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polypropylene composites
Composites Science and Technology 60 (2000) 2159±2167
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Experimental detection of a transcrystalline interphase in glass-®bre/polypropylene composites G. Pompe *, E. MaÈder Institute of Polymer Research Dresden, Hohe Str. 6, D-01069 Dresden, Germany Received 12 January 2000; received in revised form 28 April 2000; accepted 11 May 2000
Abstract Crystallization and melting behaviour of polypropylene in glass-®bre-reinforced composites have been investigated by means of dierential scanning calorimetry. Composites with dierently size-coated glass ®bres show dierent behaviour both in dynamic and isothermal crystallization. A high crystallization rate and a short induction time were observed for composites with unsized ®bres at isothermal crystallization. Glass ®bres sized with aminosilane and polypropylene ®lm-former revealed a quantitatively separable twostep crystallization process. With the assignment to the transcrystallization and spherulitic crystallization the thickness of the transcrystalline PP interphase in this composite could be estimated to about 1 mm. The in¯uence of transcrystallinity on the mechanical properties was discussed. # 2000 Elsevier Science Ltd. All rights reserved. Keywords: A. Glass ®bre; A. Polymer-matrix composites; B. Interphase; B. Mechanical properties; B. Thermal properties
1. Introduction The interphase in composites is the part of the matrix surrounding the ®bre surface and possessing local properties dierent from those of the bulk matrix. The size and type of interphase varies strongly and depends on the nature of the ®bre and its surface as well as on the polymer matrix [1,2]. Because stress transfer in the composites should be in¯uenced by the interphase, optimization of the mechanical properties of composites requires an extensive knowledge of the behaviour of the interphases and its eect on mechanical performance. The interphase results from perturbations induced at ®bre surfaces and propagated away from the ®bre surface into the matrix. Under appropiate conditions [3±7] a highly oriented, so-called transcrystalline interphase can appear in composites with semicrystalline polymer (e.g. PP) as the matrix. The structure and size of transcrystalline interphases are both system-and condition-speci®c. Some in¯uences are known, for instance the surface energy of ®bre, changed by the ®bre type, the sizing, and coupling agent [5±7], and the topography of the ®bre (level of * Corresponding author. Tel.: +49-351-4658-255; fax: +49-3514658-284. E-mail address: [email protected] (G. Pompe).
roughness) [6]. It was also observed that transcrystallization can only occur at crystallization temperatures lower than the critical temperature Ttc cr [5±7]. Our investigations were focused on the question of which surface treatment of the glass ®bres changes the crystallization behaviour in conjunction with a changed crystalline morphology of the matrix nearby the ®bre surface compared with the bulk matrix behaviour. A ®rst attempt to use dierential scanning calorimetry (DSC) for the detection of transcrystallinity was made by He and Porter [8]. In this paper, a quantitative analysis of systematic DSC measurements will be reported for PP composites with dierently sized glass ®bres. At constant ®bre volume fraction (0.5) the speci®c glass ®bre surface is varied by using dierent glass ®bre diameters. 1.1. Transcrystallinity Wu et al. [9] and also Billon et al. [10,11] de®ned transcrystallization as a nucleation-controlled process taking place under quiescent conditions in a semicrystalline polymer in contact with other materials, e.g. ®bres. At high heterogeneous nucleation ability of the surface the closely spaced nuclei hinder the lateral extension and force growth in one direction, namely perpendicular to the
0266-3538/00/$ - see front matter # 2000 Elsevier Science Ltd. All rights reserved. PII: S0266-3538(00)00120-2
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®bre surfaces. Once nucleation has occured, the growth of spherulites in the matrix and of the transcrystalline layer on the surface takes place at the same rate [5± 7,12]. A dierent nucleation rate of both crystallization mechanism is the necessary condition for the formation of a detectable extent of transcrystallinity. This dierence can amount to about three orders of magnitude found, for PTFE-®bre/PP composites [12]. Fig. 1 shows the scheme of the interphase with transcrystalline morphology in a ®bre/semicrystalline-polymer composite [13]. Three regions of dierent thickness, i,k might be distinguished. Region 1 of thickness i,1 includes the interface between the ®bre and the transcrystalline part of PP matrix (termed the `®rst' interface). Both the thickness i,1 and the speci®c surface of the ®rst interface will be determined by the roughness of the ®bre/sizing surface. Region 2 is the homogeneous part of the transcrystalline layer without boundaries (thickness i,2). Growth of the transcrystalline layer will be halted by the impingement on the spherulites nucleated and grown in the bulk matrix. As the induction period for surface nucleation is shorter [6,8,12], and the growth rate of both crystalline structures is the same [5±7,12],the thickness i,2 will be determined above all by the ability of nucleation to occur on the surface in comparison to that in the bulk. This means that i,2 can be in¯uenced by changes in the nucleation rate on the ®bre surface or sizing with a high functionality for the nucleation process as well as by changes in the nucleation rate in the bulk matrix, e.g. nucleating agents, molecular weight. The conditions of crystallization also play an important role. It was
observed that transcrystallization under quiescient conditions can only take place at crystallization temperatures, Tc, lower than a critical temperature Ttc cr [5,7]. A brief application of shear stresses along the ®bre/matrix interface permits the formation of morphology similar to that of the transcrystalline, also at Tc>Ttc cr [7]. These observations result in the assumption that stress applied externally or formed internally is a condition for the formation of transcrystallinity [7]. However, Wu et al. [9] pointed out that shearing results in cylindritic crystallites, of which the microscopic appearance is similar to that of transcrystalline morphology. Thus, the origin for the two crystallization processes is dierent. Transcrystallization requires a high heterogeneous nucleation ability of a specially prepared surface, whereas cylindritical crystallization is a self-nucleating process. Region 3 includes a further interface between transcrystalline and spherulitic morphology (termed the `second' interface). The thickness i,3 and the speci®c surface of this interface are determined by the size of spherulites formed. The in¯uence of the transcrystalline interphase on the composite properties is an important issue. The analytical detection of transcrystallinity in a real composite is required. Our work is focused on the detection of a transcrystalline interphase dependent on the ®bre sizing by investigations of the crystallization behaviour using DSC measurements. Since the sensitivity of DSC is limited a certain content of transcrystalline interphases is required. Therefore, composites both with a high speci®c ®bre surface and special crystallization conditions have been used at the investigations. However, a measurable eect of transcrystallization found at the special crystallization condition can be taken as an indication that the properties of ®bre surfaces are favourable for formation of a transcrystalline morphology. An enhancement of stress transfer via this transcrystalline interphase should result, independent of the thickness of this interphase. 2. Experimental
Fig. 1. Schematic illustration of the transcrystalline interphase in ®bre/polymer composites.
The experiments were carried out on unidirectional glass-®bre/PP composites with a ®bre content of 50 vol.%. To vary the speci®c ®bre-surface area, glass ®bres having dierent diameters were used. The glass ®bres were spun continuously in a unique melt-spinning equipment at the Institute of Polymer Research, Dresden (Germany). A maleic anhydride modi®ed polypropylene (MFI=36 g/10 min) was produced from isotactic homopolymer with the average molecular weight Mw=16104 by blending with maleic anhydride grafted PP (Polybond 3150 with Mw=9104). PP was available as a 100 mm thick ®lm.
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The composites were produced as sheets with a thickness of 0.25±0.3 mm by ®lm stacking and exact ®lament winding, which enabled a high homogeneity of impregnation and a void content below 1%. The processing conditions were 27 bar and 215 C kept for 5 min. Table 1 gives an overview of the composites investigated. The dierent ®bre diameter d fN results in a change of the speci®c ®bre surface s f by a factor of 2.5. The sizing based on g-aminopropyl-triethoxysilane (g-APS) was prepared either with a polyurethane ®lm former (sizing 1) or with a polypropylene ®lm former (sizing 2). For comparison the pure PP and the composite PP0A (with glass ®bres unsized and 13 mm thick) were investigated. The DSC measurements were performed using a DSC7 (Perkin-Elmer, USA) equipped with the Pyrissoftware in a temperature range from ÿ60 to 210 C. The temperature and heat of transition were calibrated with In and Pb standards. The temperature regime of the DSC measurements with both isothermal crystallization at 132 C and nonisothermal crystallization step is shown in Fig. 2. Additionally, the in¯uence of Tc,iso changed to 135 and 138 C on the crystallization behaviour was analysed by a further measurement. The heating rate used was always 20 K/ min and the molten state was kept for 2 min every time at 210 C to destroy the thermal history and in that way the pre-existing crystalline nuclei. The heat ¯ow related to the PP mass content is important for discussion of the results. Therefore, the PP mass content of all investigated samples was determined by Table 1 Samplesa Composite PPnN
N
d ®bre (mm) N
sfN/sfA
PPnA PPnB PPnC
A B C
13 22 32
1 0.6 0.4
a n, type of sizing: 0, unsized; 1, sizing 1; 2, sizing 2. N, designation for glass ®bre diameter.
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Fig. 2. Temperature regime of DSC measurements.
burning the samples in air using a Thermogravimetric Analyser TGA7 (Perkin-Elmer). The mass content of PP varied between 23 to 28 wt.% for the investigated composites, whereby the PP content of samples prepared from the same composite scattered by <2 wt.%. Some mechanical properties have been analysed for PPnA composites prepared by compounding and injection moulding. 3. Results The quantitative results of crystallization and heating behaviour are summarized in Table 2. The cooling scans of the dynamic crystallization are shown in Fig. 3a for pure PP and composites PP1N and in Fig. 3b for pure PP, PP0A and composites PP2N. The extrapolated onset-temperature of crystallization Tc;o (Fig. 3a) gives an information for the nucleation ability which can be changed by the special ®bre surface. A small increase of Tc;o by about 0.8 K observed in composites PP1N in comparison to pure PP is detectable, which is independent of the ®bre diameter. The experimental values of Tc;o (PP1N) scatter by 0.2 K.
Table 2 Characteristic values of melting and crystallization behavioura Sample
Pure PP PP0A PP1A PP1B PP1C PP2A PP2B PP2C a
Cooling (at 20 K/min)
Heating (aftercooling at 20 K/min)
Heating (after Tiso=132 C)
H(J/gPP)
Tc,m( C)
Tc,o( C)
H (J/gPP)
Tm( C)
H (J/gPP)
Tm( C)
97.0 96.5 94.9 95.6 95.4 95.8 94.7 94.7
117.1 121.1 117.4 117.1 117.5 117.9 118.4 117.5
121.3 127.0 122.1 121.8 122.2 124.1 124.1 122.4
100.9 101.6 96.7 98.3 98.3 97.9 97.1 97.0
161.6 162.2 160.9 161.5 160.8 161.7 161.4 161.3
109.8 108.5 106.5 107.1 107.9 106.1 104.9 106
167.1 155(sh)/ 166.5 166.5 166.8 166.4 155/ 166.5 166.5 166.7
H, transition heat of melting or crystallization, Tm/Tc,m, temperature of maximum of melting or crystallization peak or shoulder (sh); Tc,o, extrapolated onset-temperature of dynamical crystallization (Fig. 3a).
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The increase of 0.8 K can also be used as an indication of a small nucleation eect of the glass ®bres with sizing 1. The strong increase of Tc;o by 6 K was found at the composite PP0A with unsized ®bres. Sizing 2 shows also a higher nucleation ability connected with an increase of Tc;o by 3 K. Isothermal crystallization showed the same tendency regarding the nucleation ability for the dierent ®bre surfaces. At Tiso =132 C, remarkable eects were observed. Fig. 4a shows isothermal scans of pure PP and PP1N. A shorter induction time of the composites PP1N is observed in comparison to pure PP, but the change of the speci®c ®bre surface by the factor of 2.5 does not signi®cantly in¯uence the isothermal scans. However, the unsized ®bre and the sizing 2 result in strong changes of the crystallization behaviour as shown
in Fig. 4b. It is obvious, that the isothermal crystallization of the composite PP2A is a two-step process. The ®rst step is a fast process similar to the behaviour of composite PP0A followed by the second step at longer induction time. The behaviour of the second step is similar to that in the composite PP1A and also in pure PP. The ®rst step in the crystallization behaviour of PP2C is hardly detectable; however, the crystallization rate at the beginning is higher than in PP1C and the sigmoidal type of curve typically for pure PP and PP1N cannot be found for PP2C. The isothermal crystallization scans measured in the composites PP1A and PP2A at dierent temperatures are shown in Fig. 5. The higher nucleation rate observed at the beginning and caused by sizing 2 exists at all crystallization temperatures. Our results are in contrast to those of Devaux and Chabert [14], who did not
Fig. 3. Scans for the dynamical crystallization at a cooling rate of 20 K/min: (a) pure PP, PP1N; (b) pure PP, PP0A, PP2N.
Fig. 4. Scans for isothermal crystallization at 132 C: (a) pure PP, PP1N; (b) pure PP, PP0A, PP2N.
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observe any appreciable dierences between unsized ®bres and ®bres sized by a mixture of polyurethane and aminosilane, similar to our sizing 1. Fig. 6 shows the heating curves achieved for all composites with ®bre A after the isothermal crystallization at 132 C. In comparison, the signi®cant dierences in the crystallization behaviour (Fig. 4b) only re¯ect small changes in melting behaviour (Table 2). The overall PP crystallinity is about 52%. The melting heat Hf shows the same changes as the crystallization heat. The values obtained for PP2N and PP1N are similar and about 2% lower than that for PP0A, which is similar to pure PP. The qualitative melting behaviour for pure PP and composite PP1A is similar. The height of the peak at
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166 C is greater than that of composites PP0A and PP2A, which show additionally a shoulder and/or small peak at 4160 C. It is known that the main peak at 166 C is caused by melting of the monoclinic a crystallites, whereas the hexagonal b crystallites melt at about 157 C [15,16]. In this respect, a small content of b crystallites should exist in PP0A and PP2A, apparently connected with the transcrystallization process. This result could be con®rmed by WAXS investigations. A small re¯ex intensity of b crystallites was observed which is stronger in PP2A and PP0A than in PP1A. The glass-transition temperature of the PP matrix was also analysed for all samples, but a dierence could not be observed within the limit of accuracy. 4. Discussion
Fig. 5. Comparison of isothermal crystallization scans of PP1A (open circle o) with PP2A (®lled circle *) at 132, 135, and 138 C.
Fig. 6. Heating scans of pure PP and composites with dierently sized glass ®bres A after isothermal crystallization at 132 C.
The fast crystallization step clearly observed during the isothermal crystallization at 132 C for composites PP0A and PP2A (Fig. 4b) should be caused by the transcrystallization. The short induction time and the high rate at the beginning of crystallization support this hypothesis. This behaviour lead to the assumption that the whole PP matrix in the composite PP0A with the unsized ®bres isothermally crystallized at 132 C should be transcrystalline. Therefore, it was hypothesised that in the composite PP2A isothermally crystallized at 132 C both transcrystalline and spherulitic crystalline PP should occur. It can be concluded that the unsized ®bre surface has a very high nucleation ability for polypropylene. The technology of processing of composites requires in any case a sizing for glass ®bres. The surface of ®bre sized with sizing 1 does not remarkably in¯uence the morphology of the PP matrix. The absence of changes in the crystallization behaviour by changing the speci®c surface-area supports this result. In other words, the sizing 1 prevents the nucleation ability of unsized ®bre surface. In contrast, the dierent crystallization behaviour in the composites PP2N permits the conclusion that an own high nucleation ability of the sizing 2 exists. Fig. 7 shows the ®tted curves for both steps obtained using peak separation software [17] for the composites PP2A and PP2C. The quantitative results of the peaks separated are given in Table 3. The step 2 is connected with the spherulitic crystallization of a residual part of PP matrix with bulk behaviour. However, step 1 is of more interest. Its crystallization heat can be assumed as crystallization heat Htc of the transcrystalline interphase. The time of the heat ¯ow minimum is independent of the glass ®bre diameter d fN , and Htc/H is inversely proportional to d fN (Fig. 8). Using the three phase model [18±20] shown in Fig. 9 the thickness i of the transcrystalline interphase can be estimated. At the ®bre content of 50 Vol.% the value of dPP 2 can be determined
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Fig. 7. The two step-®t of isothermal crystallization at 132 C: (a) PP2A; (b) PP2C.
Table 3 Results of curve ®tting Composite
PP2A (13 mm) PP2B (22 mm) PP2C (32 mm)
Step 1
Step 2
Time of minimum tc,min (min)
Height of peak (W/gPP)
Htc/H (%)
Time of minimum tc,min (min)
Height of peak (W/gPP)
Hb/H (%)
1.8 1.8 1.8
0.28 0.14 0.1
32 25 11
4.0 3.6 4.3
0.3 0.3 0.32
68 75 89
f PP to be dPP 2 1:4d . Assuming that i i;2 << d2 , the following relation is valid:
d fibre i 4i Htc ÿ fibre 2 dfibre H d 4
1
The thickness of the transcrystalline interphase in the composite PP2A can be calculated from Eq. (1) using Htc =H 0:32 to be i 41 mm. With the assumption that the thickness i is independent of d f , the relation Htc =H 1=d fN found experimentally (Fig. 8) can be con®rmed by Eq. (1). For the composite PP0A the value of Htc =H was about 1, and consequently the whole PP between the ®bres whose average distance is about 6 mm is transcrystalline. The direct proof of the transcrystallinity by optical microscopy is hardly possible for samples with a glass ®bre content of 50 vol.%. For direct detection by optical microscopy a single ®bre sample was prepared. An unsized glass ®bre C and a glass ®bre C with sizing 1 were pressed together with a PP ®lm between two glass plates and melted at 200 C. A piece of this sample was isothermally crystallized at 132 C for 60 min in the DSC devices at analogous conditions to the DSC measurements. The morphology of crystallized PP shown in
Fig. 10 is dierent. Nearby the ®bre with sizing 1 spherulites were formed, whereas around the unsized ®bre transcrystallinity can be observed. 4.1. Mechanical properties in relation with transcrystallinity The in¯uence of the transcrystalline interphase on the mechanical properties has been controversially discussed in the literature (e.g. [21±26]). Jeng and Cheng [25] investigated the in¯uence of the transcrystalline interphase on the ¯exural failure of short carbon ®bre reinforced PEEK composites. The authors assumed the enhanced ¯exural de¯ection and strength of the composites to be caused by the transcrystalline interphase. However, they also observed that the eect of improved mechanical properties decreased with increased ®bre content. Wagner et al. [26] found that the variation of mechanical properties also depends on the crystallite type, either monoclinic or hexagonal, formed in the transcrystalline layer. Dierent possibilities for property alterations in¯uenced by transcrystalline interphases, can be discussed using the scheme given in Fig. 1. i. Region 1: Surface-induced perturbations depending on the type of sizing exist in region 1. Only these perturbations can act as nuclei for transcrystallization and
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Fig. 8. Results of the ®rst crystallization step of the isothermal crystallization at 132 C of composites PP2N: Htc =H and tc,min as a function of ®bre diameter.
Fig. 9. Schematic illustration of the three-phase model.
are responsible for the nature of this region and the ``®rst'' interface. The thickness i,1 is determined by the roughness of the ®bre/sizing surface and is independent of the thickness i,2 of the homogeneous part of the transcrystalline layer. The existence of transcrystallinity should result in a kind of `physical coupling' [25,27] between ®bre and polymer matrix. As a consequence, the stress transfer through the `®rst' interface might be enhanced. Using the fragmentation test Folkes and Wong [28] observed in glass-®bre/PP composites, an increase of the critical ®bre length as a result of to transcrystallinity in comparison to composites with PP matrix crystallized in very small spherulites. LopezManchado and Arroyo [29] support the assumption that the transcrystallinity permits an increasing stiness of composites. The problem is that an increasing stiness can also be caused by good adhesion, whereby the transcrystallinity and an enhanced adhesion could take place independently of each other [29]. ii. Region 2: The in¯uence of region 2 is connected with the change of properties of the transcrystalline crystallized polymer in comparison with that of the
Fig. 10. Optical microscopy images of dierently size-coated glass ®bres of type C embedded in PP matrix, measured in transmission with crossed polarizers: (a) unsized; (b) sizing 1.
spherulitically crystallized polymer, whereby the properties also depend on the crystallite type (a or b) [26]. Only a few experimental results of mechanical properties depending on the crystalline morphology have been published. Folkes and Hardwick [33] and Klein et al. [34] found an increase of the storage modulus E' by a factor of 2 and also increased shear and tensile yield strength for a transcrystalline layer compared with ®ne spherulitically crystallized layers. However, layers with complete transcrystalline morphology are more brittle, and the fracture energy is lower.
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iii. Region 3: The `second' interface formed in region 3 by the impingement of two crystallization fronts represents a weak fracture path [30±32]. This is caused both by low molecular weight fractions moving in front of the two growth fronts during the crystallization process [30,31] and by the lack of interlinking or entanglement of molecules between the two fronts impinged against each other [32]. The weakening eect depends on the content of low molecular weight PP species, on the thickness of the transcrystalline layer, and the size of spherulites. In summary, the existence of a transcrystalline interphase yields in contrast to each other's acting eects. The composite properties could be altered by it in dependence of the dominating eect. A further fact should be noticed. The transcrystallinity formed under quiescent conditions (transcrystallinity as de®ned in [9]) is to be distinguished from the `transcrystallinity' so-called also for cylindritic crystallites grown by stress-induced crystallization. In many papers [21,22] the `transcrystallization' is more a cylindritic growth of crystallites. It is unknown in what way both morphologies result in the same eects regarding the mechanical properties. Varga and Karger-Kocsis [15] pointed out that an improvement of the ®bre-matrix interfacial shear strength should not be expected from the presence of cylindritic crystallites grown by a self-nucleation process, whereas the transcrystalline zone formed by a heterogeneous nucleation on the ®bre surface should result in an enhancement of the stress transfer. The model composite PP0A made with unsized glass ®bres shows transcrystallinity. That means that the ®bre surface has a high nucleation ability and the `physical coupling' should improve the stress transfer through the interface. However, the acidic glass-®bre surface does not interact with an acidic-polymer matrix and, thus, the adhesion is very poor. The mechanical performance measured for injection moulded samples is very poor. Obviously, the eect of `physical coupling' by transcrystallinity cannot compensate the lack of adhesion. Sizing 1 alters the level of adhesive strength by introducing basic groups, which enable acid/base interactions at the interface. In this case, transcrystallization was found and the matrix was spherulitically crystallized. Obviously, sizing 1 inhibits the high heterogeneous nucleation ability of the pure glass-®bre surface. These composites show an average tensile strength of 57 MPa and poor Charpy unnotched impact toughness of 24 kJ/m2. An improvement of the mechanical properties has been observed in composites PP2A (sizing 2). The tensile strength increased to 86 MPa and the Charpy unnotched impact toughness to 61 kJ/m2. The adhesion to the ®bre surface has been improved by covalent bonds between NH2-groups of aminosilane and COOH-groups of maleic anhydride grafted ®lm former. On the other hand, sizing 2 has a high nucleation ability, and a detectable transcrystalline interphase was formed under special
crystallization conditions. The direct in¯uence of transcrystallinity on the mechanical performance could not be proved; however, the eect of the `physical coupling' [25,27] by transcrystallinity always exists independently of the thickness of the transcrystalline interphase. Similar results were also observed by Folkes and Hardwick [32]. The `physical coupling' by transcrystallinity and the enhanced adhesion improved the mechanical properties. A direct analysis of the in¯uence of transcrystallinity on the mechanical properties requires to investigate the mechanical properties of the composites with changed size and nature of transcrystalline interphases in future. 5. Conclusions In glass-®bre-reinforced materials with a high ®brevolume content, the in¯uence of dierent ®bre sizings on the crystallization behaviour was investigated using DSC. The thermal behaviour of the isothermal crystallization at 132 C indicates clearly two dierent crystallization processes in dependence on the nature of glass-®bre sizing. One process is very fast and has a short induction time, and should be connected with a high nucleation ability of the special ®bre surfaces. Thus, the fact supports the hypothesis of transcrystallization. This process was only observed for composites with unsized ®bres and ®bres sized with g-APS and a polypropylene ®lm former (sizing 2). The highest content of transcrystallinity was found using unsized ®bres. The crystallization behaviour of composites with dierent ®bre diameters con®rms that sizing 2 has an in¯uence on the crystalline morphology of the PP matrix correlated with varied speci®c-surface area. The thickness of the transcrystalline interphase for composite PP2A formed during isothermal crystallization at 132 C was estimated to be about 1 mm and is independent of the ®bre diameter. Moon [35] found an increase of the thickness of the transcrystalline layer with change of ®bre diameter from 19.5 up to 9.6 mm. This is in contrast with our result. The discussion of the mechanical properties related to transcrystallinity shows that it must be distinguished whether the whole PP matrix is transcrystalline or if only a transcrystalline interphase is formed. Depending of that, either an enhancement or a worsening of the composite properties can be obtained. It is also obvious that, besides the crystalline morphology, other properties of the interphase have to be considered in order to understand the alteration of mechanical performance as, for example, chemical or geometrical. For further investigations into this problem, new nondestructive methods to charaterize the interphase have been developed by Mai et al. [36]. The results obtained by these authors can be regarded as a potential for more knowledge of the nature and extent of interphases.
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Acknowledgements The research work was carried out in the frame of the Brite EuRam III programme, project no. BE 95-1710. The authors gratefully acknowledge the ®nancial support from the European Community. The authors thank H. Queck for preparation of composites, K. Arnhold and Th. Bayer for technical assistance in DSC measurements and peak separation, respectively. The authors are indebted to Dr. D. Jehnichen for the WAXS investigations. References [1] Fink BK, Mc Cullough RL. Interphase research issues. Composites, Part A 1990;30:1±2. [2] Drzal LT. Fiber±matrix interphase structure and its eect on adhesion and composite mechanical properties. In: Ishida H, editor. Proc. of ICCI III Controlled Interphases in Composites Materials, Cleveland. Elsevier Science Publishing Co., Inc., 1990. p. 309±320. [3] Ishida H, Bussi P. Surface induced crystallization in ultra-high modulus polyethylene ®ber reinforced polyethylene composite. In: Ishida H, editor. Proc. of ICCI III Controlled Interphases in Composites Materials, Cleveland. Elsevier Science Publishing Co.,Inc., 1990. p. 391±9. [4] Varga J, Karger-Kocsis J. The occurence of transcrystallization or row-nucleated cylindritic crystallization as a result of shearing in a glass-®ber-reinforced polypropylene. Composites Science and Technol 1993;48:191±8. [5] Thomason JL, van Rooyen AA. Transcrystallized interphase in thermoplastic composites Part I: In¯uence of ®bre type and crystallization temperature. J Mat Sci 1992;27:889±96. [6] Wang C, Liu C-R. Transcrystallization of polypropylene composites: nucleating ability of ®bres. Polymer 1999;40:289±98. [7] Thomason JL.van Rooyen AA. The transcrystallised interphase in thermoplastic composites. In: Ishida H, editor. Proc. of ICCI III Controlled Interphases in Composites Materials, Cleveland. Elsevier Science Publishing Co.,Inc., 1990. p. 423±30. [8] He T, Porter RS. Melt transcrystallization of polyethylene on high modulus polyethylene ®bres. J Appl Polym Sci 1988;35:1945±53. [9] Wu Ch-M., Chen M, Karger-Kocsis J. Transcrystallization in syndiotactic polypropylene induced by high-modulus carbon ®bers. Polymer Bulletin 1998;41:239±245. [10] Billon N, Haudin JM. In¯uence of transcrystallinity on DSC analysis of polymers: experimental and theoretical aspects. J Therm Analysis 1994;42:679±96. [11] Billon N, Magnet C, Haudin JM, Lefebvre D. Transcrystallinity eects in thin polymer ®lms. Experimental and theoretical approach Colloid Polym Sci 1994;272:633±54. [12] Wang Ch, Hwang LM. Transcrystallization of PTFE ®ber/PP composites (I). Crystallization kinetics and morphology. J Polym Sci, Polym Phys 1996;34:47±56. [13] Pompe G, MaÈder E. Experimental detection of transcrystalline interphases in glass ®bre/polypropylene composites by DSC. In: Proc. 12th ICCM, Paris, 1999. [14] Devaux E, Chabert B. Non-isothermal crystallization of glass ®bre reinforced polypropylene. Polymer Comm 1990;31:391±4. [15] Varga J, Karger-Kocsis J. Interfacial morphologies in carbon ®brereinforced polypropylene microcomposites. Polym 1995;36:4877±81.
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[16] Varga J, Karger-Kocsis J. Rules of supermolecular structure formation in sheared isotactic polypropylene melt. J Polym Sci, Polym Phys 1996;34:657±70. [17] Anon. Netzsch advanced software, demo version. Selb (Germany): Netzsch GeraÈtebau GmbH. [18] Lipatov Yu. Physical Chemistry of ®lled polymers. Internat. Polym. Sci and Techn Monograph No. 2, Moscow: Khimiya, 1977 (R.J. Moseley, Trans.). [19] Papanicolaou GC, Paipetis SA, Theocaris PS. The concept of boundary interphase in composite mechanics. Colloid Polym Sci 1978;256:625±30. [20] Sideridis E. The transverse elastic modulus of ®ber-reinforced composites as de®ned by the concept of interphase. J Appl Polym Sci 1993;48:243±55. [21] Hsiao BS, Chen EJH. Transcrystalline interphase in advanced polymer composites. In: Ishida H, editor. Proc. of ICCI III Controlled Interphases in Composites Materials, Cleveland 1990. Elsevier Science Publishing Co.,Inc., 1990. p. 613±21. [22] Beehag A, Ye L. Fibre/matrix adhesion in thermoplastic composites- is transcrystallinity a key? In: Proc. 11th ICCM, 1997, Australia. p. 339. [23] Huson MG, Mc Gill WJ. The eect of transcrystallinity on the behavior of ®bres in polymer matrices. J Polym Sci, Polym Phys 1985;23:121±8. [24] Schoolenberg GE, van Rooyen AA. Transcrystallinity in ®berreinforced thermoplastic composites. Composite Interfaces 1993;1:243±52. [25] Jeng C-C, Chen M. In¯uence of transcrystalline interphase on ¯exural failure of short carbon ®ber reinforced PEEK composites. In: Proc. 12th ICCM, Paris, 1999. p. 214. [26] Wagner HD, Lustiger A, Marzinsky CN, Mueller RR. Interlamellar failure at transcrystalline interfaces in glass/polypropylene composites. Comp Science and Techn 1993;48:181±4. [27] Clark Jr RC, Kander RG, Sauer BB. Nylon 66/poly(vinyl pyrrolidone) reinforced composites: 1. Interphase microstructure and evaluation of ®ber-matrix adhesion Composites, Part A 1999;30:27±36. [28] Folkes MJ, Wong WK. Determination of interfacial shear strength in ®bre-reinforced thermoplastic composites. Polymer 1987;28:1309±14. [29] Lopez-Manchado MA, Arroyo M. Crystallization kinetics of polypropylene Part 4: Eect of unmodi®ed and azide-modi®ed PET and PA short ®bres. Polymer 1999;40:487±95. [30] Schultze J, Nardin M. Interfacial adhesion. In: Ishida H, editor. Proc. of ICCI III Controlled Interphases in Composites Materials, Cleveland. Elsevier Science Publishing Co.,Inc., 1990. p. 561. [31] Teishev A, Marom G. The eect of transcrystallinity on the transverse mechanical properties of single-polymer polyethylene composites. J Appl Polym Sci 1995;56:959±66. [32] Folkes MJ, Hardwick ST. The mechanical properties of glass/ polypropylene multilayer laminates. J Mat Sci 1990;25:2598±606. [33] Folkes MJ, Hardwick ST. Direct study of structure and properties of transcrystalline layer. J Mat Sci Letters 1987;6:656±8. [34] Klein N, Marom G, Pegoretti A, Migliaresi C. Determining the role of interfacial transcrystallinity in composite materials by dynamic mechanical thermal analysis. Composites 1995;26:707± 12. [35] Moon C-K. Eect of molecular weight and ®ber diameter on the interfacial behavior in glass ®ber/PP composites. J Appl Polym Sci 1998;67:1191±7. [36] Mai K, MaÈder E, MuÈhle M. Interphase characterization in composites with new nondestructive methods. Composites Part A 1998;29A:1111±9.
www.elsevier.com/locate/compscitech
Experimental detection of a transcrystalline interphase in glass-®bre/polypropylene composites G. Pompe *, E. MaÈder Institute of Polymer Research Dresden, Hohe Str. 6, D-01069 Dresden, Germany Received 12 January 2000; received in revised form 28 April 2000; accepted 11 May 2000
Abstract Crystallization and melting behaviour of polypropylene in glass-®bre-reinforced composites have been investigated by means of dierential scanning calorimetry. Composites with dierently size-coated glass ®bres show dierent behaviour both in dynamic and isothermal crystallization. A high crystallization rate and a short induction time were observed for composites with unsized ®bres at isothermal crystallization. Glass ®bres sized with aminosilane and polypropylene ®lm-former revealed a quantitatively separable twostep crystallization process. With the assignment to the transcrystallization and spherulitic crystallization the thickness of the transcrystalline PP interphase in this composite could be estimated to about 1 mm. The in¯uence of transcrystallinity on the mechanical properties was discussed. # 2000 Elsevier Science Ltd. All rights reserved. Keywords: A. Glass ®bre; A. Polymer-matrix composites; B. Interphase; B. Mechanical properties; B. Thermal properties
1. Introduction The interphase in composites is the part of the matrix surrounding the ®bre surface and possessing local properties dierent from those of the bulk matrix. The size and type of interphase varies strongly and depends on the nature of the ®bre and its surface as well as on the polymer matrix [1,2]. Because stress transfer in the composites should be in¯uenced by the interphase, optimization of the mechanical properties of composites requires an extensive knowledge of the behaviour of the interphases and its eect on mechanical performance. The interphase results from perturbations induced at ®bre surfaces and propagated away from the ®bre surface into the matrix. Under appropiate conditions [3±7] a highly oriented, so-called transcrystalline interphase can appear in composites with semicrystalline polymer (e.g. PP) as the matrix. The structure and size of transcrystalline interphases are both system-and condition-speci®c. Some in¯uences are known, for instance the surface energy of ®bre, changed by the ®bre type, the sizing, and coupling agent [5±7], and the topography of the ®bre (level of * Corresponding author. Tel.: +49-351-4658-255; fax: +49-3514658-284. E-mail address: [email protected] (G. Pompe).
roughness) [6]. It was also observed that transcrystallization can only occur at crystallization temperatures lower than the critical temperature Ttc cr [5±7]. Our investigations were focused on the question of which surface treatment of the glass ®bres changes the crystallization behaviour in conjunction with a changed crystalline morphology of the matrix nearby the ®bre surface compared with the bulk matrix behaviour. A ®rst attempt to use dierential scanning calorimetry (DSC) for the detection of transcrystallinity was made by He and Porter [8]. In this paper, a quantitative analysis of systematic DSC measurements will be reported for PP composites with dierently sized glass ®bres. At constant ®bre volume fraction (0.5) the speci®c glass ®bre surface is varied by using dierent glass ®bre diameters. 1.1. Transcrystallinity Wu et al. [9] and also Billon et al. [10,11] de®ned transcrystallization as a nucleation-controlled process taking place under quiescent conditions in a semicrystalline polymer in contact with other materials, e.g. ®bres. At high heterogeneous nucleation ability of the surface the closely spaced nuclei hinder the lateral extension and force growth in one direction, namely perpendicular to the
0266-3538/00/$ - see front matter # 2000 Elsevier Science Ltd. All rights reserved. PII: S0266-3538(00)00120-2
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®bre surfaces. Once nucleation has occured, the growth of spherulites in the matrix and of the transcrystalline layer on the surface takes place at the same rate [5± 7,12]. A dierent nucleation rate of both crystallization mechanism is the necessary condition for the formation of a detectable extent of transcrystallinity. This dierence can amount to about three orders of magnitude found, for PTFE-®bre/PP composites [12]. Fig. 1 shows the scheme of the interphase with transcrystalline morphology in a ®bre/semicrystalline-polymer composite [13]. Three regions of dierent thickness, i,k might be distinguished. Region 1 of thickness i,1 includes the interface between the ®bre and the transcrystalline part of PP matrix (termed the `®rst' interface). Both the thickness i,1 and the speci®c surface of the ®rst interface will be determined by the roughness of the ®bre/sizing surface. Region 2 is the homogeneous part of the transcrystalline layer without boundaries (thickness i,2). Growth of the transcrystalline layer will be halted by the impingement on the spherulites nucleated and grown in the bulk matrix. As the induction period for surface nucleation is shorter [6,8,12], and the growth rate of both crystalline structures is the same [5±7,12],the thickness i,2 will be determined above all by the ability of nucleation to occur on the surface in comparison to that in the bulk. This means that i,2 can be in¯uenced by changes in the nucleation rate on the ®bre surface or sizing with a high functionality for the nucleation process as well as by changes in the nucleation rate in the bulk matrix, e.g. nucleating agents, molecular weight. The conditions of crystallization also play an important role. It was
observed that transcrystallization under quiescient conditions can only take place at crystallization temperatures, Tc, lower than a critical temperature Ttc cr [5,7]. A brief application of shear stresses along the ®bre/matrix interface permits the formation of morphology similar to that of the transcrystalline, also at Tc>Ttc cr [7]. These observations result in the assumption that stress applied externally or formed internally is a condition for the formation of transcrystallinity [7]. However, Wu et al. [9] pointed out that shearing results in cylindritic crystallites, of which the microscopic appearance is similar to that of transcrystalline morphology. Thus, the origin for the two crystallization processes is dierent. Transcrystallization requires a high heterogeneous nucleation ability of a specially prepared surface, whereas cylindritical crystallization is a self-nucleating process. Region 3 includes a further interface between transcrystalline and spherulitic morphology (termed the `second' interface). The thickness i,3 and the speci®c surface of this interface are determined by the size of spherulites formed. The in¯uence of the transcrystalline interphase on the composite properties is an important issue. The analytical detection of transcrystallinity in a real composite is required. Our work is focused on the detection of a transcrystalline interphase dependent on the ®bre sizing by investigations of the crystallization behaviour using DSC measurements. Since the sensitivity of DSC is limited a certain content of transcrystalline interphases is required. Therefore, composites both with a high speci®c ®bre surface and special crystallization conditions have been used at the investigations. However, a measurable eect of transcrystallization found at the special crystallization condition can be taken as an indication that the properties of ®bre surfaces are favourable for formation of a transcrystalline morphology. An enhancement of stress transfer via this transcrystalline interphase should result, independent of the thickness of this interphase. 2. Experimental
Fig. 1. Schematic illustration of the transcrystalline interphase in ®bre/polymer composites.
The experiments were carried out on unidirectional glass-®bre/PP composites with a ®bre content of 50 vol.%. To vary the speci®c ®bre-surface area, glass ®bres having dierent diameters were used. The glass ®bres were spun continuously in a unique melt-spinning equipment at the Institute of Polymer Research, Dresden (Germany). A maleic anhydride modi®ed polypropylene (MFI=36 g/10 min) was produced from isotactic homopolymer with the average molecular weight Mw=16104 by blending with maleic anhydride grafted PP (Polybond 3150 with Mw=9104). PP was available as a 100 mm thick ®lm.
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The composites were produced as sheets with a thickness of 0.25±0.3 mm by ®lm stacking and exact ®lament winding, which enabled a high homogeneity of impregnation and a void content below 1%. The processing conditions were 27 bar and 215 C kept for 5 min. Table 1 gives an overview of the composites investigated. The dierent ®bre diameter d fN results in a change of the speci®c ®bre surface s f by a factor of 2.5. The sizing based on g-aminopropyl-triethoxysilane (g-APS) was prepared either with a polyurethane ®lm former (sizing 1) or with a polypropylene ®lm former (sizing 2). For comparison the pure PP and the composite PP0A (with glass ®bres unsized and 13 mm thick) were investigated. The DSC measurements were performed using a DSC7 (Perkin-Elmer, USA) equipped with the Pyrissoftware in a temperature range from ÿ60 to 210 C. The temperature and heat of transition were calibrated with In and Pb standards. The temperature regime of the DSC measurements with both isothermal crystallization at 132 C and nonisothermal crystallization step is shown in Fig. 2. Additionally, the in¯uence of Tc,iso changed to 135 and 138 C on the crystallization behaviour was analysed by a further measurement. The heating rate used was always 20 K/ min and the molten state was kept for 2 min every time at 210 C to destroy the thermal history and in that way the pre-existing crystalline nuclei. The heat ¯ow related to the PP mass content is important for discussion of the results. Therefore, the PP mass content of all investigated samples was determined by Table 1 Samplesa Composite PPnN
N
d ®bre (mm) N
sfN/sfA
PPnA PPnB PPnC
A B C
13 22 32
1 0.6 0.4
a n, type of sizing: 0, unsized; 1, sizing 1; 2, sizing 2. N, designation for glass ®bre diameter.
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Fig. 2. Temperature regime of DSC measurements.
burning the samples in air using a Thermogravimetric Analyser TGA7 (Perkin-Elmer). The mass content of PP varied between 23 to 28 wt.% for the investigated composites, whereby the PP content of samples prepared from the same composite scattered by <2 wt.%. Some mechanical properties have been analysed for PPnA composites prepared by compounding and injection moulding. 3. Results The quantitative results of crystallization and heating behaviour are summarized in Table 2. The cooling scans of the dynamic crystallization are shown in Fig. 3a for pure PP and composites PP1N and in Fig. 3b for pure PP, PP0A and composites PP2N. The extrapolated onset-temperature of crystallization Tc;o (Fig. 3a) gives an information for the nucleation ability which can be changed by the special ®bre surface. A small increase of Tc;o by about 0.8 K observed in composites PP1N in comparison to pure PP is detectable, which is independent of the ®bre diameter. The experimental values of Tc;o (PP1N) scatter by 0.2 K.
Table 2 Characteristic values of melting and crystallization behavioura Sample
Pure PP PP0A PP1A PP1B PP1C PP2A PP2B PP2C a
Cooling (at 20 K/min)
Heating (aftercooling at 20 K/min)
Heating (after Tiso=132 C)
H(J/gPP)
Tc,m( C)
Tc,o( C)
H (J/gPP)
Tm( C)
H (J/gPP)
Tm( C)
97.0 96.5 94.9 95.6 95.4 95.8 94.7 94.7
117.1 121.1 117.4 117.1 117.5 117.9 118.4 117.5
121.3 127.0 122.1 121.8 122.2 124.1 124.1 122.4
100.9 101.6 96.7 98.3 98.3 97.9 97.1 97.0
161.6 162.2 160.9 161.5 160.8 161.7 161.4 161.3
109.8 108.5 106.5 107.1 107.9 106.1 104.9 106
167.1 155(sh)/ 166.5 166.5 166.8 166.4 155/ 166.5 166.5 166.7
H, transition heat of melting or crystallization, Tm/Tc,m, temperature of maximum of melting or crystallization peak or shoulder (sh); Tc,o, extrapolated onset-temperature of dynamical crystallization (Fig. 3a).
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The increase of 0.8 K can also be used as an indication of a small nucleation eect of the glass ®bres with sizing 1. The strong increase of Tc;o by 6 K was found at the composite PP0A with unsized ®bres. Sizing 2 shows also a higher nucleation ability connected with an increase of Tc;o by 3 K. Isothermal crystallization showed the same tendency regarding the nucleation ability for the dierent ®bre surfaces. At Tiso =132 C, remarkable eects were observed. Fig. 4a shows isothermal scans of pure PP and PP1N. A shorter induction time of the composites PP1N is observed in comparison to pure PP, but the change of the speci®c ®bre surface by the factor of 2.5 does not signi®cantly in¯uence the isothermal scans. However, the unsized ®bre and the sizing 2 result in strong changes of the crystallization behaviour as shown
in Fig. 4b. It is obvious, that the isothermal crystallization of the composite PP2A is a two-step process. The ®rst step is a fast process similar to the behaviour of composite PP0A followed by the second step at longer induction time. The behaviour of the second step is similar to that in the composite PP1A and also in pure PP. The ®rst step in the crystallization behaviour of PP2C is hardly detectable; however, the crystallization rate at the beginning is higher than in PP1C and the sigmoidal type of curve typically for pure PP and PP1N cannot be found for PP2C. The isothermal crystallization scans measured in the composites PP1A and PP2A at dierent temperatures are shown in Fig. 5. The higher nucleation rate observed at the beginning and caused by sizing 2 exists at all crystallization temperatures. Our results are in contrast to those of Devaux and Chabert [14], who did not
Fig. 3. Scans for the dynamical crystallization at a cooling rate of 20 K/min: (a) pure PP, PP1N; (b) pure PP, PP0A, PP2N.
Fig. 4. Scans for isothermal crystallization at 132 C: (a) pure PP, PP1N; (b) pure PP, PP0A, PP2N.
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observe any appreciable dierences between unsized ®bres and ®bres sized by a mixture of polyurethane and aminosilane, similar to our sizing 1. Fig. 6 shows the heating curves achieved for all composites with ®bre A after the isothermal crystallization at 132 C. In comparison, the signi®cant dierences in the crystallization behaviour (Fig. 4b) only re¯ect small changes in melting behaviour (Table 2). The overall PP crystallinity is about 52%. The melting heat Hf shows the same changes as the crystallization heat. The values obtained for PP2N and PP1N are similar and about 2% lower than that for PP0A, which is similar to pure PP. The qualitative melting behaviour for pure PP and composite PP1A is similar. The height of the peak at
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166 C is greater than that of composites PP0A and PP2A, which show additionally a shoulder and/or small peak at 4160 C. It is known that the main peak at 166 C is caused by melting of the monoclinic a crystallites, whereas the hexagonal b crystallites melt at about 157 C [15,16]. In this respect, a small content of b crystallites should exist in PP0A and PP2A, apparently connected with the transcrystallization process. This result could be con®rmed by WAXS investigations. A small re¯ex intensity of b crystallites was observed which is stronger in PP2A and PP0A than in PP1A. The glass-transition temperature of the PP matrix was also analysed for all samples, but a dierence could not be observed within the limit of accuracy. 4. Discussion
Fig. 5. Comparison of isothermal crystallization scans of PP1A (open circle o) with PP2A (®lled circle *) at 132, 135, and 138 C.
Fig. 6. Heating scans of pure PP and composites with dierently sized glass ®bres A after isothermal crystallization at 132 C.
The fast crystallization step clearly observed during the isothermal crystallization at 132 C for composites PP0A and PP2A (Fig. 4b) should be caused by the transcrystallization. The short induction time and the high rate at the beginning of crystallization support this hypothesis. This behaviour lead to the assumption that the whole PP matrix in the composite PP0A with the unsized ®bres isothermally crystallized at 132 C should be transcrystalline. Therefore, it was hypothesised that in the composite PP2A isothermally crystallized at 132 C both transcrystalline and spherulitic crystalline PP should occur. It can be concluded that the unsized ®bre surface has a very high nucleation ability for polypropylene. The technology of processing of composites requires in any case a sizing for glass ®bres. The surface of ®bre sized with sizing 1 does not remarkably in¯uence the morphology of the PP matrix. The absence of changes in the crystallization behaviour by changing the speci®c surface-area supports this result. In other words, the sizing 1 prevents the nucleation ability of unsized ®bre surface. In contrast, the dierent crystallization behaviour in the composites PP2N permits the conclusion that an own high nucleation ability of the sizing 2 exists. Fig. 7 shows the ®tted curves for both steps obtained using peak separation software [17] for the composites PP2A and PP2C. The quantitative results of the peaks separated are given in Table 3. The step 2 is connected with the spherulitic crystallization of a residual part of PP matrix with bulk behaviour. However, step 1 is of more interest. Its crystallization heat can be assumed as crystallization heat Htc of the transcrystalline interphase. The time of the heat ¯ow minimum is independent of the glass ®bre diameter d fN , and Htc/H is inversely proportional to d fN (Fig. 8). Using the three phase model [18±20] shown in Fig. 9 the thickness i of the transcrystalline interphase can be estimated. At the ®bre content of 50 Vol.% the value of dPP 2 can be determined
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Fig. 7. The two step-®t of isothermal crystallization at 132 C: (a) PP2A; (b) PP2C.
Table 3 Results of curve ®tting Composite
PP2A (13 mm) PP2B (22 mm) PP2C (32 mm)
Step 1
Step 2
Time of minimum tc,min (min)
Height of peak (W/gPP)
Htc/H (%)
Time of minimum tc,min (min)
Height of peak (W/gPP)
Hb/H (%)
1.8 1.8 1.8
0.28 0.14 0.1
32 25 11
4.0 3.6 4.3
0.3 0.3 0.32
68 75 89
f PP to be dPP 2 1:4d . Assuming that i i;2 << d2 , the following relation is valid:
d fibre i 4i Htc ÿ fibre 2 dfibre H d 4
1
The thickness of the transcrystalline interphase in the composite PP2A can be calculated from Eq. (1) using Htc =H 0:32 to be i 41 mm. With the assumption that the thickness i is independent of d f , the relation Htc =H 1=d fN found experimentally (Fig. 8) can be con®rmed by Eq. (1). For the composite PP0A the value of Htc =H was about 1, and consequently the whole PP between the ®bres whose average distance is about 6 mm is transcrystalline. The direct proof of the transcrystallinity by optical microscopy is hardly possible for samples with a glass ®bre content of 50 vol.%. For direct detection by optical microscopy a single ®bre sample was prepared. An unsized glass ®bre C and a glass ®bre C with sizing 1 were pressed together with a PP ®lm between two glass plates and melted at 200 C. A piece of this sample was isothermally crystallized at 132 C for 60 min in the DSC devices at analogous conditions to the DSC measurements. The morphology of crystallized PP shown in
Fig. 10 is dierent. Nearby the ®bre with sizing 1 spherulites were formed, whereas around the unsized ®bre transcrystallinity can be observed. 4.1. Mechanical properties in relation with transcrystallinity The in¯uence of the transcrystalline interphase on the mechanical properties has been controversially discussed in the literature (e.g. [21±26]). Jeng and Cheng [25] investigated the in¯uence of the transcrystalline interphase on the ¯exural failure of short carbon ®bre reinforced PEEK composites. The authors assumed the enhanced ¯exural de¯ection and strength of the composites to be caused by the transcrystalline interphase. However, they also observed that the eect of improved mechanical properties decreased with increased ®bre content. Wagner et al. [26] found that the variation of mechanical properties also depends on the crystallite type, either monoclinic or hexagonal, formed in the transcrystalline layer. Dierent possibilities for property alterations in¯uenced by transcrystalline interphases, can be discussed using the scheme given in Fig. 1. i. Region 1: Surface-induced perturbations depending on the type of sizing exist in region 1. Only these perturbations can act as nuclei for transcrystallization and
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Fig. 8. Results of the ®rst crystallization step of the isothermal crystallization at 132 C of composites PP2N: Htc =H and tc,min as a function of ®bre diameter.
Fig. 9. Schematic illustration of the three-phase model.
are responsible for the nature of this region and the ``®rst'' interface. The thickness i,1 is determined by the roughness of the ®bre/sizing surface and is independent of the thickness i,2 of the homogeneous part of the transcrystalline layer. The existence of transcrystallinity should result in a kind of `physical coupling' [25,27] between ®bre and polymer matrix. As a consequence, the stress transfer through the `®rst' interface might be enhanced. Using the fragmentation test Folkes and Wong [28] observed in glass-®bre/PP composites, an increase of the critical ®bre length as a result of to transcrystallinity in comparison to composites with PP matrix crystallized in very small spherulites. LopezManchado and Arroyo [29] support the assumption that the transcrystallinity permits an increasing stiness of composites. The problem is that an increasing stiness can also be caused by good adhesion, whereby the transcrystallinity and an enhanced adhesion could take place independently of each other [29]. ii. Region 2: The in¯uence of region 2 is connected with the change of properties of the transcrystalline crystallized polymer in comparison with that of the
Fig. 10. Optical microscopy images of dierently size-coated glass ®bres of type C embedded in PP matrix, measured in transmission with crossed polarizers: (a) unsized; (b) sizing 1.
spherulitically crystallized polymer, whereby the properties also depend on the crystallite type (a or b) [26]. Only a few experimental results of mechanical properties depending on the crystalline morphology have been published. Folkes and Hardwick [33] and Klein et al. [34] found an increase of the storage modulus E' by a factor of 2 and also increased shear and tensile yield strength for a transcrystalline layer compared with ®ne spherulitically crystallized layers. However, layers with complete transcrystalline morphology are more brittle, and the fracture energy is lower.
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iii. Region 3: The `second' interface formed in region 3 by the impingement of two crystallization fronts represents a weak fracture path [30±32]. This is caused both by low molecular weight fractions moving in front of the two growth fronts during the crystallization process [30,31] and by the lack of interlinking or entanglement of molecules between the two fronts impinged against each other [32]. The weakening eect depends on the content of low molecular weight PP species, on the thickness of the transcrystalline layer, and the size of spherulites. In summary, the existence of a transcrystalline interphase yields in contrast to each other's acting eects. The composite properties could be altered by it in dependence of the dominating eect. A further fact should be noticed. The transcrystallinity formed under quiescent conditions (transcrystallinity as de®ned in [9]) is to be distinguished from the `transcrystallinity' so-called also for cylindritic crystallites grown by stress-induced crystallization. In many papers [21,22] the `transcrystallization' is more a cylindritic growth of crystallites. It is unknown in what way both morphologies result in the same eects regarding the mechanical properties. Varga and Karger-Kocsis [15] pointed out that an improvement of the ®bre-matrix interfacial shear strength should not be expected from the presence of cylindritic crystallites grown by a self-nucleation process, whereas the transcrystalline zone formed by a heterogeneous nucleation on the ®bre surface should result in an enhancement of the stress transfer. The model composite PP0A made with unsized glass ®bres shows transcrystallinity. That means that the ®bre surface has a high nucleation ability and the `physical coupling' should improve the stress transfer through the interface. However, the acidic glass-®bre surface does not interact with an acidic-polymer matrix and, thus, the adhesion is very poor. The mechanical performance measured for injection moulded samples is very poor. Obviously, the eect of `physical coupling' by transcrystallinity cannot compensate the lack of adhesion. Sizing 1 alters the level of adhesive strength by introducing basic groups, which enable acid/base interactions at the interface. In this case, transcrystallization was found and the matrix was spherulitically crystallized. Obviously, sizing 1 inhibits the high heterogeneous nucleation ability of the pure glass-®bre surface. These composites show an average tensile strength of 57 MPa and poor Charpy unnotched impact toughness of 24 kJ/m2. An improvement of the mechanical properties has been observed in composites PP2A (sizing 2). The tensile strength increased to 86 MPa and the Charpy unnotched impact toughness to 61 kJ/m2. The adhesion to the ®bre surface has been improved by covalent bonds between NH2-groups of aminosilane and COOH-groups of maleic anhydride grafted ®lm former. On the other hand, sizing 2 has a high nucleation ability, and a detectable transcrystalline interphase was formed under special
crystallization conditions. The direct in¯uence of transcrystallinity on the mechanical performance could not be proved; however, the eect of the `physical coupling' [25,27] by transcrystallinity always exists independently of the thickness of the transcrystalline interphase. Similar results were also observed by Folkes and Hardwick [32]. The `physical coupling' by transcrystallinity and the enhanced adhesion improved the mechanical properties. A direct analysis of the in¯uence of transcrystallinity on the mechanical properties requires to investigate the mechanical properties of the composites with changed size and nature of transcrystalline interphases in future. 5. Conclusions In glass-®bre-reinforced materials with a high ®brevolume content, the in¯uence of dierent ®bre sizings on the crystallization behaviour was investigated using DSC. The thermal behaviour of the isothermal crystallization at 132 C indicates clearly two dierent crystallization processes in dependence on the nature of glass-®bre sizing. One process is very fast and has a short induction time, and should be connected with a high nucleation ability of the special ®bre surfaces. Thus, the fact supports the hypothesis of transcrystallization. This process was only observed for composites with unsized ®bres and ®bres sized with g-APS and a polypropylene ®lm former (sizing 2). The highest content of transcrystallinity was found using unsized ®bres. The crystallization behaviour of composites with dierent ®bre diameters con®rms that sizing 2 has an in¯uence on the crystalline morphology of the PP matrix correlated with varied speci®c-surface area. The thickness of the transcrystalline interphase for composite PP2A formed during isothermal crystallization at 132 C was estimated to be about 1 mm and is independent of the ®bre diameter. Moon [35] found an increase of the thickness of the transcrystalline layer with change of ®bre diameter from 19.5 up to 9.6 mm. This is in contrast with our result. The discussion of the mechanical properties related to transcrystallinity shows that it must be distinguished whether the whole PP matrix is transcrystalline or if only a transcrystalline interphase is formed. Depending of that, either an enhancement or a worsening of the composite properties can be obtained. It is also obvious that, besides the crystalline morphology, other properties of the interphase have to be considered in order to understand the alteration of mechanical performance as, for example, chemical or geometrical. For further investigations into this problem, new nondestructive methods to charaterize the interphase have been developed by Mai et al. [36]. The results obtained by these authors can be regarded as a potential for more knowledge of the nature and extent of interphases.
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