Barrier and Mechanical Properties of Pulp Fiber/Polymer Laminates and Blends.

M. KROOK [a,b]

M. S. HEDENQVIST [b]

A.-C. ALBERTSSON [b]

A. HELLMAN [b]

T. IVERSEN [c]

U. W. GEDDE [b*]

Paperboard laminates coated with two grades of poly([epsilon]-caprolactone) (PCL), poly(hydroxy bulyrate-co-valerate) (PHBV) or a liquid crystalline copolyester (LOP) were prepared by compression molding, and the influence of the processing conditions and polymer content of the laminate on the laminate properties was studied. Ligno-cellulose fiber/polymer blends were prepared from wet pulps and PCL and PHBV. The morphology, water vapor transmission rates, creasability, curl and twist and mechanical properties of the laminates and blends were studied. LOP and slowly cooled high molar mass PCL laminated paperboards showed the best creasing properties and the paperboards that were penetrated by the polymer showed the smallest degree of curl and twist. Extensive penetration occurred during compression molding of the paperboard with the low molar mass PCL at all temperatures and with PHBV and LCP at the higher molding temperatures. The water vapor transmission rates ranged from 1 to 300 times that of polyethylene depending on the polymer used and on the thermal history. In the case of blends, competitive properties were obtained only in those with a high polymer content. The laminate stiffness decreased and the strength increased in two polymer concentration regions, at 20 wt% due to fiber-fiber separation and at 60 wt% due to phase inversion.

INTRODUCTION

Biodegradable polymers are of interest in the context of environmental protection and conservation [1]. In the 1980s, packaging waste was generally disposed of by landfilling or incineration. This caused serious water pollution and waste location problems when the landfills became saturated, With incineration, air pollution also became a problem [2]. Today. waste management is extremely important and there is a need for environmentally compatible solutions because the use of nondegradable components in packagings is increasing. The replacement of metals and glass in packaging by polymers reduces manufacturing costs, the weight-to-strength ratio and the amount of energy and water required [3]. These advantages have led to an increased use of polymeric containers (e.g. poly(ethylene terephthalate)-bottles) and polymer coatings. Biodegradable polymers that are comparable with existing packaging polymers in terms of barrier properties, processability and mechanical properties constitute a tremendous possibility.

In packaging requiring extremely good barrier properties towards e.g. permanent gases, current biodegradable polymers are insufficient. For such applications, liquid crystalline polymers have been shown to be an interesting alternative [4-8]. These polymers exhibit extremely good barrier properties and are at present the only polymer alternative to aluminum and glass.

In this work, the use of biodegradable polymers, two different grades of poly([epsilon]-caprolactone) and a commercial grade of poly(hydroxybutyrate-co-valerate), as barrier components on or in paperboard were studied. The possibility of obtaining a continuous barrier component in paperboard, either by simple coating or by blending the polymer with cellulose fibers, was explored. The latter method will, if it can be successfully implemented in the early stage of papermaking, make the coating step unnecessary and permit the production of a biodegradable mono-material having sufficient barrier properties. The properties of paperboard coated with a liquid crystalline polymer were also studied.

EXPERIMENTAL

Materials

Two different grades of poly([epsilon]-caprolactone were kindly supplied by Union Carbide as granules, TONE P-300 ([M.sub.w] = 10,000 g/mol, [M.sub.w]/[M.sub.n] = 1.4 and TONE P767E (M.sub.w] = 69,000 g/mol, [M.sub.w]/[M.sub.n] = 1.7). These grades are hereafter denoted PCL10 and PCL69 respectively. A commercial grade of poly(hydroxybutyrate-co-valerate), BIOPOL D400P ([M.sub.w] = 320,000 g/mol, [M.sub.w]/[M.sub.n] = 2.9; valerate content = 7-9 mol%), hereafter referred to as PHBV, was kindly supplied by Monsanto (Belgium) in powder form. The liquid crystalline polymer (LCP) studied was an experimental grade of the thermotropic liquid crystalline copolyester, Vectra RD 501, kindly supplied by Ticona, Germany. The LCP was delivered as granules. The kraft pulp ligno-cellulosic fibre material, having k = 60, was kindly supplied by AssiDoman Frovi, Sweden. The cellulosic material used for the coating process was a four-ply paperboard of Duplex quality with a basis weight of 260 g/[m.sup.2]. For the polymer/cellulose-fiber blends, the cellulosic material was delivered as a 2.9% wet pulp taken directly from the mill stream.

Sample Preparation

Cooting

The polymers were compression molded onto paper-board using a water-cooled Schwabenthan compression molding machine (Polystat 400s). PCL10, PCL69 and PHBV were molded onto the paperboard using two 15-[micro]m-thick metal plates made of Sandvik-stainless steel (11 R 51) as release agent. The steel plates provided an extremely smooth surface. Polyimide sheets, Kapton l00HN, 25 [micro]m thick, were used as release agent for the LCP. The PCL10 and PCL69 granulate materials were ground in a Retsch ZM 1 mill grinder prior to the compression molding. To prevent melting, the polymers were cooled in liquid nitrogen prior to grinding and were exposed to [CO.sub.2]-ice during grinding. All laminates were kept for three minutes at the initial molding temperature before cooling. Except during quenching, the samples were pressurized throughout the compression molding cycle. The PCL69 was compression molded at temperatures ranging between 70[degrees]C and 130[degrees]C at a pressure of 20 MPa and subsequently cooled with a cooling rate of 4[degrees]C/min to 25[degrees]C, where the temperature was held constant for 7 min before cooling to 20[degrees]C. These samples are referred to as "slowly cooled PCL69". PCL69 was also compression molded at temperatures between 60[degrees]C and 100[degrees]C at a pressure of 10 MPa followed by quenching to -50[degrees]C. The latter step was accomplished by placing the sample on an aluminum plate cooled with liquid nitrogen. The quenched samples are referred to as "rapidly cooled PCL69". PCL10 was compression molded at 100[degrees]C at a pressure of 20 MPa and was subsequently cooled at 4[degrees]C/min to 25[degrees]C, where the temperature was held constant for 7 min before final cooling to 20[degrees]C. PHBV was compression molded at temperatures ranging between 150[degrees]C and 220[degrees]C at a pressure of 10 MPa. The samples were cooled at 7[degrees]C/min to 55[degrees]C where the temperature was held constant for 5 min before subsequent cooling to 2500 at a rate of 2-3[degrees]C/min. LOP was compression molded at temperatures ranging between 180[degrees]C and 260[degrees]C at a pressure of 20 MPa. The samples were cooled at 13[degrees]C/min to 100[degrees]C followed by a slower cooling to ambient temperature. These samples are referred to as "laminates".

Blending

The cellulosic wet pulp (2.9%) was diluted to 0.29% in water and mixed at 2000 rpm for 5 mm. PCL10, PCL69 and PHBV were mixed in powder form into the diluted pulps. The amount of pulp needed to make a standard laboratory sheet, i.e. 165 mm X 165 mm 60 g/[m.sup.2] paper, being added to each blend. Four different sheets containing 10, 30, 50 and 70% by weight of polymer were made for each of the three polymer grades. Manual mixing of the polymer/pulp provided a uniform distribution of the polymer in the pulp. Square sheets were made from the polymer/pulp using a wire screen under suction according to SCANC 26:76 [9]. To prevent the polymer powder from sticking to the stirrer and the walls of the sheet former, the former was carefully washed with water during the entire sheet-forming process. The sheets were subsequently stacked according to SCAN-C 26:76 [9] and then placed in a Lorentzen and Wettre compression machine at 0.35 MPa for 5.5 mm. Subsequently the sheets were firmly attached to the drying plates and , before a second compression, which lasted for 2 min, the order of the sheets was reversed and all the blotters were replaced. The drying plates and the paper sheets were separated from the blotters and conditioned (2300, 50% relative humidity) as specffied in SCAN-P2 [9]. The sheets were dried for 24 h and shrinkage was prevented by using drying plates during the entire drying period. The dried polymer/pulp sheets were placed between glass-fiber-reinforced poly(tetrafluoroethylene) sheets before the final compression molding in a Pasadena Hydraulics 0230H compression machine. The sheets containing PCL10 and PCL69 were compression molded at 100[degrees]C for 2 min at 10 MPa and the sheets containing PHBV were compression molded at 180[degrees]C for 2 min at 10 MPa, in both cases followed by cooling at 15[degrees]C/min to room temperature. These samples are referred to as "blends".

Methods

Differential Scanning Calorimetry 

The melting endotherms of the polymers were obtained using a Mettler DSC 820 and the recorded heats of fusion ([delta][h.sub.f]) were transformed into mass crystallininity ([w.sub.c]) using the total enthalpy method (10), with a 155.9 kJ/kg (11) (PCL) as the heat of fusion ([delta][[h.sup.0].sub.f]) for 100% crystalline polymer at the equilibrium melting point (349.3 K) (12) (PCL), according to:

[W.sub.c] = [delta][h.sub.f]/[delta][[h.sup.0].sub.f] - [[[integral of].sup.349.3].sub.[t.sub.1]] ([C.sub.pa] - [C.sub.pc]) dT (1)

where [T.sub.1] is an arbitrary temperature below the melting range and, [C.sub.pa] and [C.sub.pc] are the specific heats of the amorphous and crystalline components respectively. Data for [C.sub.pa] and [C.sub.pc] provided by Wunderlich [13] on PCL have been used. Due to a lack of experimental data, the crystallinity of PHBV was estimated without the integral correction term in Eq 1 and using [delta][[h.sup.0].sub.f] = 146.8 kJ/kg, which is the heat of fusion of 100% crystalline poly(hydroxybutyrate) (PHB) (14). It was believed that the error due to the neglect of the temperature-dependence of [delta][[h.sup.0].sub.f] would be small. The Thomson-Gibbs equation (15) was used to calculate the crystal thickness ([L.sub.c]):

[L.sub.c] = 2[[zigma].sub.[epsilon][[T.sup.0].sub.m]/([[T.sup.0].sub.m] - [T.sub.m]) [delta][[h.sup.0].sub.f][P.sub.c] (2)

where [[zigma].sub.[epsilon] is the fold surface free energy (101 mJ/[m.sup.2] (PCL) [12] and 38 mJ/[m.sup.2] (PHBV) [14]), [T.sub.m] and [[T.sup.0].sub.m] are the experimental and equilibrium melting peak temperatures respectively ([[T.sup.0].sub.m] = 470.2 K (PHB) [14] and [p.sub.c] is the crystal density. The amorphous interlayer thickness was estimated according to:

[L.sub.a] = [L.sub.c](1 - [W.sub.c])/[W.sub.c]/[p.sub.c]/[p.sub.a] (3)

where [p.sub.a] is the amorphous density. For PCL [p.sub.c] = 1187 kg/[m.sup.3] and [p.sub.a] = 1094 kg/[m.sup.3] [16], and for PHBV [p.sub.c] = 1260 kg/[m.sup.3] and [p.sub.a] = 1177 kg/[m.sup.3] [17] were used. The volume fraction of amorphous material was calculated from the mass crystallinity and crystal and amorphous densities:

[v.sub.a] = 1 - [W.sub.c]/[p.sub.a]/[W.sub.c]/[p.sub.c] + 1 - [W.sub.c]/[p.sub.a] (4)

Water Vapor Transmission Rate

The water vapor transmission rate was measured using a Mocon Permatran-W Twin at 23[degrees]C and 100% relative humidity, according to ASTM F 1249-90. The water vapor transmission rate Q was normalized with respect to the polymer layer thickness t to yield Q[degrees]:

Q[degrees] = Q X t (5)

The film thickness was obtained from cross-sections of freeze-fractured samples using a JEOL JSM-5400 scanning electron microscope.

Creasability Testing

The creasability was determined at 25[degrees]C and 50% relative humidity in a clamshell press, Econo cut model 20 X 27. Three samples, 15 mm X 60 mm in size, from each of the coated paperboards were tested. The crease depth was calculated according to

[t.sub.c] = 0.78 + [t.sub.s] - 0.7 (6)

where [t.sub.c] is the crease depth, [t.sub.s] is the thickness of the sample and the constant 0.78 is the matrix thickness, all in mm. The constant 0.7 corresponds to the height of the creasing rule substracted from the knife height.

Water Absorbency 

The water absorbency was obtained as the [Cobb.sub.60] value according to SCAN-P 12:64 [9]. The [Cobb.sub.60] value in g [H.sub.2]O/[m.sup.2] was calculated according to:

X = 100(a - b) (7)

where a and b are the weights (g) of the test piece after 60 s of wetting and before wetting respectively.

Tearing Resistance

The tearing resistance was determined at 25[degrees]C and 50% relative humidity according to ISO 1974[12.3] using a Lorentzen & Wettre tear tester of the Elmendorf pendulum type. The mass of the pendulum was 68.8 g. Samples 62 mm X 50 mm in size were cut from the molded sheets of the blends as well as from polymer-free paper sheets. An average value of the tearing resistance was calculated according to

[T.sup.b] = S X P (8)

where the tearing resistance ([T.sup.b]) is given in mN, s is the mean scale reading in the tearing direction and p is the pendulum factor (p = 2).

Tensile Testing

The stress-strain behavior were obtained at 25[degrees]C and 50% relative humidity according to SCAN-P 38:80 [9] using an Alwetron TCT 10 tensile tester. Three samples 50 mm X 15 mm in size were tested for each blend as well as paper sheets containing no polymer. The thicknesses of the samples were calculated as the average of five measurements. The strain rate was 100 mm/min and the strain was measured from the separation of the clamps.

RESULTS AND DISCUSSION

The structures of the coated paperboards could be divided into two main categories: [1] laminates with polymer that had penetrated the paperboard (Fig. la); [2] laminates with polymer forming a coating layer (Fig. 1b). PHBV- and LOP-coatings. which were applied at or above 180[degrees]C and 220[degrees]C respectively, and PCL10 belong to the first category whereas the rest of the laminates belong to the second category. The low melt viscosities of LOP and PCL10 explains why these polymers readily penetrate the paperboard. In the case of PHBV, it is believed that the penetration, which occurred only at the higher molding temperatures was, due to a lowering of melt viscosity induced by molar mass reduction caused by thermal degradation involving random chain scission of ester groups [17, 18]. The thickness of the pure polymer coating layer as assessed by scanning electron microscopy decreased at the onset of penetration of the polymer through the paperboard for PHBV and LOP (Fig. 2).

The low molar mass PCL10 had a very low melt viscosity and it was necessary to use more than 33 wt% polymer to obtain a continuous polymer coating layer. The laminate containing 50 wt% PCL10 showed spots of uncoated fibers whereas the laminate with 60 Wt% PCL10 was uniformly coated with polymer. The water vapor transmission rate of the laminate with 60 wt% PCL10 was considerably lower than that of the laminates with the lower PCL10 content (Fig. 3). The spots of uncoated fibers were evidently regions with a greater water vapor transmission. The higher water vapor transmission rate of the laminate with 50 wt% PCL10 than to that of the laminate with 33 wt9/o PCL10 shown in Fig. 3 seems anomalous. There is a significant uncertainty in the assessment of the thickness of the polymer layer, particularly in the case of the 33 wt9/o PCL10 laminate. PCL10 is a very brittle polymer which also readily cracks when subjected to small stresses and strains. Cracks were more readily formed in the laminate with 50 wt% PCL10, and this may explain its higher measured water vapor transmission rate. The laminate with 60 wt% PCL10 showed a greater degree of polymer penetration into the paperboard and cracks formed in the outer regions were arrested in the paperboard-polymer core, thus preventing transport of water vapor through the laminate through continuous cracks.

Figure 4 shows the water vapor transmission rates for laminates with continuous and uniform polymer coatings. The aliphatic polyesters showed water vapor transmission rates 10-300 times that of polyethylene. The molding temperature had no effect except in the case of LCP (Fig. 4). Molding at temperatures below the crystal melting point, [tilde]220[degrees]C (Fig. 5), resulted in brittle, non-continuous layers with many regions of high water vapor transmission. Molding above 220[degrees]C resulted in uniformly coated paperboard with a low and molding-temperature-independent water vapor transmission rate (Fig. 4). The aliphatic polyesters showed variations in Q[degrees] within an order of magnitude which evidently correlated negatively with the degree of crystallinity of the polymer. Rapid cooling and an increase in molar mass favor a lower degree of crystallinity and these materials exhibited higher transmission rates.

Figure 6 presents water vapor transmission rates as a function of the amorphous volume fraction ([v.sub.a]) for the aliphatic polyesters. It has been known for several decades (see e.g. the review by Hedenqvist and Gedde [19] that the crystallites are inpenetrable for all non-reactive molecules except for the smallest (e.g. He). Lasoski and Cobbs [20] found that the water vapor permeability of poly(ethylene terephthalate), polyamide 6,10 and polyethylene followed the relationship:

P = [P.sub.a][[v.sub.a].sup.2] (9)

where [P.sub.a] is the permeability of the fully amorphous polymer. This relationship can also be obtained theoretically assuming the validity of Henry's law (P = D S, where D = diffusivity and S = solubility) and that S = [v.sub.a] Sa and D = [v.sub.a] [D.sub.a]. By letting Q[degrees] equal to P, the present data displayed in Fig. 6 fit Eq 9 fairly well although the [v.sub.a] ranges covered by each polymer were not sufficient for a critical test of the equation. The following values for the amorphous permeabilities (Pa)were obtained from the fitting: [P.sub.a] = 1243 g mm/([m.sup.2] day atm) (PCL69, rapidly cooled); 804 g mm/([m.sup.2] day atm) (PCL69, slowly cooled) and 195 g mm/([m.sup.2] day atm) (PHBV). The data for PCL thus indicate that the post-molding cooling rate determining the temperature range at which the material crystallized had some influence on the structure of the amorphous component.

Figure 7 shows the amorphous permeability plotted as a function of the amorphous layer thickness. None of the studied polymers showed any pronounced variation in the amorphous permeability with [L.sub.a]. The decrease in [P.sub.a] with increasing [L.sub.a] for rapidly cooled PCL seems artificial and not related to [L.sub.a]. The high [P.sub.a] value found for the low molar mass PCL is due to the great concentration of cilia (loose chain ends) in its amorphous phase. The molar mass of crystalline stems is for this particular polymer of the order of 3000 g/mol, which may be compared with the average molar mass of 10,000 g/mol. The concentration of cilia is thus on an average 1 out of three amorphous entries. Slowly and rapidly cooled PCL showed very different amorphous permeabilities despite the fact that they have similar amorphous interlayer spacings (Fig. 7). It thus seems that the amorphous chains are less well packed in rapidly cooled PCL than in slowly cooled PCL. The low amorphous permeability of the PH BV samples may to some extent be explained by their very thin amorphous interlayers (Fig. 7). The amorphous chain segments are shorter and more constrained by the crystallites, and this reduces both D and S (Fig. 8). It is important to emphasize that among the different PHBV's there was no clear trend [P.sub.a] vs. [L.sub.a] despite the fact that [L.sub.a] varied by a factor of two among the PHBV samples studied.

The very slow crystallization of PHBV is a practical problem when PHBV- coated packaging materials are being used shortly after manufacture. The liquid barrier properties of PHBV are usually determined for a "fully" semicrystalline material, but as can be seen in Fig. 9, the material may be almost completely amorphous immediately after cooling from the melt, and hence may possess inferior barrier properties. It should be pointed out that the crystallinity immediately after the cooling procedure is above 50% for PHBV molded at 220[degrees]C, suggesting the occurrence of extensive molar mass reduction during the molding. The degradation facilitated crystallization increasing the degree of crystallinity but yielded a material with thinner crystals. This conclusion is drawn from the observation that the melting point was lower for samples molded at 220[degrees]C than for samples molded at a lower temperature.

In Fig. 10, the crease depth is plotted as a function of the molding temperature. Since the polymer has a lower modulus than the fiber component, the crease depth obtained at constant force increases with increasing polymer coating thickness, as is evident from a comparison of the data presented in Figs. 2 and 10. In general, the crease depth obtained for the different samples was acceptable for bending with limited force. The PCL69-coated samples (slowly and rapidly cooled) delaminated during creasing with the exception of the PCL69 samples molded at 80 and 100[degrees]C and then slowly cooled, which showed signs of delamination only after bending. Rapidly cooled PCL69 exhibited more extensive delamination than the slowly cooled samples, both before and after bending. Due to the high fracture toughness of PCL69, no cracks developed during bending regardless of the cooling conditions. The creasing of the more brittle PCLO led to serious delamination from the paperboard accompanied by the formation of grea t many cracks. PHBV molded at lower temperatures delaminated extensively, but cracks were absent. Delamination decreased with increasing molding temperature indicating a better polymer-paperboard adhesion after molding at higher temperatures due to more complete wetting resulting from the lower melt viscosity. Molding at temperatures above 200[degrees]C caused degradation and the laminates delaminated during bending. Extensive degradation at 220[degrees], led to a poor surface, and the creasing properties could not be determined. During creasing of LCP-coated samples molded above the melting point, the polymer showed good adhesion to the paperboard with no delamination or crack initiation. However, creasing the laminates molded at temperatures below 220[degrees]C caused the LCP component to break into two pieces and cellulose fiber breakage was also observed. Cracks in the polymer component were observed in samples molded at 240[degrees]C and 260[degrees[C.

An important paperboard property is the curl and twist resistance. Curl and twist are due to differences in thermal expansion properties of polymer and the paperboard, leading to shrinkage-induced mechanical stresses exerted on the paperboard from the solidifying polymer. The lowest curl and twist were obtained with the PCL10 which penetrated the paperboard and with the LCP laminates. Penetration involves a more homogeneous distribution of polymer through the laminate thickness, and this reduces the tendency for curl and twist. Liquid crystalline polymers are known to exhibit a low thermal expansion coefficient [21], which is beneficial for obtaining a laminate with low curl and twist.

The thicknesses of the blends are shown as a function of the content of polymer in Fig. 11. The thickness increased only moderately with increasing polymer content below 30 wt% (PCL69) and below 50 wt% (PCL10 and PHBV). The differences between PCL69 and PCL1O may be due to differences in melt viscosity. The blends containing the lower viscosity material, PCL10, became more densely packed than the PCL69/paper blends. At low contents of polymer, the polymer distribution in the paperboard was heterogeneous and therefore a content of polymer of 70 wt% or more was therefore needed to give a smooth polymer surface without areas of uncoated fibers.

Figure 12 shows the water absorbency as a function of the polymer content for the blend samples. At a low mass fraction of polymer ([less than] 30 wt%) the water uptake data showed a significant variation between different samples due to the non-uniform polymer distribution. At 50 wt% polymer a continuous layer was formed, but there were still areas of uncoated fibers causing a high water absorbancy. At 70 wt% polymer, the surface became smooth and since all fibers were coated the amount of water absorbed was low. To be a commercially interesting alternative in packaging, the amount of polymer should not exceed 10 wt% in the polymer/fiber blends, which is far below the polymer contents giving satisfactory results here. At a polymer content of 70 wt%, the blend may be considered to be a fiber-reinforced polymer rather than a polymer coated/impregnated fiber material. The water absorbency for PCL10 was lower than for PCL69 and PHBV, owing to its higher crystallinity. Furthermore, the lower melt viscosity of PC L10 yielded a more uniform distribution of polymer in the blend that gave low water vapor transmission rates.

Figure 13 shows the stress/strain behavior of the PCL69 blends. Above 10wt% polymer, the Young's modulus and the fracture stress decreased and the fracture strain increased with increasing polymer content. The disturbance of the fiber-fiber bonds by the addition of polymer seems to have a greater impact on the modulus than on the fracture strain, as indicated by the greater decrease in modulus occurring at a lower polymer content than the associated increase in fracture strain. The PHBV blends showed a higher modulus than both the PCL10 and PCL69 blends, which may be explained by the higher polarity of PHBV giving better adhesion between polymer and fiber (Fig. 14). The moduli of the pure polymers are also different; the modulus of PHBV being higher than of PCL [22,23]. The difference in polymer-fiber adhesion also leads to a higher fracture stress for PHBV blends (Fig. 15). The lower fracture strain for the PHBV blends may be explained by the low fracture strain of the pure PHBV [20] (Fig. 16). The toughest blends are obtained by mixing cellulosic fibers with PCL69 as is shown by their higher fracture strain (Fig. 16) and larger tearing resistance (Fig . 17). This is clearly due to the higher fracture toughness of PCL69 [23]. At low contents of polymer, the blends containing the more brittle PCL10 show a higher or similar toughness to the blends containing PCL69 (Figs. 16, 17) due to the more uniform distribution of the polymer in the PCL10 composites. The most dramatic changes in mechanical properties occurred when the polymer content was increased from 10% to 30% and from 50% to 70%. The first change occurs due to the separation of the fibers caused by the addition of polymer and the second change is due to phase inversion when the polymer phase became continuous.

CONCLUSIONS

Two different types of coating structures were observed in the paperboard laminates. The first type consisted of a polymer coating on top of the paperboard. Besides the coating, the second type invoked polymer penetration into the paperboard. The high molar mass poly([epsilon]-caprolactone) always belonged to the first type whereas the low molar mass poly([epsilon]-caprolactone) belonged to the second type. Poly(hydroxybutyrate-co-valerate) and liquid crystalline copolyester-coated paperboards exhibited both types of structures depending on the molding temperature. Water vapor transmission rates ranged between 1 and 300 times that of polyethylene, being lowest for the liquid crystalline copolyester-coated paperboard and highest for the rapidly cooled high molar mass poly([epsilon]-caprolactone) paperboard. The slowly cooled high molar mass poly([epsilon]-caprolactone) paperboard and the liquid crystalline copolyester-coated paperboard showed the best creasing properties. Commercially interesting properties o f the blends as packaging materials were obtained only for systems with large polymer contents. Water absorbency, Young's modulus and fracture stress decreased and fracture strain increased with increasing polymer content in the blends.

ACKNOWLEDGMENTS

Mr. Lars Norberg, STEI, Mr. Birger Edholm, Ms. Jinghua Wan and Mrs. Christina Olmers, Packforsk-Swedish Packaging Institute, Stockholm, are thanked for providing guidance during the course of this work. Mr. Magnus Johansson, Assi-Doman Frovi, Sweden, is thanked for fruitful discussions.

(a.) Packforsk-Swedish Packaging Research Institute P. 0. Box 9, s-164 93 Kista, Sweden

(b.) Royal Institute of Technology, Department of Polymer Technology S-100 44 Stockholm, Sweden

(c.) Address: STFI, Box 5604. S-114 86 Stockholm, Sweden.

(*.) To whom correspondence should he addressed.

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