpolymers that belong to the same polymer family) willbe pushed forward. Differences in the tacticity resultingin semicrystalline and amorphous versions (e.g., isotac-tic or syndiotactic PMMA reinforcements in amorphousPMMA matrices) as well as the phenomenon of polymor-

phism in semi-crystalline polymers will be favored topicsof related research and development activities. New SRPMsand SRPCs will be produced by combining new methods(e.g., electrospinning of fibers) with well established ones(e.g., film stacking).

olymer S

1308 Á. Kmetty et al. / Progress in P

Acknowledgements

The authors want to thank the Hungarian ScientificResearch Fund (OTKA K75117). T. Bárány is thankful forthe János Bolyai Research Scholarship of the HungarianAcademy of Sciences.

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[3] Prox M, Pornnimit B, Varga J, Ehrenstein GW. Thermoanalyti-cal investigations of self-reinforced polyethylene. J Therm Anal1990;36:1675–84.

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