Amandine Viretto

C2MA Rodolphe Sonnier
email: rodolphe.sonnier@mines-ales.fr

Aurélie Taguet

Belkacem Otazaghine

Laurent Ferry

José-Marie Lopez-Cuesta

Christian Lagreve

Jose-Marie Lopez-Cuesta

Christian Lagrève

Thermal degradation of polyesters filled with magnesium dihydroxide and magnesium oxide

Keywords: flame retardancy, polyesters, poly(butylene terephthalate), magnesium dihydroxide

à la diffusion de documents scientifiques de niveau recherche, publiés ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

INTRODUCTION

The wide use of polymer materials in everyday life is led by their outstanding combination of properties, low weight, low cost, and ease of processing. However, polymers are also known for their high flammability usually accompanied by the release of corrosive or toxic gases and smoke during combustion. Accordingly, improving the flame retardancy of polymers is a key issue for extending their use to most applications [START_REF] Camino | Overview of fire retardant mechanisms[END_REF][START_REF] Laoutid | New prospects in flame retardant polymer materials: from fundamentals to nanocomposites[END_REF].

The scientific literature has presented various strategies for improving polymer flame retardancy, which depend mainly on the nature of the polymer [START_REF] Levchik | Flame retardancy of thermoplastic polyestersa review of the recent literature[END_REF][START_REF] Weil | Flame retardants in commercial use or development for polyolefins[END_REF]. The main route to reduce the flammability of polymers is the incorporation of fillers and additives, which can act according to different modesof-action in condensed or gas phases [START_REF] Laoutid | New prospects in flame retardant polymer materials: from fundamentals to nanocomposites[END_REF]. Until recently, many flame retardants were halogen based. These compounds act mainly in gas phase by scavenging the high energy radicals (H° and OH°) responsible for flame propagation. This flame inhibition effect leads to incomplete combustion with intense smokes and toxic products [START_REF] Hull | Fire retardant action of mineral fillers[END_REF]. Moreover, some halogenated flame retardants are considered toxic [START_REF] Silva | Brominated organic micropollutantsigniting the flame retardant issue[END_REF] and persistent, with a significant long term environmental impact [START_REF] Wit | An overview of brominated flame retardants in the environment[END_REF]. The alternatives to halogenated flame retardants include metal hydroxides, hydrocarbonates, borates, phosphorous compounds, nitrogen compounds, silicon compounds, and nanoparticles [START_REF] Laoutid | New prospects in flame retardant polymer materials: from fundamentals to nanocomposites[END_REF].

In this paper, we will mainly focus on the hydrated mineral fillers and especially on magnesium dihydroxide (MDH). Metal hydroxides provide effective flame retardancy when they are used in polymers. Aluminim trihydroxide (ATH) and MDH are the most widely used [START_REF] Hollingbery | The thermal decomposition of huntite and hydromagnesitea review[END_REF]. MDH is a wellknown flame retardant and smoke suppressor in thermoplastic polymers (polyolefin [START_REF] Rothon | Flame retardant effects of magnesium hydroxide[END_REF] and polyamide [START_REF] Hornsby | Thermal decomposition behaviour of polyamide fire-retardant compositions containing magnesium hydroxide filler[END_REF][START_REF] Fei | Synergistic effects of novolac-based char former with magnesium hydroxide in flame retardant polyamide-6[END_REF]) and in thermoplastic elastomers [START_REF] Durin-France | Magnesium hydroxide/zinc borate/talc compositions as flame-retardants in EVA copolymer[END_REF]13]. High loading levels of 50-60wt.% are generally required to achieve good fire performances in polymers [START_REF] Hollingbery | The thermal decomposition of huntite and hydromagnesitea review[END_REF]. MDH decomposes through an endothermic reaction to give magnesium oxide (MgO) with water release. The decomposition starts at about 300-330 °C giving off about 31% of its mass as water. The associated endothermic energy of this reaction is about 1450 J/g [START_REF] Rothon | Flame retardant effects of magnesium hydroxide[END_REF]. MDH acts as a fire retardant through five main mechanisms [13]: [START_REF] Camino | Overview of fire retardant mechanisms[END_REF] the dilution of the organic phase by replacement of the polymer matrix; (2) its endothermic decomposition, which slows down the thermal decomposition of the polymer; (3) the release of water, which dilutes the fuels in vapor phase; [START_REF] Weil | Flame retardants in commercial use or development for polyolefins[END_REF] the absorption of energy due to the heat capacity of MDH and MgO, which reduces the thermal energy available to degrade the polymer (the heat capacity of the filler is generally lower than that of the polymer; nevertheless, the residue from filler decomposition remains in the condensed phase and absorbs heat, whereas the polymer is completely or almost completely volatilized during pyrolysis); and (5) the mineral residue, which accumulates at the surface and insulates the substrate from the heat source. If MDH or MgO promotes also charring, the organomineral residue should be much more efficient as barrier layer. It is the objective of this article to study the char promotion of both fillers in relation to the structure of the host polymer.

Various polyesters with different chemical structures were selected. The polybutylene terephtalate (PBT), an aromatic polyester, was especially investigated. PBT is a saturated thermoplastic polyester produced by polycondensation reaction of terephtalic acid and 1,4-butanediol [START_REF] Carothers | Studies on polymerization and ring formation. I. An introduction to the general theory of condensation polymers[END_REF]. This polymer has remarkable properties such as high rate of crystallization, good solvent resistance, thermal stability, and excellent processability [START_REF]Handbook of Thermoplastic Polyesters[END_REF]. PBT is quite flammable, and it is necessary to improve its flame retardancy.

The thermal stability and flammability of thermoplastic polyesters, mainly filled with MDH, are investigated in this study. Many polymers, such as polyamides and polyesters, are sensitive to hydrolysis, which leads to an earlier degradation of material properties. This explains the great interest addressed to the hydrolytic degradation mechanisms and kinetics of these polymers [START_REF] Bellenger | Lifetime prediction in the hydrolytic ageing of polyesters[END_REF][START_REF] Golovoy | Hydrolysis of polycarbonate/polybutylene terephthalate blend[END_REF]. The degradation of PBT properties is due to hydrolysis of ester groups in the backbone leading to chain scissions [START_REF] Kelleher | Hydrolysis of poly(butylene terephthalate)[END_REF]. This hydrolysis phenomenon could be promoted by the water release from hydroxides. Using metal oxides instead of hydroxides may limit the hydrolytic degradation of PBT while acting as a flame retardant. Recent literature reports an interest in the use of oxide particles [START_REF] Gallo | Halogen-free flame retarded poly(butylene terephthalate) (PBT) using metal oxides/PBT nanocomposites in combination with aluminium phosphinate[END_REF][START_REF] Laoutid | Red phosphorus/aluminium oxide compositions as flame retardants in recycled poly(ethylene terephthalate)[END_REF][START_REF] Laachachi | Effect of Al2O3 and TiO2 nanoparticles and APP on thermal stability and flame retardance of PMMA[END_REF][START_REF] Lewin | Catalysis of intumescent flame retardancy of polypropylene by metallic compounds[END_REF] as synergistic agents in addition to the usual fire retardant additives. Metal oxide particles showed an ability to change the pathway of polymer decomposition and promote the formation of a protective barrier during the thermal degradation. This barrier reduces the mass transfer from the sample to the flame and the heat transfer from the flame to the condensed phase [START_REF] Gallo | Halogen-free flame retarded poly(butylene terephthalate) (PBT) using metal oxides/PBT nanocomposites in combination with aluminium phosphinate[END_REF]. In this work, MgO in the micronic range was used in replacement of MDH to avoid hydrolysis of PBT, and its effect on the fire behavior was studied.

In the first part of this article, the influence of the presence of MDH on the thermal degradation of 12 polymers was investigated. This study allowed selecting PBT as good charring polymer in the presence of MDH. The second part attempted to optimize the char promotion effect in PBT through two ways. In the first approach, MDH content was varied to identify a threshold value above which charring is not further promoted. The second approach dealt with the replacement of MDH by MgO to avoid early degradation of PBT by hydrolysis.

EXPERIMENTAL PART

Materials

The polymers used in this study were various aromatic and aliphatic polyesters, polymers with pendant ester groups and polycarbonate (PC). According to their structure, these polymers were divided into four series A, B, C, and D (Table I). Series A gathers thermoplastic aliphatic polyesters (polylactique acid

Processing

The polymers filled with MDH or MgO were obtained using a micro-compounder (DSM Xplore, 15 cm 3 ). The blends were prepared under the following conditions: 60 rpm, 10 min, and processing temperatures T melt described in Table I. It should be noticed that the decomposition temperature range of MDH is above the processing temperature of all the studied polymers, except LCP. In the case of this last formulation, MDH probably starts to degrade during processing.

Polybutylene terephtalate was also compounded with 10 or 25wt.% of MDH using a co-rotating twinscrew extruder Clextral BC21 and injection-molded using a 50-t Krauss-Maffei equipment (Munich, Germany) to prepare specimens for cone calorimeter tests. It is believed that compounding of PBT-MDH is quite difficult because of the hydrolytic degradation of PBT in the presence of MDH. We have not observed any problem to prepare these blends, maybe because of the moderate content of MDH.

Characterizations

Thermogravimetric analyses (TGAs) were carried out using a PerkinElmer Pyris-1 TGA thermo balance for samples of 10 ± 2 mg. For each sample, a run was carried out at a heating rate of 10 °C/min from room temperature to 900 °C under nitrogen flow of 20 mL/min. The gases released from the pyrolysis were analyzed by Fourier transformed infrared (FTIR) spectroscopy. The gases were sent to the gas cell through a transfer line heated at 200 °C. The FTIR analyses were carried out using a Bruker IFS66 spectrometer (wavelength range of 4000-400 cm À1 and resolution of 2 cm À1 ). Moreover, the FTIR spectrometer was also used in transmission mode (KBr pellets) to characterize the solid residues. The solid residues were also studied using energy-dispersive X-ray spectroscopy (environmental scanning electron microscopy was equipped with an Oxford energy dispersive X-ray spectrometry (EDX) sensor).

Flammability was investigated using a pyrolysis combustion flow calorimeter (PCFC), which was developed by Lyon et al. [START_REF] Lyon | Pyrolysis combustion flow calorimetry[END_REF]. The sample (2 ± 0.5 mg) was first heated from 80 to 750 °C at 1 °C/s in a pyrolyzer under nitrogen flow, and the degradation products were sent to a combustor where they mixed with oxygen in excess at 900 °C. In such conditions, these products were fully oxidized. Heat release rate (HRR) was then calculated by oxygen depletion according to Huggett's relation (1 kg of consumpted oxygen corresponds to 13.1 MJ of released energy) [START_REF] Huggett | Estimation of rate of heat release by means of oxygen consumption measurements[END_REF]. Very small samples (several milligrams) could be analyzed using PCFC. At this microscale, some flame retardant effects (as barrier effect) are not effective. Nevertheless, the temperature of the different steps of degradation, the heat released during each step, and the char promotion effect are very important parameters to evaluate the flammability of a material and the PCFC allows studying all these parameters. The evolved gases during thermal degradation were also analyzed using pyrolysis-gas chromatography/mass spectrometry (Py-GC/MS) analysis. A Pyroprobe 5000 pyrolyzer (CDS Analytical, Oxford, PA, USA) was used to flash pyrolyze the samples in a helium environment. The pyrolyzer is supplied with an electrically heated platinum filament. One coil probe enables the pyrolysis of samples (less than 1 mg) placed in a quartz tube between two pieces of quartz wool. The same sample was heated at 900 °C. The pyroprobe 5000 is interfaced to a 450-GC gas chromatograph (Varian) by means of a chamber heated at 270 °C. In the oven, the initial temperature of 70 °C was raised to 250 °C at 10 °C/min. The column is a Varian Vf-5ms capillary column (30 m × 0.25 mm), and helium (1 mL/min) was used as the carrier gas; a split ratio was set to 1:50. The gases were introduced from the GC transfer line to the ion trap analyzer of the 240-MS mass spectrometer (Varian) through the direct-coupled capillary column. Identification of the products was achieved comparing the observed mass spectra to those of the National Institute of Standards and Technology (NIST) mass spectral library.

Fire behavior was studied using a cone calorimeter (Fire Testing Technology), which is a powerful tool to investigate the fire behavior of polymers. A horizontal sample sheet of 100 × 100 × 4 mm 3 was placed at 2.5 cm below a conic heater and isolated by rock wool. The samples were exposed to 50 kW/m 2 irradiance in well-ventilated conditions (air rate 24 L/s) and in the presence of a spark igniter to force the ignition. HRR was determined according to oxygen depletion (Huggett's relation [START_REF] Huggett | Estimation of rate of heat release by means of oxygen consumption measurements[END_REF]). Mass loss, CO and CO 2 release rates, and smoke opacity were also recorded continuously. This test was performed according to the ISO 5660 standard.

RESULTS AND DISCUSSION

Influence of MDH on thermal degradation of various polyesters

In this section, the influence of MDH on the degradation of various polyesters from the series A to D was investigated by TGA, PCFC, and Py-GC/MS.

The addition of MDH significantly changes the thermal stability of the studied polymers. Figures 1234show the thermogravimetric curves and the HRR curves of four polymer matrices: PLA, PBT, EMA, and LCP with and without 10wt.% of MDH. They are representative of the series A, B, C, and D, respectively.

Pyrolysis combustion flow calorimeter and TGAs are in good agreement for all studied (pure and filled) polymers. Contrarily to TGA, the release of non-combustible gases is not detected in PCFC. Therefore, since peaks of HRR at low temperature could be observed for PLA, PBT, and LCP filled with MDH (series A, B, and D, respectively), it is confirmed that the earlier degradation (also noticed from thermogravimetric curves) of these materials corresponds not only to the water release from MDH but also to the pyrolysis of the polymers. Pure PLA (series A) exhibits only one peak of degradation centered on 390 °C (Figure 1(a)). When it is filled with MDH, the degradation occurs in several steps. The main peak of degradation of PLA (380 °C) vanishes (only a very small peak is also observed at 350 °C), and the polymer degrades almost completely at lower temperature (300 °C). The first peak is detected much earlier (the first peak at 280 °C). The two first peaks in the range 270-350 °C correspond not only to water release from MDH decomposition but also to early pyrolysis of PLA (as confirmed by the presence of a peak of HRR in the same range in PCFC -Figure 1(b)). PLA is very sensitive to hydrolysis at high temperature. Therefore, the release of water from the decomposition of MDH leads to its quick degradation.

The last peak is observed at higher temperature (450-500 °C) when PLA is filled with MDH. PCFC curve (Figure 1(b)) confirms that it corresponds to the release of combustible gases. This peak can be assigned to the degradation of an organic residue that will be called char in the following. Therefore, it can be assumed that MDH promotes charring, but the thermal stability of this char is low according to its decomposition temperature.

Polybutylene terephtalate (series B) is also very sensitive to hydrolysis, and its degradation also occurs earlier in the presence of MDH (Figure 2(a)). Nevertheless, the main peak is still observed at the same temperature as for pure PBT (approximately 400 °C). Like PLA, PBT filled with MDH exhibits an additional step of degradation at high temperature, indicating that MDH also promotes charring of aromatic polyesters. However, its temperature of degradation is significantly higher (600 °C). Such observation is particularly interesting while the temperature in the condensed phase could reach 500-700 °C during a fire (e.g., in the cone calorimeter test, according to the heat flux [START_REF] Schartel | Temperature inside burning polymer specimens: pyrolysis zone and shielding[END_REF]). HRR curve (Figure 2(b)) is perfectly in accordance with TGA.

On the contrary, when EMA (series C) is filled with MDH, there is no important change of the degradation pathway. The degradation occurs in one main step with and without MDH (a small shoulder is observed at low temperature for EMA filled with MDH). Peak of mass loss rate is slightly shifted to higher temperature, probably due to endothermic effect related to the water release when MDH is decomposed (from 350 °C). Its value decreases slightly from 25 to 20%/min. A shoulder is also observed at lower temperature. No char is observed.

Similarly, the temperature of the peak of HRR is not significantly modified (Figure 3(b)). The intensity of the peak is lowered, probably as a result of the endothermic effect of water release, which slows down the pyrolysis of the polymer.

A significant decrease of the thermal stability is observed when MDH is added to LCP (series D). The incorporation of MDH leads to a first peak of mass loss rate at low temperature (450 °C) while pure LCP starts degrading at 480 °C with a peak of mass loss rate close to 540 °C (Figure 4(a)). Moreover, the final residue is lowered in the presence of MDH (43 and 37% for pristine LCP and LCP filled with MDH, respectively).

The incorporation of MDH into LCP also leads not only to a peak of HRR at lower temperature but also to a slight increase in Total Heat Release (from about 7 to 8 kJ/g -Figure 4(b)). This could be assigned to the decrease in char content due to MDH: A higher content of fuels is released in gas phase.

Figures 5 and6 summarize the results for all formulations to establish some relations between the chemical structure of polymers and the MDH influence on the thermal stability. Figure 5 plots the onset temperature of degradation for the filled polymers versus the pure ones. It has been previously explained that the onset temperature of degradation in TGA could be due to the mass loss of the polymer or alternatively of the MDH. Therefore, the onset temperature in PCFC was preferentially considered. This temperature corresponds to an arbitrary HRR of 30 W/g, which ensures that the polymer has started degrading at this temperature.

All polymers from series A, B, and D exhibit a lower onset temperature when they are filled with 10wt.% of MDH. This result tends to prove the role of the water released from MDH in promoting the hydrolytic degradation of the polymers at high temperature. Most of these polymers exhibit a decrease in onset temperature higher than 50 °C. PHB and PHBV are highly sensitive to hydrolysis, but the decrease in their onset temperatures is quite limited (10-25 °C). Pure PHB and PHBV degrade at low temperature (from 275 °C). Their stability is only slightly disturbed by the water release from MDH.

On the contrary, polymers from series C exhibit similar onset temperatures with and without MDH. These polymers seem to be not sensitive to hydrolysis.

Char promotion and its thermal stability are other important parameters. Figure 6 plots the char promotion versus its degradation temperature. Char promotion corresponds to the difference between the char content for polymers filled with MDH and for pure polymers. The mineral fraction of the residue (6.9wt.% due to MgO) is not considered to calculate the char content. For series A and B, the char percentage was calculated from the mass loss at high temperature. For series D, the char was calculated using the value of residual mass at 900 °C because this char is still thermally stable at this temperature.

Four groups (corresponding to the four series) could be identified. For series C, no char was formed in presence of MDH. For series D, the MDH is not a charring promoter but a charring inhibitor because the char content is lower in the presence of MDH for these polymers. MDH exhibits a significant char promotion effect for polymers from series A and B. The increase of char content is in the range 14-23wt.% for these polymers. However, a significant difference could be noticed from the temperature of the char degradation: The char is much more thermally stable for aromatic polyesters (series B) than for aliphatic polyesters (series A). Its degradation temperature is in the ranges 600-650 °C and 400-500 °C for series B and A, respectively. The thermal degradation has been also studied by Py-GC/MS analysis. The chromatograms obtained at a pyrolysis temperature of 900 °C for the four polymers representative of series A, B, C, and D are shown in Figure 7. PLA (from series A) and, to a less extent, PBT (from series B) exhibit a significantly different chromatogram when they are filled with MDH. On the contrary, chromatograms of EMA and LCP (from series C and D, respectively) are very similar with and without MDH.

According to the literature [START_REF] Aoyagi | Thermal degradation of poly[(R)-3-hydroxybutyrate], poly[ε-caprolactone], and poly[(S)-lactide][END_REF][START_REF] Kopinke | Mechanistic aspects of the thermal degradation of poly(lactic acid) and poly(βhydroxybutyric acid)[END_REF][START_REF] Kopinke | Thermal decomposition of biodegradable polyesters -II. Poly(lactic acid)[END_REF][START_REF] Omura | Thermal degradation behavior of poly(lactic acid) in a blend with polyethylene[END_REF], PLA pyrolysis is initiated by cis-elimination and transesterification. The main products obtained [START_REF] Kopinke | Thermal decomposition of biodegradable polyesters -II. Poly(lactic acid)[END_REF] are isomers of lactide (DL-lactide and meso-lactide) and volatile products such as CO, CO2, acetaldehyde, and acrylic acid. For a pyrolysis at 900 °C, the chromatogram of PLA shows essentially the formation of acrylic acid (peak at 2.9 min). In presence of MDH, the thermal degradation of PLA changes, and various new peaks are observed on the chromatogram. The presence of lactide (two peaks at 9.8 and 10.4 min corresponding to meso, D and L compounds), acrylic acid (peak at 2.9 min), and aromatic compounds (benzene and toluene at 3 and 4.2 min, respectively) can be noted.

Concerning the degradation of PBT, the appearance (at 2.7 min) and disappearance (at 13.4 and 13.7 min) of small peaks and evolution of intensities of some signals can be remarked when MDH was added. The most intense peaks correspond to aromatic compounds (derivatives of benzene and benzoic acid, etc.) and butadiene (peak at 1.7 min). The degradation of unfilled and filled PBT will be discussed in more detail later.

The chromatogram obtained for EMA corresponds to that described in the literature [START_REF] Cox | Pyrolysis-photoionization mass spectrometry of ethylene-methyl acrylate copolymers[END_REF]. Indeed, characteristic products of degradation of both parts of the copolymer (ethylene and methyl acrylate parts) were observed. The chromatogram shows the presence of methylacrylate and alkyl and alkene compounds. The presence of MDH does not affect the thermal degradation of EMA during the pyrolysis step of Py-GC/MS analysis.

The analysis of the LCP sample showed mainly the formation of phenol. Some small peaks attributed to aromatic compounds were also observed. Sato et al. [START_REF] Sato | Thermal degradation and combustion process of liquid crystalline polyesters studied by directly coupled thermal analysis-mass spectrometry[END_REF] studied the thermal degradation of the same LCP compound (Vectra A950), and they described that the main evolved gaseous products resulting from thermal degradation are H 2 O, CO, CO 2 , and phenol (phenol was released from the PHB moiety). In their study, other minor components were also detected (benzene, benzoic acid, naphthalene, etc.). As for EMA, the presence of MDH does not affect the thermal degradation of LCP during the pyrolysis step of Py-GC/MS analysis.

To conclude, it was highlighted that hydrolysis and char formation in presence of MDH as well as the char stability and the nature of the released gases are dependent on the polyester structure. It was shown that MDH can be a good char promoter for PBT and PET (aromatic polyesters from series B). Indeed, the hydrolysis of these polymers appears limited in comparison to aliphatic polyesters as PLA. The char amount is relatively high (20wt.% for PBT) and relatively thermally stable up to around 700 °C. It could be reasonably assumed that the char degradation temperature is higher and its protective effect is stronger. Therefore, in the following, the effect of MDH on the thermal stability and the flammability of PBT has been evaluated in details.

Effect of MDH on the thermal degradation and flammability of PBT

In the first part of this study, the MDH has been identified as a charring promoter for some polyesters, particularly for PBT. Formulations based on PBT filled with various MDH content were prepared and studied by TGA and PCFC to optimize the char promotion. Figure 8(a) shows the mass loss rate and HRR measured using TGA and PCFC for pristine PBT and PBT filled with MDH. As noticed previously, the degradation of PBT occurs in one single step (peak of heat release rate (pHRR) at 400 °C), whereas PBT filled MDH exhibits three steps: The first one (at 310 °C) is related to the early degradation of PBT due to its hydrolysis by water, the second step (at 400 °C) corresponds to the main degradation peak of PBT, and the third step (at 640 °C) corresponds to the degradation of the char. Another significant parameter in fire reaction is the effective heat of combustion (EHC). This parameter is calculated as the ratio between the total heat release (THR) measured in PCFC and the mass loss fraction measured in TGA. EHC was calculated for the whole degradation and separately for the first degradation steps (called step A and corresponding to the sum of the first step and the main degradation step in Figure 8(a)) and for the last degradation step (called step B and corresponding to the degradation of the char). The EHC values for pure PBT and PBT filled with 10wt.% of MDH are given in Table III. The EHC of step A for PBT filled with MDH is equal to 21.1 kJ/g. This value is very close to the value obtained for pure PBT. On the contrary, the EHC of step B for PBT filled MDH is much lower (12 kJ/g). Figure 8(b) shows also the change in the EHC values according to the MDH content, from 10 to 40wt.%. The EHC of step B does not change when increasing the MDH content. Therefore, the composition of the char could be assumed constant whatever the MDH content. Despite the scattering of results, it could be observed that the EHC of step A decreases slightly as a result of the dilution of fuel gases by water released during MDH decomposition. Figure 8(c) shows the char content according to the MDH content. The char content increases up to almost 30wt.% when increasing MDH content. Nevertheless, above a threshold value around 17wt.% of MDH, the char content levels off and starts decreasing at MDH content higher than 25wt.%. It means that the char promotion could be one of the main modes-ofaction of MDH at low or moderate MDH content. At higher MDH content, this effect is probably lesser in comparison with the other modes-of-action (dilution of condensed phase, endothermic effect, and dilution of fuel gases).

Thermal degradation and flammability of MDH-filled and MgO-filled PBTs

Magnesium dihydroxide not only promotes the charring of PBT but also leads to its early degradation by hydrolysis. If charring and hydrolysis are two decoupled phenomena, MgO could exhibit char promotion effect without hydrolysis. PBT formulations filled with 10wt.% of MDH or MgO were studied using TGA, PCFC, TGA-FTIR, and Py-GC/MS to compare both fillers as char promoters.

The thermal degradation of PBT starts earlier with MDH than with MgO, as shown in Figure 9. The onset temperature in PCFC is 370 and 300 °C for PBT and MgO-filled PBT and MDH-filled PBT, respectively. The peak at low temperature on thermogravimetric or HRR curves for PBT filled with MDH is not observed for PBT filled with MgO. Therefore, it could be assumed that there is no hydrolytic degradation of PBT in the presence of MgO. The main peak at 400-450 °C is significantly reduced in the presence of MDH and MgO, indicating a lower flammability. However, the decrease of pHRR is more pronounced with MDH probably because a part of the polymer is already decomposed at lower temperature in this case.

A peak of mass loss rate (and HRR) could be detected at high temperature for MDH-filled and MgO-filled PBTs. Therefore, MgO promotes charring as MDH, but the char content is significantly lower (10wt.% versus 17wt.%). However, because the particle size could affect the charring promotion, it should be noticed that both studied fillers exhibit very different particle size. The thermal stability of both chars is similar; their degradation temperatures are in the range 600-650 °C.

Magnesium dihydroxide and MgO are efficient char promoter for PBT. A stable char is formed, and it degrades at high temperature (in the range 600-650 °C) under non-oxidizing atmosphere. TGA in nitrogen flow of PBT filled with MDH and MgO was stopped at 550 °C, and the residue was kept to be further analyzed by FTIR spectroscopy. A significant difference at macroscopic scale is observed between both residues; it is intumescent only for PBT filled with MDH probably because of the water release from the hydrated filler.

As an example, Figure 10 shows the FTIR spectrum of the residue obtained from PBT-MgO pyrolysis (the FTIR spectrum is the same for PBT-MDH). Carboxylate functions (asymmetric and symmetric modes are located at 1580 and 1400 cm À1 , respectively) are identified. Two large absorption bands at 3300 and at 550 cm À1 could be attributed to water and MgO, respectively. Another band at 720 cm À1 is observed, but it was not attributed.

The EDX analyses indicate similar elementary percentages for both residues. Three main elements were found: C, O, and Mg, whose mass and molar percentages are reported in Table IV. Significant C percentage is observed, confirming the organo-mineral feature of the residue. The O percentage is too high to correspond only to the oxygen in MgO. The excess of O could be attributed to carboxylate groups, detected by FTIR analyses.

The thermal decomposition of PBT has been widely studied by TGA-FTIR [START_REF] Brehme | Phosphorus polyesteran alternative to low-molecular-weight flame retardants in poly(butylene terephthalate)?[END_REF][START_REF] Levchik | A review on thermal decomposition and combustion of thermoplastic polyesters[END_REF] or Py-GC/MS [START_REF] Ohtani | Analysis of thermal degradation of terephthalate polyesters by high-resolution pyrolysis-gas chromatography[END_REF][START_REF] Vijayakumar | Pyrolysis studies of aromatic polyesters[END_REF]. The degradation of PBT during pyrolysis is initiated by random scissions of the ester bonds [START_REF] Buxbaum | The degradation of poly(ethylene terephthalate)[END_REF]. The scission is achieved through a six-member cyclic transition state to obtain alkenyl and carboxyl end groups. The main product of the PBT degradation is butadiene, but tetrahydrofuran (THF) is also released. The scission of the C-O ester bond and intra-molecular or inter-molecular hydrogen transfer lead to the formation of terminal hydroxyl compounds such as 1,4-butanediol, which leads to THF through an acid-catalyzed dehydration reaction [START_REF] Vijayakumar | Pyrolysis studies of aromatic polyesters[END_REF]. The formation of water during the decomposition of the polymer contributes to the acid catalyzed hydrolysis of the ester groups and to the release of THF [START_REF] Lum | Thermal decomposition of poly(butylene terephthalate)[END_REF]. The decomposition starts from successive scissions of C-O alkyl or acyl groups to form products with vinyl, carboxylic, aldehyde, or hydroxyl functions. The products functionalized with carboxylic groups are converted at high temperature by CO 2 elimination, into more stable compounds [START_REF] Sugimura | Studies on thermal degradation of aromatic polyesters by pyrolysis-gas chromatography[END_REF] such as benzene or naphthalene. The Py-GC/MS (Figure 11) and the TGA/FTIR (Figure 12) analyses show significant differences between pure PBT, MDH-filled PBT, and MgO-filled PBT. The Py-GC/MS chromatograms at 900 °C for the three samples show the presence of butadiene, benzene, benzoic acid, naphthalene, and diphenyl (peaks 1, 3, 5, 6, and 9 on chromatograms, respectively). However, the ratios of peak heights are different for these different samples. This reflects a difference of degradation of these PBT composites during pyrolysis. The chromatogram of PBT filled with MDH is significantly different (with a decrease in peaks 1 and 5 and an increase of peaks 2 and 9), whereas the chromatograms of PBT and PBT filled with MgO exhibit similar patterns. This observation shows that water release from MDH decomposition influences the degradation pathway of PBT and changes the ratios of pyrolytic products.

It must be noticed that some other components were not identified (peaks 2, 4, 7, 8, 10, and 11). The TGA/FTIR analyses have also highlighted differences during the thermal decomposition of pure PBT and MDH-filled or MgO-filled PBT. The results are summarized in Figure 12. The FTIR spectra recorded at different temperatures show differences in terms of decomposition for the three samples.

The degradation of pure PBT, at 420 °C (Figure 12(a)) during the maximum of decomposition of PBT, releases butadiene, benzoic acid, butyl benzoate, CO 2 , and CO as main volatile decomposition products in accordance with the literature [START_REF] Brehme | Phosphorus polyesteran alternative to low-molecular-weight flame retardants in poly(butylene terephthalate)?[END_REF].

Magnesium dihydroxide or MgO modifies the PBT decomposition pathway below 420 °C. In comparison with pure PBT, an earlier release of CO 2 is observed from 240 °C for MDH-filled PBT and 330 °C for MgO-filled PBT. The presence of MDH leads to an early extensive hydrolytic degradation of PBT due to the water release from the filler. However, in both cases, we can observe a significant release of THF from 330 °C, which is more important for MDH-filled PBT. The THF formation is the consequence of the ring formation reaction from butanediol, which is formed by hydrolytic ester bond scission [START_REF] Levchik | A review on thermal decomposition and combustion of thermoplastic polyesters[END_REF][START_REF] Lum | Thermal decomposition of poly(butylene terephthalate)[END_REF]. Moreover, in the presence of basic Mg ions, the formation of terephthalate salts is preferred against the terephthalic acid formation. The hydrolytic bond scission is enhanced by the water release of MDH. Above 420 °C, the degradation pathway seems very close to pure PBT with the presence of benzoic acid, butyl benzoate, CO 2 , and CO except butadiene (this identification is not obvious).

We must emphasize that an important release of CO and CO 2 is still observable at 640 °C for MDHfilled and MgO-filled PBTs. This release at high temperature in the presence of MDH and MgO is due to the char decomposition. This char stable at high temperature results from the terephthalate salts formation.

The results obtained by Py-GC/MS and TGA/FTIR are complementary and allow to draw conclusions about the nature and the release temperature of pyrolysis gases. The release of water due to MDH led to change the decomposition mechanism of PBT with lower release temperatures of volatile compounds such as THF and CO 2 . Moreover, the addition of MgO unlike MDH does not seem to affect the nature of PBT decomposition gases, except that the degradation of char led to the release of CO 2 at higher temperature (CO 2 is still present at 640 °C for MDH-filled or MgO-filled PBT).

Cone calorimeter tests

The main objective of this article is to evaluate the char promotion of MDH into polyesters and especially into PBT mainly from TGA and PCFC results. It is well known that the flame retardancy could not be evaluated properly using microscale analyses. It is beyond the scope of this article to establish the relations between PCFC and other tests. Such analyses could be found elsewhere [START_REF] Lyon | Screening flame retardants for plastics using microscale combustion calorimetry[END_REF][START_REF] Sonnier | Combining cone calorimeter and PCFC to determine the mode of action of flame-retardant additives[END_REF].

Nevertheless, it is of great interest to check if the selected system PBT-MDH could exhibit significant flame retardancy in the cone calorimeter test. Pure PBT and PBT filled with 10 and 25wt.% of MDH were tested at an irradiance of 50 kW/m 2 (Figure 13). At these low or moderate MDH contents, main modes-of-action of MDH (as dilution of organic phase, endothermic effect, and dilution of fuels by water) are believed to be none or poorly efficient.

However, the incorporation of MDH leads to a significant decrease in pHRR (from about 1000 to less than 300 kW/m 2 with only 10wt.% of MDH, i.e., a decrease of 75%). The curve shape also changes: Pure PBT exhibits a typical thermally thin behavior, whereas filled PBT exhibits a thermally thick charring behavior confirming the barrier effect of the organo-mineral layer formed on the sample surface. PBT burns approximately in 300 s. On the contrary, filled PBT is not fully degraded after 1200 s. Intumescence and presence of char at the end of the test are observed for filled PBT. Finally, ignition occurs earlier when MDH is added into PBT (32 s for filled PBT vs. 40 s for pure PBT).

These observations (charring and early ignition due to hydrolysis) seem to be in good agreement with microscale analyses.

CONCLUSION

Magnesium dihydroxide was investigated as a char promoter in various polyesters using microscale analytical methods. Relationships between its presence and the chemical structure of the polyesters were studied. Aromatic groups in the polymer backbone allow significant charring with a high thermally stable char. Aliphatic polyesters also show charring, but the char exhibits poor thermal stability. The decomposition of polymers with pendant ester groups are not affected by the presence of MDH. Finally, MDH is detrimental for highly charring polymers as PC and LCP. A detailed study with PBT shows that char yield does not increase anymore from an optimal ratio of 25wt.% of MDH. Therefore, this proves that char promotion appears as one of the main modes-of-action at low MDH content, whereas other mechanisms of fire retardancy become more influent at high MDH content.

Magnesium oxide also allows charring while it avoids early degradation of PBT by hydrolysis. Nevertheless, the char content is significantly lower. It means that the char promotion effect could not be entirely decoupled to the hydrolysis. In other words, the release of water from MDH decomposition promotes not only hydrolysis but also charring through a mechanism that should be further investigated. Water also changes the pyrolysis products from PBT degradation to a certain extent. However, the structure of the residue appears to be the same for MDH and MgO.

Finally, it could be assumed that MDH or MgO are good flame retardants for aromatic polyesters as char promoter at low content, and synergy with other flame retardants should be investigated to achieve high rating levels in fire tests. Preliminary tests in cone calorimeter confirm that MDH reduces significantly the flammability of PBT even at low MDH content due to the charring promotion and the associated barrier effect.

Figure 1 .

 1 Figure 1. (a) Weight (%) and derivative weight (%/min) versus temperature with heating rate β = 10 °C/min obtained by thermogravimetric analyses under nitrogen atmosphere for PLA with and without 10wt.% of MDH. (b) HRR (W/g) versus temperature obtained by PCFC analyses with a heating rate of β = 1 °C/s under nitrogen atmosphere for PLA with and without 10wt.% of MDH.

Figure 3 .

 3 Figure 3. (a) Weight (%) and derivative weight (%/min) versus temperature with heating rate β = 10 °C/min obtained by thermogravimetric analyses under nitrogen atmosphere for EMA with and without 10wt% of MDH. (b) HRR (W/g) versus temperature obtained by PCFC analyses with a heating rate of β = 1 °C/s under nitrogen atmosphere for EMA with and without 10wt.% of MDH.

Figure 4 .

 4 Figure 4. (a) Weight (%) and derivative weight (%/min) versus temperature with heating rate β = 10 °C/min obtained by thermogravimetric analyses under nitrogen atmosphere for LCP with and without 10wt.% of MDH. (b) HRR (W/g) versus temperature obtained by PCFC analyses with a heating rate of β = 1 °C/s under nitrogen atmosphere for LCP with and without 10wt.% of MDH.

Figure 2 .

 2 Figure 2. (a) Weight (%) and derivative weight (%/min) versus temperature with heating rate β = 10 °C/min obtained by thermogravimetric analyses under nitrogen atmosphere for PBT with and without 10wt.% of MDH. (b) HRR (W/g) versus temperature obtained by PCFC analyses with a heating rate of β = 1 °C/s under nitrogen atmosphere for PBT with and without 10wt.% of MDH.

Figure 6 .

 6 Figure 6. Char promotion versus its degradation temperature.

Figure 5 .

 5 Figure 5. Onset temperature of degradation for MDH-filled polymers versus polymer degradation temperature for pure polymers (the shaded area corresponds to the temperature range of MDH decomposition).

Figure 7 .

 7 Figure 7. Py-GC/MS chromatograms of (a) PLA, PLA/MDH; (b) PBT, PBT/MDH; (c) EMA, EMA/MDH; and (d) LCP and LCP/MDH samples.

Figure 8 .

 8 Figure 8. PCFC and TGA curves for (a) pure PBT and PBT filled with 10wt.% of MDH, (b) EHC versus MDH content in PBT filled with MDH, and (c) char yield versus MDH content in PBT filled with MDH (c).

Figure 9 .Figure 11 .

 911 Figure 9. (a) Thermal stability in TGA (10 °C/min, 50-900 °C under N 2 ) and (b) flammability in (1 °C/min, 50-750 °C under N 2 , and 900 °C under O 2 ) for pure PBT, MDH-filled PBT, and MgO-filled PBT.

Figure 10 .

 10 Figure 10. FTIR spectrum of PBT-MgO residue.

Figure 12 .

 12 Figure 12. FTIR spectra obtained by TGA-FTIR at different temperatures (240, 330, 420, and 640 °C) for (a) pure PBT, (b) PBT/MDH, and (c) and PBT/MgO.

Table I .

 I Chemical structure and processing temperatures of studied polymers.

	Processing temperatures (PLA), polycaprolactone, polybutylene succinate, polyhydroxybutyrate (PHB) and PHB-co-valerate T dry = 60 °C T melt = 160 °C T dry = 40 °C T melt = 80 °C T dry = 60 °C T melt = 120 °C T dry = 60 °C T melt = 180 °C T dry = 60 °C T melt = 160 °C T dry = 120 °C T melt = 260 °C T dry = 100 °C T melt = 240 °C (Continues) T dry = 70 °C T melt = 220 °C T dry = T amb T melt = 160 °C T dry = T amb T melt = 160 °C T dry = 120 °C T melt = 300 °C T dry = 120 °C T melt = 260 °C (PHBV)), whereas series B corresponds to thermoplastic aromatic polyesters (polyethylene terephtalate (PET) and PBT). Series C includes thermoplastic polymers or copolymers with pendant ester groups (ethylene vinyl acetate, ethylene methyl acrylate (EMA), and polymethyl-methacrylate). Whereas the polymers from series A, B, and C hardly form char during pyrolysis, series D gathers polymers able to char easily, which present an ester group such as liquid crystal polyester (LCP) or a Processing temperatures carbonate group such as PC.
	The characteristics of MDH and MgO used in this study are shown in Table II.
		Suppliers and grades Suppliers and grades	Natureplast, PLE 005 Arkema, Altuglas V825T	Perstorp Caprolactone, Capa 6500 Exxon Mobil, Escorene Ultra UL00226CC	Xinfu Pharm, 1903F	Arkema, Lotryl 24MA005	Biocycle, PHB	Ticona, Vectra A950	Biocycle, PHBV		Bayer, Makrolon AL2647	Tergal, Reseko	Degussa, Vestodur 3000
		Chemical structure Chemical structure
		Name Name	Polylactique acid (PLA) Polymethyl-methacrylate (PMMA)	Polycaprolactone (PCL) Ethylene vinyl acetate (EVA)	Polybutylene succinate (PBS)	Ethylene methyl acrylate (EMA)	Polyhydroxybutyrate (PHB)	Liquid-crystal polyester (LCP)	Polyhydroxybutyrate-co-valerate	(PHBV)	Polycarbonate (PC)	Polyethylene terephtalate (PET)	Polybutylene terephtalate (PBT)
	Table I. (Continued)	Type Type	Series A: Aliphatic Series C: Pendant ester groups polyesters					Series D: Charring	polymers			Series B: Aromatic	polyesters

Table II .

 II Physicochemical properties of studied MDH and metal oxides.

	Name and stands	Suppliers and product references	Particle size (μm)
	Magnesium dihydroxide, Mg(OH) 2	Albemarle Corp., Magnifin H10	d 10 = 0.4
			d 50 = 0.8
			d 90 = 1.6
	Magnesium oxide, MgO	Prolabo, 52126	d 10 = 6.1
			d 50 = 110.2
			d 90 = 189.5

Table III .

 III Mass loss, total heat release, and effective heat of combustion for pure PBT and PBT filled with 10wt.% of MDH.

			Mass loss (%)	THR (kJ/g)	EHC = THR/ML (kJ/g)
	PBT		93	20.2	21.7
	PBT þ MDH	Step A Step B	65 17	13.7 2.1	21.1 12.4


Aurélie Taguet
email: aurelie.taguet@mines-ales.fr

Martin N Bureau

Michel A Huneault

Basil D Favis

Toughening mechanisms in interfacially modified HDPE/thermoplastic starch blends

Keywords:

The mechanical behavior of polymer blends containing 80 wt% of HDPE and 20 wt% of TPS and compatibilized with HDPE-g-MA grafted copolymer was investigated. Unmodified HDPE/TPS blends exhibit high fracture resistance, however, the interfacial modification of those blends by addition of HDPE-g-MA leads to a dramatic drop in fracture resistance. The compatibilization of HDPE/TPS blends increases the surface area of TPS particles by decreasing their size. It was postulated that the addition of HDPE-g-MA induces a reaction between maleic anhydride and hydroxyl groups of the glycerol leading to a decrease of the glycerol content in the TPS phase. This phenomenon increases the stiffness of the modified TPS particles and stiffer TPS particles leading to an important reduction in toughness and plastic deformation, as measured by the EWF method. It is shown that the main toughening mechanism in HDPE/TPS blends is shear-yielding. This article demonstrates that stiff, low diameter TPS particles reduce shear band formation and consequently decrease the resistance to crack propagation.

Introduction

Starch is a low cost, hydrophilic and biodegradable renewable resource with limited mechanical performance. Its ductility can be improved by the addition of a plasticizer such as glycerol to water-swollen starch granules in the presence of heat and high shear [START_REF] Averous | Biodegradable multiphase systems based on plasticized starch[END_REF][START_REF] Otey | Biodegradable films from starch and ethylene-acrylic acid copolymer[END_REF][START_REF] Swanson | Starch-plastic materials-preparation, physical properties, and biodegradability (a review of recent USDA research)[END_REF]. The resulting plasticized starch is known as thermoplastic starch (TPS). To decrease its hydrophilicity and improve its mechanical properties, TPS can be melt-blended with polyolefins, such as polyethylene (PE) [START_REF] Bikiaris | LDPE/plasticized starch blends containing PE-g-MA copolymer as compatibilizer[END_REF][START_REF] Bikiaris | LDPE/plasticized starch blends containing PE-g-MA copolymer as compatibilizer[END_REF][START_REF] Rodriguez-Gonzalez | High performance LDPE/thermoplastic starch blends: a sustainable alternative to pure polyethylene[END_REF][START_REF] St-Pierre | Processing and characterization of thermoplastic starch/polyethylene blends[END_REF][START_REF] Wang | Preparation and characterization of compatible thermoplastic starch/polyethylene blends[END_REF]. PE and TPS form immiscible blends with a high interfacial tension. Their interfacial interactions can be improved by blending with copolymer interfacial modifiers [START_REF] Cerclé | High performance polyethylene/thermoplastic starch blends through controlled emulsification phenomena[END_REF][START_REF] Prachayawarakorn | Properties of thermoplastic rice starch composites reinforced by cotton fiber or low-density polyethylene[END_REF][START_REF] Yin | Rheological and thermal behavior of starch/LDPE blends containing EAA[END_REF], the most common of which is PE grafted with maleic anhydride, PE-g-MA [START_REF] Bikiaris | LDPE/plasticized starch blends containing PE-g-MA copolymer as compatibilizer[END_REF][START_REF] Bikiaris | LDPE/plasticized starch blends containing PE-g-MA copolymer as compatibilizer[END_REF][START_REF] Sailaja | Use of maleic anhydride-grafted polyethylene as compatibilizer for polyethylene-starch blends: Effects on mechanical properties[END_REF][START_REF] Wang | Effects of glycerol and PE-g-MA on morphology, thermal and tensile properties of LDPE and rice starch blends[END_REF]. In a previous article on HDPE/TPS blends [START_REF] Taguet | Interface/morphology relationships in polymer blends with thermoplastic starch[END_REF], it was shown that the glycerol-rich domains of TPS dispersed phase migrated to the interface, and played a lubricant role, which resulted in very high mechanical properties. Moreover, upon interfacial modification of HDPE/TPS with HDPE-g-MA, maleic anhydride can react with glycerol, decreasing the glycerol content of the TPS particle and lead to micelle formation in the matrix. It was shown that a higher glycerol content leads to a higher amount of copolymer necessary to saturate the interface. It was also demonstrated that high tensile properties were obtained for modified HDPE/TPS blends due to a double mechanism at the interface, related to the formation of glycerol and copolymer layers. [START_REF] Sailaja | Effect of epoxy functionalized compatibilizer on the mechanical properties of low-density polyethylene/plasticized tapioca starch blends[END_REF] reported a similar improvement of the mechanical properties of LDPE with varying proportions of plasticized tapioca starch in presence of a compatibilizer (poly(ethylene-co-glycidyl methacrylate)). Micrographs of impact and tensile fractured surfaces showed a brittle-ductile failure transition with increasing amounts of copolymer at TPS contents higher than 30 wt%. Regarding the literature on polyolefin/TPS blends, a few articles dealing with the mechanical performances of such blends have been reported [START_REF] Gupta | Preparation and characterization of potato starch based low density polyethylene/low density polyethylene grafted maleic anhydride biodegradable polymer composite[END_REF][START_REF] Ning | The influence of citric acid on the properties of thermoplastic starch/linear low-density polyethylene blends[END_REF][START_REF] Tena-Salcido | Effect of morphology on the biodegradation of thermoplastic starch in LDPE/TPS blends[END_REF] which did not address failure mechanisms. Compatibilized polyolefin/TPS blends have been shown to be complex systems due to the presence of glycerol and its possible reactivity with the compatibilizer.

Fracture toughness measurement is an interesting tool to characterize the state of the interface in a polymer blend [START_REF] Ahn | Rubber toughening of nylon 6 nanocomposites[END_REF][START_REF] Bucknall | Notched impact behavior of polymer blends: Part 1: New model for particle size dependence[END_REF][START_REF] Laura | Effect of rubber particle size and rubber type on the mechanical properties of glass fiber reinforced, rubber-toughened nylon 6[END_REF][START_REF] Laurens | Adhesion promotion mechanisms at isotactic polypropylene/polyamide 6 interfaces: role of the copolymer architecture[END_REF]. The main toughening mechanisms in polymer blends are identified to be cavitation, crazing and/or shear-yielding [START_REF] Mai | Application of fracture mechanics for characterization of toughness of polymer blends[END_REF][START_REF] Wu | Study on morphology and toughening mechanisms in polymer blends by microscopic techniques[END_REF][START_REF] Yee | Toughening mechanisms in elastomer-modified epoxies. Part 1. Mechanical studies[END_REF]. While crazing is the main toughening mechanism of glassy polymers [START_REF] Brown | A molecular interpretation of the toughness of glassy polymers[END_REF], cavitation or void growth is mostly observed in rubber toughened polymers [START_REF] Bagheri | Role of particle cavitation in rubber-toughened epoxies: 1. Microvoid toughening[END_REF][START_REF] Bucknall | A model for particle cavitation in rubber-toughened plastics[END_REF]. Shearyielding is the process by which most ductile polymers extend to high strains in standard tests. Chain segments slip past each other in response to shear stresses, resulting in localized deformations in the yield zone while remaining close to constant volume [START_REF] Bowden | Formation of microshear band in polystyrene and poly(methylmethacrylate)[END_REF][START_REF] Brady | Yielding behavior of glassy amorphous polymers[END_REF][START_REF] Brady | Similarity between craze morphology and shear-band morphology in polystyrene[END_REF]. Other phenomena such as crack bifurcation and/or deflection by rubber particles [START_REF] Kaber | Crack deflection processes-I[END_REF][START_REF] Pearson | Influence of particle size and particle size distribution on toughening mechanisms in rubber-modified epoxies[END_REF], rubber-particle stretching and bridging [START_REF] Kunz-Douglass | A model for the toughness of epoxy-rubber particulate composites[END_REF][START_REF] Pearson | Influence of particle size and particle size distribution on toughening mechanisms in rubber-modified epoxies[END_REF] and formation of a plastic zone ahead of the crack [START_REF] Bascom | Effect of temperature on the adhesive fracture behavior of an elastomer-epoxy resin[END_REF][START_REF] Bascom | Fracture of epoxyand elastomer-modified epoxy polymers in bulk and as adhesives[END_REF][START_REF] Bascom | Fracture reliability of structural adhesives[END_REF][START_REF] Bascom | Adhesive fracture behavior of CTBN-modified epoxy polymers[END_REF] contribute to the toughening mechanism of polymer blends. In most polymer blends, these different toughening mechanisms occur simultaneously [START_REF] Brady | Similarity between craze morphology and shear-band morphology in polystyrene[END_REF][START_REF] Bucknall | A model for particle cavitation in rubber-toughened plastics[END_REF].

A tensile dilatometry test can be performed by measuring the volume strain as a means of identification of the deformation mechanism(s) taking place during the course of a usual stress-strain experiment [START_REF] Bucknall | Dilatometric studies of crazing in rubbertoughened plastics[END_REF]). Indeed, it is possible to differentiate deformation mechanisms that induce volume changes, such as crazing and cavitation, or not, such as shear-yielding. [START_REF] Bucknall | Dilatometric studies of crazing in rubbertoughened plastics[END_REF] were the first to introduce the tensile dilatometry approach to evaluate quantitatively the deformational mechanisms of rubber-toughened polystyrene. The rate of volume change was related directly to the rate of crazing. Indeed, during a tensile or a creep test, the slope of the volume strain, V/V 0 against longitudinal (also called axial) strain, ε a , curve equals zero for a pure shear-flow process, and equals unity for a pure crazing or cavitation process. The tensile dilatometry tests were performed to study the contribution of cavitation and crazing to the fracture resistance in many rubber-toughened plastics, such as elastomermodified epoxies [START_REF] Yee | Toughening mechanisms in elastomer-modified epoxies. Part 1. Mechanical studies[END_REF], PMMA and PP [START_REF] Naqui | Tensile dilatometric studies of deformation in polymeric materials and their composites[END_REF], rubber-modified nylon [START_REF] Borggreve | Impact behavior of nylon-rubber blends 6. Influence of structure on voiding processes: Toughening mechanism[END_REF], HDPE/nylon 6 [START_REF] Fellahi | Tensile dilatometry of injection-molded HDPE/PA6 blends[END_REF]. The characterization of the toughness of some polymer blends has been well-studied [START_REF] Mai | Application of fracture mechanics for characterization of toughness of polymer blends[END_REF][START_REF] Williams | Fracture mechanics of polymers[END_REF][START_REF] Wong | Fracture behavior of polymer blends[END_REF]). In the case of ductile fracture the well-known linear elastic fracture mechanics (LEFM) concept cannot be applied, because considerable energy is dissipated in the material creating extensive plasticity ahead of the crack tip. [START_REF] Kim | High strength, high fracture toughness fiber composites with interface control-a review[END_REF], [START_REF] Mai | Application of fracture mechanics for characterization of toughness of polymer blends[END_REF], Mai andCotterell (1985, 1986), [START_REF] Mai | The essential work of plane stress ductile fracture of linear polyethylenes[END_REF], [START_REF] Wu | The essential fracture work concept for toughness measurement of ductile polymers[END_REF] have developed a methodology based on [START_REF] Broberg | Crack-growth criteria and non-linear fracture mechanics[END_REF][START_REF] Broberg | On stable crack growth[END_REF] theory of fracture to characterize the fracture behavior of polymeric materials. According to this essential work of fracture (EWF), and under a plane-stress state, the total work of fracture generated during crack growth, W f , can be divided into two components: (i) the essential work of fracture (W e ), associated to the work spent in the inner fracture process zone where W e is a property of the material and, (ii) the non-essential work of fracture (W p ), dissipated in the outer plastic zone. W p is geometry dependent. The EWF method is a useful technique to measure the toughness in ductile materials as it separates the work dissipated in the plastic deformation zone from the work spent to create new surfaces under plane stress conditions. Most other methods require plane strain conditions. In 2010, an extensive review dealing with the application of the EWF concept was published [START_REF] Barany | Application of the essential work of fracture (EWF) concept for polymers, related blends and composites: A review[END_REF], listing all the articles on EWF applied to polymers, blends and composites. [START_REF] Chaléat | Properties of a plasticised starch blend. Part 1: Influence of moisture content on fracture properties[END_REF] evaluated the influence of the moisture content on the tensile properties of plasticized starch/high molecular weight polyol blends by using EWF. As the moisture content increases, the material becomes more ductile and EWF measurements highlight a strong dependence of the fracture behavior on the moisture content of this material.

In the present article, the evaluation of the toughness of HDPE/TPS and especially the identification of the mechanisms responsible for fracture resistance are addressed by investigating the resistance to crack initiation and/or propagation of the biobased TPS phase.

Experimental

Materials

The high density polyethylene (HDPE), supplied by Nova Chemicals (Calgary, Canada), was a Sclair ® HDPE2710 (MFI = 17). The native wheat starch (Supergell 1203-C) and the glycerol were obtained from ADM (Candiac, Canada) and Laboratoire MAT (Beauport, Canada), respectively. The wheat starch was composed of 25 wt% of amylose and 75 wt% of amylopectin. The glycerol was pure at 99.5 wt% and contained 0.5 wt% water. The HDPE copolymer grafted with maleic anhydride (HDPE-g-MA) was purchased from Dupont (Fusabond ® E MB-265D with a MFI = 12). The wt% MA content was calculated on a weight basis via elemental analysis and was equal to 3.9%.

Blend preparation

Starch granules were gelatinized, plasticized with glycerol and water and blended with HDPE and compatibilizer in a one-step extrusion process. Blends were prepared containing 20 wt% of TPS and 80 wt% of HDPE. The compatibilizer was added with the HDPE, at various rates. All compatibilizer concentrations are based on the TPS content.

The processing of the polyethylene/thermoplastic starch blends was based on a process developed previously (Favis, Rodriguez-Gonzalez, & Ramsay, 2005;Favis, Rodriguez, & Ramsay, 2003;[START_REF] Rodriguez-Gonzalez | High performance LDPE/thermoplastic starch blends: a sustainable alternative to pure polyethylene[END_REF]. The extrusion system was composed of a single-screw extruder (SSE) connected midway to a co-rotating twin-screw extruder (TSE). A starch/glycerol/water suspension (detailed information for this starch suspensions are given elsewhere [START_REF] Taguet | Interface/morphology relationships in polymer blends with thermoplastic starch[END_REF]) was fed in the first zone of the TSE. Native starch was gelatinized and plasticized and volatiles were extracted in the first part of the TSE. Molten HDPE and copolymer (T = 160 • C) were fed from the SSE to midway on the TSE. TPS, HDPE and copolymer were then mixed in the latter part of the TSE. The TSE screw speed was 150 rpm for all blends. A three-hole strand die (diameter 3 mm) was used and strands were water cooled, followed by air cooling and then pelletized. After plasticization of the starch, most of the water was removed during venting. Hence, the final TPS can be considered as a binary blend containing starch and glycerol. Under such conditions, the final glycerol contents of the TPS after extrusion were 24 wt%, 28 wt%, 36 wt% and 40 wt%, and are referred to as TPS24, TPS28, TPS36 and TPS40, respectively. The final blends are identified by the following nomenclature: for example, 20TPS36C9 stands for a blend containing 20 wt% of TPS36 and 9 wt% of HDPE-g-MA (based on TPS); indicating that this blend contains 78.2 wt% of HDPE, 20 wt% of TPS36 and 1.8 wt% of HDPE-g-MA.

Tensile properties and fracture mechanisms

After extrusion, samples were injection-molded into dogbone specimens of 57.00 mm in gage length, 9.70 mm in width and 3.05 mm in thickness. Tensile properties were investigated and the profile of the elongated specimen was studied to investigate the fracture mechanism using tensile dilatometry. Tensile tests were performed according to ASTM D638 with an Instron 4400R universal testing machine at a crosshead speed of 50 mm/min. Elastic modulus, elongation at break and strain at yield were measured from the obtained stress-strain curves. The test was performed at a crosshead speed of 5 mm/min and stopped just prior to necking and samples were cryofractured.

Scanning electron microscopy and image analysis

The cryofractured surfaces were observed under a Jeol SM 840 scanning electron microscope (SEM) operating at a voltage of 5 kV. Specimens were primarily coated with a thin layer of gold-platinum by physical vapor deposition. The method employed for particle size measurements is given elsewhere [START_REF] Taguet | Interface/morphology relationships in polymer blends with thermoplastic starch[END_REF].

Tensile dilatometry

Volume change measurements were carried out from dogbone specimens using an Instron universal mechanical tester 55R1123 equipped with a video-extensometer (Norwood, MA). A constant strain rate of 50 mm/min was used. Axial and transverse extensions were measured with a video extensometer. It is assumed, according to Poisson's ratio, that the deformation in the thickness is equal to that in the width. The volume strain V/V 0 is calculated using the formula:

V V 0 = (1 + ε a) (1 + ε T ) 2 -1 (1)
where V and V 0 are the change in volume and the original volume, respectively, and ε a and ε T represent the axial and the transverse strain, respectively. Note that ε T assumes a negative value.

d V/V 0 dε a = (1 + ε T ) 2 (2)
Hence, at the yield point a, d V/V 0 /dε a = 0 indicates that shear deformation is the only non-Hookean deformation mechanism. On the contrary, at the yield point a, d V/V 0 /dε a = 1 indicates that cavitation and/or crazing are the only mechanisms responsible for volume changes. V/V 0 was thus plotted vs. ε a for HDPE, unmodified and interfacially-modified 20TPS24, 20TPS28 and 20TPS36 blends. For each blend the wt% of copolymer corresponds to the critical concentration (C crit ) in emulsification curves. In all experiments the sample formed a neck and further interpretation of the V/V 0 vs. axial strain curve is no longer possible due to the non-homogeneous deformation along the specimen gage length.

Notched Izod impact tests

After extrusion, samples were injection molded with a Sumitomo SE50S injection machine under rectangular bars. The dimension of the specimens was 62.10 × 12.40 × 4.62 mm 3 . The notched samples were cut in the middle of the length. Samples were conditioned for 48 h at 23 • C under 50% of humidity. Then, notched specimen measurements were performed with a Resil 25 Izod impact tester from Ceast and according to ASTM D256 and ASTM 4812, respectively. Seven specimens were tested and their average value was reported.

Essential work of fracture

Fracture tests were performed by the EWF method using double edge-notched specimens of 80.45 mm in length, 30.60 mm in width and 3.20 mm in thickness. Specimens were tested with a 500 kN load cell, at a cross-head rate of 5 mm/min. At least fifteen specimens were tested in a ligament length ranging from 5 to 23 mm.

Under plane-stress conditions, the total work of fracture (W f ) is divided into an essential part (W e ) and a non-essential one (W p ), corresponding to the inner or fracture process zone and to the outer plastic deformation zone, respectively. Hence, W e represents all energy absorbed to create new surfaces, whereas W p gathers all energy dissipated outside of the fracture surface.

W f = W e + W p (3)
where W e is a function of the area l × t (W e = w e × l × t), W p is proportional to l 2 × t (W p = w p × l 2 × t), where l and t are the ligament length and thickness, respectively.

W f = W e + 1 × ˇwP (4)
where ˇis a shape factor related to the geometry of the plastic zone [START_REF] Mai | Application of fracture mechanics for characterization of toughness of polymer blends[END_REF]; Test Protocol for Essential Work of Fracture (Version 6)). Finally, w f is determined by calculating the integral under the load-displacement curve for every ligament length (l in mm). Therefore, w e and ˇwp are obtained by least-squares regression of w f vs. l; w e being the y-axis intercept and ˇwp , the slope of the line. EWF testing was only performed on plasticized 20TPS28 and 20TPS36 blends.

Results

The results of the particle size analysis and mechanical testing of interfacially-modified and unmodified blends containing TPS24, TPS28, TPS36 and TPS40 are presented in Table 1. Both interfaciallymodified and unmodified blends show a decrease in dispersed phase size when glycerol content is increased, pointing out the importance of the plasticization level. It was shown elsewhere [START_REF] Taguet | Interface/morphology relationships in polymer blends with thermoplastic starch[END_REF]) that 24% of glycerol in starch is not enough to effectively plasticize starch; to be efficient as a plasticizer, glycerol must be added with at least 28% in starch. Table 1 shows that the elastic modulus is increased when stiff TPS24 particles are added into a HDPE matrix, whereas the addition of TPS particles with higher starch content (TPS28, TPS36 and TPS40) leads to a leveling off or a drop of elastic modulus. A glycerol content of 28 and 36 wt% led to the highest elongation at break. The addition of a compatibilizer tends to increase the stress at yield ( max ) due to the improved interfacial adhesion between PE and TPS.

In light of these results, the tensile dilatometry tests were performed on HDPE, 20TPS24 and 20TPS28 and 20TPS36 blends in order to study their toughening mechanisms. A previous article [START_REF] Taguet | Interface/morphology relationships in polymer blends with thermoplastic starch[END_REF] studied the emulsification curves of all TPS blends as a function of the wt% of glycerol and reported the critical concentration (C crit ) that represents the amount of copolymer needed to saturate the interface. This value of C crit depends on the wt% of glycerol and is given in Table 1. For more clarity, the stress and volume change vs. axial strain curves are only presented for unmodified and interfacially-modified 20TPS36 blends in Figs. 1 and2, respectively. Table 2 summarizes the slope of the volume change vs. axial strain at the yield point obtained for Table 1 Particle size dn and tensile mechanical properties (Young's modulus E, strain at break ε b , stress at yield max and strain at yield εmax) for pure HDPE and blends containing 80% HDPE and 20% TPS (TPS24, TPS28, TPS36 and TPS40). Blends are unmodified or interfacially-modified with HDPE-g-MA.

wt%

20TPS36C0

Axial strain (%) pure HDPE and interfacially-modified and unmodified 20TPS24, 20TPS28 and 20TPS36 blends. The yield point is defined as the first maximum on the stress-strain curve, and depends on the material. Figs. 1 and2 (obtained for 20TPS36C0 and 20TPS36C9, respectively) exhibit a yield point at 12.2% and 11.6% of axial strain, respectively. For both interfacially-modified and unmodified 20TPS36 blends, the volume change is very low. Moreover, Table 2 shows that the slope of V/V 0 vs. ε a curves at the yield point is very low for wellplasticized 20TPS28 and 20TPS36 blends, ranging from 0 to 0.06. However, the slope increases to 0.20 and 0.25 for unmodified and interfacially-modified 20TPS24 blends, respectively. In general, the slope increases from unmodified to modified blends. The very low slopes obtained for 20TPS28 and 20TPS36 blends provide indications that their main toughening mechanism is shear-yielding. The results of tensile dilatometry also show that the interfacial modification with HDPE-g-MA copolymers tends to increase cavitation as a toughening mechanism. The results of the notched impact tests are presented in Fig. 3. The highest notched impact energy noted for unmodified blends was obtained for a glycerol content of 36 wt% (it increased from 24 wt% to 36 wt% then decreased at 40 wt%). Interfacial modification using grafted maleic anhydride led to an overall reduction in the notched impact strength for all blends except for 20TPS40. From these results, it appears that the notched impact strengths of the unmodified blends are governed by the plasticization level of the TPS particles. Low amounts of glycerol (24 wt%) lead to stiff TPS particles and high amount of glycerol (40 wt%) soften the materials leading to low notched impact strengths. In agreement with results presented elsewhere [START_REF] Taguet | Interface/morphology relationships in polymer blends with thermoplastic starch[END_REF], the results show that a glycerol content of 36 wt% appears to present a balance between a too low and a too high plasticization level. Upon interfacial modification, the surface area of TPS particles increases as a result of the lower TPS particle size. Since TPS is stiffer and more brittle than PE, even when plasticized, TPS particles act in the soft PE 2 Fig. 3. Notched Izod impact strength (in kJ/mm 2 ) of HDPE and unmodified and interfacially modified blends of compatibilized HDPE matrix (98.2% HDPE and 1.8% HDPE-g-MA) and 20% TPS with various amounts of glycerol ranging from 24% to 40%.

Table 3

Values obtained after EWF tests on pure HDPE and blends containing 80% HDPE, 20% TPS (TPS28 and TPS36) and HDPE-g-MA copolymer: the essential work (We), the non-essential work (ˇWp) of fracture and the total fracture work for a given ligament length (l = 17.8 mm) are reported.

Samples

We matrix as stress concentration sites. Failure in the notched impact specimens therefore occurs by crack propagation at the TPS-matrix interface. When more such crack propagating particles are present, the notched impact strength is reduced.

The toughness of the plasticized TPS/HDPE blends (i.e. containing TPS28 and TPS36 blends) was evaluated by means of the essential work of fracture (EWF) testing. The fracture properties of HDPE/TPS24 blend was not evaluated because of poor tensile properties. Prior to EWF data treatment, all load at yield vs. ligament length data were analyzed to make sure they were obtained under plane stress conditions. Only those ligament lengths that gave a load at yield data within the ±10% of the mean load at yield were selected. The load vs. extension curves were then examined for self-similarity. The values of W e and ˇWp were then calculated according to Eq. ( 4) from the linear regression of the specific work of fracture vs. ligament length (data are shown in Fig. 4 and Table 3). The total work of fracture (W f ) is dramatically reduced when TPS is added to HDPE, and further reduced when a HDPE-g-MA compatibilizer is added to the blends. For example, the value of W f is reduced by 72% and 55% when TPS28 and TPS36 are added to HDPE, respectively. The value of W f is further reduced by 90% and 84% when the copolymer is added to 20TPS28 and 20TPS36 blends, respectively. All the non-EWF values decrease upon TPS addition, but follow a different trend with respect to the glycerol content. For example, non-EWF values dropped by, respectively, 83% and 57% when 20 wt% of TPS28 and TPS36 are added to HDPE. Copolymer addition also led to a very large reduction in non-EWF values (91% and 86% less when copolymer is added to 20TPS28 and 20TPS36, respectively). The EWF and non-EWF results indicate that the introduction of more rigid TPS particles in a ductile HDPE matrix causes the reduction of the total work of fracture, and this reduction is more important for the blend containing lower glycerol content. They also indicate that compatibilization of this blend results in a further very significant reduction of the total work of fracture. When more particles are present in the matrix, as in the case of higher glycerol content or when the copolymer is added, the blends show a lower intrinsic fracture resistance, as indicated by the EWF values. The effect on plastic dissipation (non-EWF) is slightly different; since they act as fracture promoting sites, when a higher number of stiff particles are present in the matrix (e.g. at lower glycerol content (TPS28) or with the copolymer), the process zone extends less, resulting in a lower non-EWF. Indeed, stiff TPS particles reduce localized shear bands formation, effectively reducing overall plastic deformation [START_REF] Brady | Similarity between craze morphology and shear-band morphology in polystyrene[END_REF].

Fig. 5 represents micrographs of unmodified and interfacially modified 20TPS24 and 20TPS36 blends after elongation. The micrograph 1 of Fig. 5 clearly identifies void growth around TPS24 particles poorly bonded to the matrix after 55% of elongation. Moreover the TPS24 droplets are poorly elongated. Micrographs 2 and 4 in Fig. 5 represent compatibilized 20TPS24 and 20TPS36 blends after 30.8% and 54.2% elongation, respectively. Those micrographs exhibit no elongation of the TPS particles and void growth can be seen around TPS 24 particles. Finally, the micrograph 3 of Fig. 5, which represents the unmodified 20TPS36 blends after 59.2% of elongation, clearly identifies the high deformation of the TPS particles in the direction of the applied strain, with no cavitation around the particles. Those four micrographs in Fig. 5 confirm the results obtained in Table 2 (up to the yield point): the tensile deformation of well-plasticized HDPE/TPS blends is governed by a shear-yielding toughening mechanism, with a small increase of cavitation phenomenon after interfacial modification. But, if TPS particles are not sufficiently plasticized (TPS24), cavitation slightly increases.

Discussion

In the studied blends, 20 wt% of the HDPE is replaced by a plasticized wheat starch that is a renewable resource of very low cost. It was expected that the addition of an interfacial modifier, such as HDPE-g-MA copolymer would improve the adhesion between both polymer phases and the resulting mechanical properties, especially the fracture behavior (toughness). In reality, the fracture behavior observed, based on EWF and notched Izod impact results, shows that the maleic anhydride compatibilization drastically decreases both toughness and plastic deformation.

The introduction of TPS particles in an HDPE matrix and further compatibilization of the blend leads to important changes in their morphology and their mechanical properties. When the glycerol content is too high (40 wt%), it leads to an overall softening of the blend (cf. reduction in Young's modulus and the stress at yield in Table 1). On the contrary, if the plasticization level of the TPS dispersed phase is insufficient (as in the case of 20TPS24 blends), the TPS particles are too stiff and not prone to deformation.

The presence of these stiff particles in a soft matrix leads to a reduction in ductility and impact strengths (Table 1, Fig. 3). This is explained by the volume increase measured by tensile dilatometry (Table 2) indicating voiding, presumably at TPS particles, as observed on fracture surfaces (micrograph 1 in Fig. 5 showing voiding at TPS24 particles). For the well-plasticized blends (20TPS28 and 20TPS36), the addition of a grafted copolymer leads to particle size reduction (as shown in Table 1), but it also decreases

No elongation

After elongation 20TPS28 or 20TPS36 20TPS24 C0 C9 C0 C6

Fig. 6. Schematic representation of the deformation (matrix tearing and decohesion) during a tensile test for blends containing 20% of TPS24 modified with 0 and 6% of copolymer and for 20% of TPS28 or 36 modified with 0 and 9% of copolymer. Black space around the TPS particles characterizes void growth.

the glycerol content of the TPS particle, as a result of the reaction between the maleic anhydride groups of the copolymer and the hydroxyl groups of the glycerol, as reported elsewhere [START_REF] Taguet | Interface/morphology relationships in polymer blends with thermoplastic starch[END_REF]. This reaction results in a drastic increase of the stiffness of the TPS droplet. Not surprisingly, higher amount of stiff particles in an HDPE matrix reduces the toughness of the blend. Indeed, for all blends, tensile dilatometry tests demonstrate that the main toughening mechanism up to the yield point is shearyielding, but compatibilization is accompanied by voiding at the TPS particles (Table 2 and Micrographs 2 and 4 in Fig. 5). Fig. 6 summarizes the evolution of the microstructure of 20TPS24 and 20TPS28 (or 36) blends during a tensile test. This schematic representation of the modified or unmodified TPS particles after a tensile test indicate void growth around the particles in the case of 20TPS24 blends and compatibilized 20TPS28 or 36. In this case, TPS particles are stiff and have low deformability, which leads to voiding observed at TPS particles. This result can also explain the impact and EWF measurements. Whereas, for the unmodified 20TPS28 or 36 blends, the flow-induced orientation of the particles and the high deformability of TPS particle shows little cavitation and a dominant shear-yielding mechanism as shown in micrograph 3 in Fig. 5.

Notched impact strengths such as EWF results show the influence of the stiffness and the amount of particles per unit area on the intrinsic fracture resistance. During a fracture test, the crack propagates around the TPS particles. Hence, the higher the amount and the higher the stiffness of the particles, the lower the energy needed to propagate a crack. This explains the toughness decrease when TPS particles are added to HDPE matrix, and when this copolymer is added to the blend. A higher amount of TPS particles leads to the reduction of the process zone extension ahead of the crack tip, which reduces the plastic deformation and consequently the value of the non-EWF.

Conclusions

HDPE blended with thermoplastic starch can lead to materials with relatively high fracture resistance. The high toughness of HDPE can be maintained by maintaining a low amount of highly deformable TPS particles. That is to say, by adding well-plasticized non-modified TPS particles. Indeed, by decreasing the glycerol content of TPS particles, the addition of HDPE-g-MA dramatically increases the amount of stiff particles. Thus, even if compatibilization increase the adhesion between the phases [START_REF] Taguet | Interface/morphology relationships in polymer blends with thermoplastic starch[END_REF], it decreases the toughness and lead to the reduction of the plastic deformation zone ahead of the crack tip. This contributes to clearly decrease the fracture resistance. Finally, the unmodified blend containing 36 wt% of glycerol is the one with the highest resistance to fracture.

Fig. 1 .Fig. 2 .

 12 Fig. 1. Stress-axial strain-volume change curves of 20TPS36C0.

Fig. 4 .

 4 Fig. 4. Specific work of fracture vs. ligament length for HDPE and blends containing 20% of TPS28 and TPS 36 with 0% and 9% of HDPE-g-MA.

Fig. 5 .

 5 Fig. 5. SEM micrographs of: 20TPS24C0 (1), 20TPS24C6 (2), 20TPS36C0 (3) and 20TPS36C6 (4) blends, elongated at 55, 30.8, 59.2 and 54.2%, respectively, prior to cryofracture.

Table 2

 2 Slope of the curves plotting V/V0 vs. εa at the yield point (εmax) forHDPE and 20TPS24C0, 20TPS24C6, 20TPS28C0, 20TPS28C9, 20TPS36C0 and 20TPS36C9 blends.

	Blends	Slope of the V/V0 vs. εa curve at εmax
	HDPE	0.08
	20TPS24C0	0.20
	20TPS24C6	0.25
	20TPS28C0	0.02
	20TPS28C9	0.06
	20TPS36C0	0
	20TPS36C6	0.06

Present address: Chemical & Biotechnological Engineering, Université de Sherbrooke, Sherbrooke, QC, Canada J1K 2R1.

Acknowledgements

The authors gratefully acknowledge the financial support received from the Canadian Biomass Innovation Network and the Natural Sciences and Engineering Research Council of Canada through a Strategic Grant. Support from the National Research Council of Canada for experimental testing is also acknowledged.