Erlantz Lizundia
email: erlantz.liizundia@ehu.eus

Marie Cristina Penayo

Alain Guinault

Jose Luis Vilas

Sandra Domenek
email: sandra.domenek@agroparistech.fr

Impact of ZnO nanoparticle morphology on relaxation and transport properties of PLA nanocomposites

Keywords:

In this work we study the effect of zinc oxide (ZnO) nanoparticle morphology and concentration on the resulting relaxation and transport properties of polylactide (PLA) nanocomposites. Films containing spherical and rodshaped ZnO nanoparticles were incorporated into an amorphous polylactide (PLA) matrix through a solventprecipitation and compression moulding method. Morphological analyses carried out by scanning electron microscopy (SEM) together with ultraviolet-visible (UV-Vis) spectroscopy and thermogravimetric analysis (TGA). Results indicate a much better distribution of rod-shaped ZnO within PLA matrix. Relaxation experiments reveal faster physical aging kinetics of PLA in presence ZnO nanoparticles notwithstanding their shape, suggesting the presence of non-interacting surfaces between the amorphous PLA matrix and the ZnO nanoparticles. Interestingly, both helium and oxygen permeability remained stable or increase upon nanoparticle addition, and anisole sorption kinetics showed faster mass transport in the nanocomposites, suggesting that the low interfacial adhesion between PLA and ZnO brings supplementary voids to the material increasing mass transport. Overall, the experimental findings here reported provide a deeper understanding on the influence of metal oxide nanoparticle morphology on the resulting relaxation and gas transport properties of amorphous polymeric nanocomposites.

Introduction

One of the most promising biopolymers able to replace the petroleum-derived polymers for industrial applications is poly (lactide) (PLA). PLA is a linear aliphatic thermoplastic polyester derived from 100% renewable resources such as sugar, corn, potatoes, cane, beet, etc. [START_REF] Ducruet | Characteristics and applications of poly(lactic acid)[END_REF]. PLA possesses interesting physical properties together with biocompatibility and biodegradability character, which are all strongly influenced by its stereochemistry and molecular weight. Industrial PLA is mostly polymerized from L-lactide and D,L-lactide, where the Lisomer is the main fraction. PLA emerges as a potential candidate for consumer products such as packaging due to its transparency, degradability, low toxicity and environmentally benign characteristics. Nevertheless, there are some drawbacks, such as its high brittleness, poor crystallization behaviour and low gas barrier properties that limit its current use [START_REF] Lizundia | PLLA-grafted cellulose nanocrystals: role of the CNC content and grafting on the PLA bionanocomposite film properties[END_REF][START_REF] Armentano | Multifunctional nanostructured PLA materials for packaging and tissue engineering[END_REF].

In order to maintain food quality, food packaging materials should protect the product from the environment. When using PLA as a packaging material, this requires minimising both UV and gas permeability of the neat polymer, critical properties to increase the packagedfood shelf-life as oxygen and water vapour permeation through the film would cause oxidation of lipids and microbial growth respectively [START_REF] Lee | Permeation of gas and vapor[END_REF]. Usually, traditional polymers such as polyethylene terephthalate (PET), polyethylene (PE) and polypropylene (PP), are laminated to other materials (barrier films or aluminium foil), coextruded with barrier resins (PVOH, EVOH), coated or metalized to improve their barrier properties. This implies the use of a considerable amount of plastic, materials and adhesives, which can make the recycling processes difficult [START_REF] Ortenzi | Evaluation of crystallinity and gas barrier properties of films obtained from PLA nanocomposites synthesized via ˮin situˮ polymerization of L-lactide with silanemodified nanosilica and montmorillonite[END_REF]. In contrast, blending polymers with nanoscale particles have been shown as a cost-effective and easily scalable approach to tackle this matter, making the hosting PLA a competitive alternative to currently used petrochemical polymers [START_REF] Davis | Liquid water transport in polylactide Homo and graft copolymers[END_REF].

To improve the light barrier properties of PLA, we chose to add inorganic nanoparticles, namely zinc oxide (ZnO). ZnO is a multifunctional inorganic material with: i) an optical band gap in the UV region (3.37 eV) that makes it appropriate as UV-light absorber additive similarly to other inorganic materials [START_REF] Lizundia | Light and gas barrier properties of PLLA/metallic nanoparticles composite films[END_REF] ii) high thermal stability and proper antimicrobial properties [START_REF] Lizundia | Metal nanoparticles embedded in cellulose nanocrystal based films: material properties and post-use analysis[END_REF], iii) high photostability and has been approved by the American Food and Drug Administration (FDA) due to its low toxicity [START_REF] Trujillo | PLLA/ZnO nanocomposites: dynamic surfaces to harness cell differentiation[END_REF]. In this line, recent studies have highlighted the efficiency of ZnO inclusion into a PLA matrix in order to obtain films which improve the shelf-life of packaged fresh-cut products [START_REF] Li | Effects of PLA film incorporated with ZnO nanoparticle on the quality attributes of fresh-cut apple[END_REF], and effectively inactivate microbial growth when in contact with food [START_REF] Zhang | ZnO-PLA nanocomposite coated paper for antimicrobial packaging application[END_REF].

Whereas numerous aspects of the inorganic-polymeric nanocomposites have been extensively studied, ranging from their processing to the characterization of their mechanical and physical properties, the effect of nanoscale fillers on the so-called physical aging behaviour of glassy polymers remains relatively uncharted. This aging behaviour deals with the evolution of thermodynamic properties (enthalpy, volume, etc.) towards equilibrium, occurring below the glass transition temperature (T g ) [START_REF] Rittigstein | Model polymer nanocomposites provide an understanding of confinement effects in real nanocomposites[END_REF][START_REF] Boucher | Enthalpy recovery of PMMA/silica nanocomposites[END_REF]. Since physical aging may result in many deleterious effects ranging from embrittlement to reduction in permeability and dimensional instability, it has been extensively studied for bulk polymer systems in the past years [START_REF] Pluta | Plasticized polylactide/clay nanocomposites. II. The effect of aging on structure and properties in relation to the filler content and the nature of its organo-modification[END_REF][START_REF] Pluta | Polylactide compositions. The influence of ageing on the structure, thermal and viscoelastic properties of PLA/calcium sulfate composites[END_REF][START_REF] Lizundia | Crystallization, structural relaxation and thermal degradation in Poly(L-lactide)/cellulose nanocrystal renewable nanocomposites[END_REF][START_REF] Lizundia | Physical aging and mechanical performance of poly(L-lactide)/ZnO nanocomposites[END_REF]. In the case of polymer nanocomposites and, more generally, nanoscale confined materials such as polymer thin films or polymers in nanopores, the physical aging process can be dramatically modified in comparison to the same process in bulk polymers [START_REF] Cangialosi | Physical aging in polymers and polymer nanocomposites: recent results and open questions[END_REF][START_REF] Cangialosi | Effect of nanostructure on the thermal glass transition and physical aging in polymer materials[END_REF]. This is due to the fact that nanoscale confinement may induce two kinds of effects on polymer segmental mobility: i) molecular mobility in glass-forming systems is generally associated with a typical length scale and the presence of a physical barrier in confined materials may modify the size and the temperature dependence of such a length scale; ii) the surface effect is responsible for mobility changes at the interface: polymer-nanoparticle interface in nanocomposites; polymer-matrix in nanoporous materials; polymer-air or substrate in thin films. These peculiarities of physical aging in nanoscale confined polymers explain the emerging interest of the scientific community in the subject [START_REF] Cangialosi | Physical aging in polymers and polymer nanocomposites: recent results and open questions[END_REF].

In this work, a non-crystallisable polylactide has been selected to develop nanocomposites having both spherical (PLA/ZnO_s) and rodshaped ZnO nanoparticles (PLA/ZnO_r). Contrarily to the majority of the works reported so far, a non-crystallisable polymer has been selected as a model matrix in order to avoid the effects associated with the presence of crystalline phases on the resulting physical aging of the nanocomposite. The effect of nanoparticles on the glass transition, physical aging of the polymer and polymer relaxation during mass transport studies has been investigated. Physical aging has been evaluated through the study of mass transport across aged samples, using Helium, Oxygen and Anisole, an organic vapour able to swell PLA.

Materials and methods

Starting materials

PLA with a commercial name Ingeo™ 4060D was purchased from NatureWorks LLC. Rod-shaped ZnO nanoparticles (termed as ZnO_r) were purchased by L'Urederra technological centre (Spain), while spherical ZnO nanoparticles (termed as ZnO_s) were supplied by Plasmachem GmbH (Germany). Dichloromethane (CH 2 Cl 2 ), methanol (CH 4 O) and anisole were obtained from LabScan, Panreac and Sigma Aldrich (France) respectively.

Sample preparation

Samples have been prepared by solvent-precipitation followed by compression moulding. Nanoparticles (NPs) were homogeneously suspended in CH 2 Cl 2 via sonication (and they were added to previously dissolved PLA to yield nanocomposites containing 0, 0.2, 0.5, 1, 5 and 10 wt%. Another sonication step has been applied to dispersions before precipitating them in an excess of cold methanol. After vacuum drying (48 h at 60 °C), films were fabricated by compression moulding at 200 °C for 3 min under a pressure of 150 MPa (films were water quenched to allow their solidification). The films were stored in a desiccator over P 2 O 5 .

Morphological analysis

Transmission electron microscopy (TEM) analyses were carried out using a Philips CM120 Biofilter apparatus with STEM module at an acceleration voltage of 120 kV. Raw ZnO nanoparticles were analyzed depositing a droplet of nanoparticle suspension (0.1% (w/w)) in distilled water on a carbon-coated grid. The dispersion ZnO within the PLA matrix was evaluated using ultrathin sections of the nanocomposites (thickness of approx. 150 nm; samples cutted using a Leica Ultracut UCT ultramicrotome equipped with a diamond knife).

Thermogravimetric analysis

Thermal stability of nanocomposites was analyzed in a TGA MET-TLERTOLEDO 822e Thermal Gravimetric Analysis (TGA) instrument by heating the samples from room temperature to 500 °C at 10 °C/min under N 2 atmosphere.

Ultraviolet-visible (UV-Vis) spectroscopy

UV-Vis transmittance spectra were recorded with a Shimadzu MultiSpec-1501 spectrophotometer. Total transmittance experiments have been analyzed in the range of 190-800 nm with a sampling interval of 1 nm and 25 accumulations.

Differential Scanning Calorimetry analysis (DSC)

Thermal analyses of the samples were analyzed by means of the Differential Scanning Calorimetry analysis (DSC) (DSC-Q100 from TA Instruments) apparatus under 50 mL/min N 2 -flow. The samples (3-5 mg) were sealed in TZero aluminium pans. The DSC was temperature calibrated with Zn and In standards. The calibration of the heat-flow was carried out with In-standards and the calibration of the baseline with a sapphire disk. The samples were scanned by a heat-coolheat loop from 0 to 180 °C at 10 °C/min and the glass transition temperature (T g ) was determined as the mid-point of the heat capacity step at the second heating scan. The value of the glass transition temperature of PLA was used for the determination of the temperature program of physical aging with aging temperature equal to T a = T g -15 °C = 39 °C. Fig. S1 in the Supporting Information represents the applied aging program.

For the study of physical aging, all experiments began with a heating ramp until 180 °C in order to erase the material's previous thermal history. The samples were subsequently cool-heat loop from 0 to 150 °C at a scanning speed of 10 °C/min.

Physical aging times t a varied from 0 to 120 h in the DSC. Furthermore, a longer aging experiment was carried out in an external oven at 39 °C ( ≥ t weeks 2 a ). These samples were scanned in the DSC with the same temperature program (heat-cool-heat cycle from 0 to 150 °C at 10 °C/min). The enthalpy of recovery was calculated by integration of the recovery peak using Eq. ( 1): where C p a and C p u are respectively the specific heat of the aged and unaged samples normalized by the PLA content, and T x and T y are re- spectively temperatures below and above the calorimetric T g . For each polymer, T x and T y were chosen as the temperatures at which the un- aged and the aged sample scans superimposed below and above T g , respectively.

∫ = - H T t C T C T dT Δ ( , ) ( ( ) ( )) ,

Gas permeability testing

The oxygen permeability was measured using a Systec (France) permeability tester 8001 at 23 °C and 0% Relative humidity (RH) and a pressure gradient of 1 atm. The recorded oxygen transmission rate was normalized by the sample surface and thickness (measured at 10 points) for the calculation of the oxygen permeability.

Based on the same principle the helium permeability was measured by a home-made apparatus developed at CNAM with the help of a mass spectrometer. The measurements were performed at 23 °C and 0% RH on the sample that was used for measuring oxygen permeability.

Anisole sorption experiments

The gravimetric method was used for the measurement of the anisole sorption kinetics (see Table S1 for the physicochemical properties of anisole) [START_REF] Espino-Pérez | Designed cellulose nanocrystal surface properties for improving barrier properties in polylactide nanocomposites[END_REF]. The sorption curves were recorded at 40 °C and 0% RH, using an electronic microbalance operating at a fixed partial pressure (Intelligent Gravimetric Analyser 002, Hiden Isochema Ltd, Warrington, UK, (precision 0.2 μg). Films of ~15 mg and thickness ~100-300 μm were suspended from the microbalance by a stainless steel pan, which was contained in a thermo-regulated cell. The microbalance itself was maintained at 50 °C to ensure stability during the weight measurement and prevent solvent condensation. To ensure low water content, the samples were stored in a desiccator over P 2 O 5 before use.

After the samples had been loaded into the microbalance, they were purged for 24 h at 10 -5 mbar (10 -3 Pa) to remove all volatile compounds, which were sorbed in the film and present in the chamber. The anisole partial pressure, which corresponded to an activity of 0.2 (Table S1, Supporting Information), was set in the measuring chamber, which was connected to a tank with a vapour phase saturated in anisole. The partial pressure was regulated using a pressure transducer (Baratron ® , MKS Instruments, Wilmington, MA, USA). The mass uptake of the samples was recorded automatically over 5 days. The results of the preliminary experiments showed that there is some inevitable leakage past the copper seals, which, over time, leads to the dilution of the organic vapours by incoming air. Therefore, a pressure perturbation was introduced at regular time intervals to ensure a constant partial pressure while the measurement was being performed. This pressure perturbation forces the electronics of the microbalance to re-regulate the partial pressure by pumping part of the atmosphere in the chamber and re-regulating by connecting the organic vapour reservoir. It has a time constant, which is very different from the characteristic time of diffusion. Therefore, it was easily cancelled out in the curve analysis. In the following, the pressure perturbation is visible in the experimental data and corresponds to the spikes in the mass uptake curves.

The data treatment was performed using Crank's solution of Fick's second law for diffusion in a plane sheet [START_REF] Crank | The Mathematics of Diffusion[END_REF]. At short times and assuming constant surface concentrations, the solution becomes:

∑ = ⋅ ⎡ ⎣ ⎢ + ⎤ ⎦ ⎥ = ∞ - m m Dt L π n erf nL Dt 4 1 2 ( ) 2 , t eq n n 2 0 1 (2)
with m t being the relative mass uptake at time t, m eq being the relative mass uptake at Fickian equilibrium, L being the half sample thickness, (3)

The diffusion coefficient can be obtained from the slope to the curve

⎜ = ⎛ ⎝ f m m t L t eq
). In the case of thin samples, non-Fickian diffusion of anisole in PLA can be observed [START_REF] Espino-Pérez | Designed cellulose nanocrystal surface properties for improving barrier properties in polylactide nanocomposites[END_REF]. In this case, the equilibrium mass uptake was assumed to correspond to the change of slope of the

⎜ = ⎛ ⎝ f m m t L t eq
) curve. This assumption was validated with Espino's methodology [START_REF] Espino-Pérez | Designed cellulose nanocrystal surface properties for improving barrier properties in polylactide nanocomposites[END_REF]. Because of the high sample thicknesses of the PLA/ZnO nanocomposites, the equilibrium sorption could be not reached within the experimentally feasible time. In that case, the uncertainty on the determination of the diffusion coefficient was high. For that, the transport rate was discussed with the help of the normalized slope of the experimental curve.

Results and discussion

Nanoparticle dispersion

It is well-established that the macroscopic properties of nanocomposites closely depend on the nanoparticle dispersion within the hosting matrix [START_REF] Lizundia | Physical aging in poly(L-lactide) and its multi-wall carbon nanotube nanocomposites[END_REF]. To that end, transmission electron microscopy (TEM) has been applied to both PLA/ZnO_r and PLA/ZnO_s nanocomposites. TEM has firstly carried out with raw ZnO nanoparticles to deter, mine their size and morphology. TEM micrographs in Fig. 1a for spherical nanoparticles (ZnO-s) and in Fig. 1d for rod-shaped nanoparticles (ZnO-r) confirm the nanoscale nature of used fillers, which dimensions are 43 ± 24 nm in lenght and ~20 nm in width for ZnO-r and 25 nm in diameter for ZnO-s. Micrographs of nanocomposites reveal a non-oriented isotropic distribution of both spherical and rodshaped ZnO nanoparticles (see Fig. 1b andc and Fig. 1e and f respectively) over the entire polymeric matrix. While nanocomposites containing spherical nanoparticles present entanglements between adjacent NPs, ZnO_r remain randomly distributed within the PLA. As observed in Fig. 1b-c, when spherical nanoparticle concentration increases from 1 to 10 wt% the entanglements become more evident, resulting in large aggregated structures. On the contrary, rod-shaped nanoparticles could be found over the entire composite for all the studied composition range (Fig. 1e-f). In the light of these results, it could be expected that, for a given concentration, barrier properties would be improved to a larger extent in PLA/ZnO_r nanocomposites because PLA/ZnO_s showed a microcomposite structure rather than nanocomposite morphology, resulting less effective for the formation of a tortuous pathway [START_REF] Rezakazemi | State-of-the-art membrane based CO2 separation using mixed matrix membranes (MMMs): an overview on current status and future directions[END_REF]. More precisely, the achieved zinc oxide surface area exposed to PLA chains (and thus to the gas permeating through the material) is much larger in the case of PLA reinforced by rod-shaped ZnO.

Thermal stability of PLA/ZnO_r and PLA/ZnO_s nanocomposites

Thermogravimetric analysis (TGA) has been carried out to evaluate the thermal stability of PLA/ZnO_r and PLA/ZnO_s nanocomposites (see Table 1 for the characteristic thermodegradation temperatures and maximum degradation rates). As shown in Fig. 2, where the thermogravimetric traces are shown in the left and weight loss rates in the right, one main loss step characterizes the thermal degradation of PLA nanocomposites. In fact, degradation of neat PLA begins at 330.8 °C (determined by the first 5 wt% weight loss) and reaches its maximum weight loss rate at 366.2 °C. When ZnO is added, the peak thermal degradation temperature decreases up to 276.8 °C for the PLA/ZnO_r 10 wt% composition.

The marked decrease in thermal stability upon ZnO loading arises from the fact that, as other metal such as Al, Sn and Al, ZnO catalyse PLA depolymerisation generating unzipping depolymerisation and intermolecular transesterification reactions upon heating [START_REF] Lizundia | Towards the development of ecofriendly disposable polymers: ZnO-initiated thermal and hydrolytic degradation in poly(L-lactide)/ZnO nanocomposites[END_REF]. It should be noted that, especially for concentrations up to 0.5 wt%, the catalyzing effect of zinc oxide nanoparticles is much more marked in the case of nanocomposite containing rod-shaped NPs (T 5% , the temperature value of the first 5% weight loss, decreases by 73.2 °C in comparison with 53.2 °C for PLA/ZnO_r and PLA/ZnO_s containing 0.2 wt% respectively). Since the catalytic reactions take place at the ZnO interfaces, the amount of ZnO-PLA interfaces would determine the thermal degradation behaviour of nanocomposites. Therefore, these experimental findings suggest that ZnO_r nanoparticles are better dispersed through the polymer than ZnO_s and larger concentrations than 0.5 wt% yield similar dispersion degrees for both systems. Finally, the residual mass steeply increases upon loading suggesting that ZnO nanoparticles are not degraded in the studied temperature range.

UV-Vis spectroscopy

UV-Vis transmittance spectroscopy has been carried out to evaluate the UV-shielding behaviour of fabricated nanocomposites. As depicted in Fig. 3, while neat PLA presents high transparency over the 250-800 nm region, the transmittance of nanocomposites, which accounts for the relative number of photons that pass through the film under an incident light, continuously decrease upon ZnO loading.

Table 2 summarizes the transmitted light in both UV and visible regions, where remarkable differences could be observed between both types of nanoparticles, especially in the UV region. Unlike ZnO_r nanoparticles, which show a strong exciton absorption band at 363 nm, ZnO_s does not show any characteristic exciton absorption in the studied range. Accordingly, the addition of ZnO_r reduces the amount of transmitted ultraviolet light from 94.8% for PLA to 21.5% for its PLA/ ZnO_r 10 wt% nanocomposite. This may be ascribed to the fact that ZnO nanoparticles absorb UV light because of their electronic structure characterized by an empty conduction band and a filled valence band. It is shown that spherical zinc oxide does not substantially block UV light, which may be due to aggregation of NPs. Moreover, the transparency of nanocomposites which, according to the ASTM D1746-03 standard, is defined as the amount of transmitted light in the 540-560 nm range, drops from 96% for neat PLA to 73.3 and 70.9% for nanocomposites containing 10 wt% of ZnO_r and ZnO_s, respectively. The lower decrease in transparency for the PLA/ZnO_r system could be explained by the fact that, according to the Rayleigh scattering, upon particle aggregation a transparency loss and film whitening occurs. This is in agreement with TEM micrographs shown in Fig. 2, where the dispersion of ZnO_r within PLA was more homogeneous than that using ZnO_s.

Glass transition and physical aging of PLA, PLA/ZnO_r and PLA/ZnO_s nanocomposites

The impact of the two types of nanoparticles on the polymer relaxation phenomena was probed by analysing the glass transition and the physical aging kinetics. The main results are presented in Table 3 and the DSC curves are shown in the Supporting Information (Fig. S2 to Fig. S5). The glass transition temperature (T g ) of the PLA/ZnO_r nanocomposites was slightly inferior to that of PLA. This can be explained by a probable degradation of the PLA chains during film preparation, although the material was thoroughly dried before hot-moulding of the films. Indeed, ZnO particles are known to be able to catalyse the degradation reaction of PLA [START_REF] Lizundia | Towards the development of ecofriendly disposable polymers: ZnO-initiated thermal and hydrolytic degradation in poly(L-lactide)/ZnO nanocomposites[END_REF][START_REF] Murariu | High-Performance polylactide/ZnO nanocomposites designed for films and fibers with special end-use properties[END_REF]. Some works have reported on the use of surface treatments such as silanization to protect the polymer during melt processing. Pantani et al. [START_REF] Pantani | PLA-ZnO nanocomposite films: water vapor barrier properties and specific end-use characteristics[END_REF] and Murariu et al. [START_REF] Murariu | Current progress in the production of PLA-ZnO nanocomposites: beneficial effects of chain extender addition on key properties[END_REF] did not observe a shift to T g in PLA/ZnO nanorods composites produced by melt compounding using surface compatibilisation, while Anzlovar et al. [START_REF] Anzlovar | Degradation of PLA/ZnO and PHBV/ZnO composites prepared by melt processing[END_REF] reported that the average molar mass of PLA in PLA/ZnO nanocomposites obtained by melt extrusion was almost halved. They observed a T g increase of 36 °C upon inclusion of 1% ZnO nanoparticles by injection moulding, which was a much larger shift than that observed in the present work. The inclusion of PLA/ZnO_s nanocomposites at low concentrations had no impact on the T g (Table 3). Starting with a concentration of 1 wt%, the T g shifted to lower temperatures. The behaviour is consistent with the data of the thermal stability of the nanocomposites. Fig. 2 shows that nanocomposites of ZnO_s 0.2 and 0.5 wt% are slightly more stable than the other materials. Furthermore, the height of the heat capacity step at the glass transition of all materials was constant (Table 3, column ΔC p ). The width of the glass transition is a measure of local dynamic heterogeneity in a polymer and extensive literature exists investigating the impact of nanocharges on the local mobility of macromolecules [START_REF] Jancar | Current issues in research on structure-property relationships in polymer nanocomposites[END_REF]. Table 3 shows that the width of the glass transition increased in the case of PLA/ZnO_r 5 and 10 wt%, but only in the case of PLA/ZnO_s 10 wt%. The presence of the ZnO nanoparticles in the material could therefore impact the local chain mobility. The ZnO_s particles had higher tendency to aggregate, which might explain that the effect is only observed at the highest concentration [START_REF] Yousefzade | Cooperative rearranging region and dynamical heterogeneity of nanocomposites in poly(l-lactide) and functionalized carbon nanotubes systems[END_REF], as previously showed for the inclusion of carbon nanotubes into PLA, which resulted in an increase of dynamic heterogeneity.

In the aim to investigate further the impact of the presence of the ZnO nanoparticles on the non-equilibrium structural relaxation of PLA, a study of the physical aging was performed. The total recovery of enthalpy ( ∞ H Δ ) was estimated after the relation

= - ∞ H C T T Δ Δ ( ). p g a ( 4
)
The results are given in Table 3 and show the impact of the degree of undercooling on the estimated value. At equal degree of undercooling, the presence of ZnO particles had no impact, being a result of the absence of change in C Δ p . The increase of the recovered enthalpy with increasing aging times of the PLA/ZnO nanocomposites is shown in Fig. 4. The quantity of recovered enthalpy at the longest aging times was near to the estimated total enthalpy loss, which showed that the theoretical thermodynamic equilibrium (within the measurement uncertainty) was reached within the experimental time. The H Δ values of PLA reached after 1000 min aging are comparable to already published data [START_REF] Delpouve | Evidence of two mobile amorphous phases in semicrystalline polylactide observed from calorimetric investigations[END_REF][START_REF] Dobircau | Molecular mobility and physical ageing of plasticized poly (lactide)[END_REF][START_REF] Monnier | Reduced physical aging rates of polylactide in polystyrene/polylactide multilayer films from fast scanning calorimetry[END_REF]. The curves in Fig. 4a-b shows hardly any difference among the PLA/ZnO_r nanocomposites with different particle concentrations. All nanocomposites curves are having higher recovery of enthalpy than PLA at a given aging time. Dispersion of experimental data was higher in the case of ZnO_s nanocomposites (Fig. 4b). For quantification purposes, the relaxation rate was calculated using the data in Fig. 4 after the following equation [START_REF] Pan | Enthalpy relaxation and embrittlement of Poly(L-lactide) during physical aging[END_REF]:

= ∂ ∂ β δ t [ / (log )] , H H a q q T , , a 1 2 (5)
where β H is the relaxation rate in J . g -1 per decade, δ H the enthalpy loss in J . g -1 , t a the aging time, q 1 the cooling rate in °C. min -1 , q 2 the heating rate in °C. min -1 and T a the aging temperature. The rates q 1 and q 2 were fixed at 10 °C. min -1 . The results are listed in Table 3. The observed relaxation rate of the neat PLA was faster than the value published by Pan et al. [START_REF] Pan | Enthalpy relaxation and embrittlement of Poly(L-lactide) during physical aging[END_REF] (1.77 J g -1 per decade) at similar undercooling (18 °C), and slower than the results reported by Lizundia et al. [START_REF] Lizundia | Physical aging and mechanical performance of poly(L-lactide)/ZnO nanocomposites[END_REF], who observed a quite high relaxation rate (3.2 J g -1 per decade) at 16 °C undercooling (in the latter a different experimental protocol for performing the sample aging ex situ in an oven was done). Average molecular mass differences and experimental uncertainty could be a reason for the observed discrepancies, while differences in the PLA stereochemistry seem to be less important [START_REF] Aou | Roles of conformational and configurational defects on the physical aging of amorphous poly(lactic acid)[END_REF]. For example, the relaxation rate of polystyrene decreased with increasing molecular weight [START_REF] Málek | Volume and enthalpy relaxation rate in glassy materials[END_REF]. Because of the differences between the degree of undercooling and a decrease in molecular weight of PLA in the ZnO nanocomposites, a comparison is not evident. The higher undercooling would cause faster relaxation rates of PLA, while higher molecular weight would slower relaxations. Slower aging kinetics of PLA/nanoclay composites were reported by Ref. [START_REF] Maspoch | Ductile-brittle transition behaviour of PLA/o-MMT films during the physical aging process[END_REF]. On the other side, Lizundia et al. [START_REF] Lizundia | Physical aging and mechanical performance of poly(L-lactide)/ZnO nanocomposites[END_REF] showed with a different experimental protocol that the inclusion of up to 2 w% ZnO_r decelerated the relaxation rate (observed values between 0.9 and 1.25 J/g), and at 5 w% ZnO_r the relaxation rate was almost equal to that of neat PLA. In the present case, the initial relaxation rate of the ZnO_r nanocomposites was accelerated compared to PLA and the increase of the ZnO_r concentration decreased the relaxation rate. The presence of small concentrations of ZnO_s in PLA had no impact on the relaxation rate, only with 5 and 10 wt% a deceleration was found (Table 3).

There is a large discussion on the impact of nanoparticles on the physical aging rates of polymers and results in the literature are contradictory [START_REF] Cangialosi | Physical aging in polymers and polymer nanocomposites: recent results and open questions[END_REF]. For further insight, we investigated the recovery function, calculated from:

= - ∞ ∞ φ t H H t H ( ) Δ Δ ( ) Δ a a (6)
The recovery functions of the different samples are shown in Fig. 5. The inspection of the curves gives supplementary information to the relaxation rate, i.e. the speed of the aging process characterized by the time to reach equilibrium. Fig. 5a-b shows that the nanocomposites reach the equilibrium time at shorter times compared to PLA.

The counterintuitive enhancement of the physical aging upon ZnO_r and ZnO_s incorporation may be explained in terms on non-attractive interactions between the hosting amorphous PLA phase and the surface of the non-compatibilized inorganic nanoparticles. A similar effect was firstly reported by Boucher et al. [START_REF] Boucher | Enthalpy recovery of PMMA/silica nanocomposites[END_REF][START_REF] Boucher | Accelerated physical aging in PMMA/silica nanocomposites[END_REF], who ascribed the accelerated physical aging rates in poly (methyl methacrylate)/silica nanocomposites to a mismatch between the hosting polymer and the reinforcing phase, facilitating the diffusion of free volume holes to the interfaces polymer/nanoparticle interfaces. This effect has been found to occur also in other polymer nanocomposite systems such as polycarbonate nanocomposites containing grafted nanosilica particles [START_REF] Ramakrishnan | Physical aging in polycarbonate nanocomposites containing grafted nanosilica particles: a comparison between enthalpy and yield stress evolution[END_REF]. As the diffusion of the free volume is governed by the amount of available ZnO surfaces [START_REF] Boucher | Time dependence of the segmental relaxation time of poly(vinyl acetate)-silica nanocomposites[END_REF], the occurring nanoparticle aggregation as concentration increased lowers the effect of ZnO on the overall aging kinetics. It was possible to superpose the relaxation functions of the nanocomposites and the PLA on a single master curve by a horizontal shift as proposed by Boucher et al. [START_REF] Boucher | Enthalpy recovery of PMMA/silica nanocomposites[END_REF]. Fig. 5c andd show the obtained master curves. This behaviour suggested, that the underlying macromolecular relaxation phenomena of PLA was not impacted by the presence of the nanoparticles notwithstanding their shape factor.

In any case, those reported results here for amorphous PLA/ZnO nanocomposites follow opposite trend reported by Lizundia et al. [START_REF] Lizundia | Physical aging in poly(L-lactide) and its multi-wall carbon nanotube nanocomposites[END_REF] for the crystalline PLLA/ZnO system, suggesting that the presence of crystalline phases needs to be taken into account when analysing physical aging in polymers.

Transport properties of PLA/ZnO nanocomposites

The mass transport properties of PLA/ZnO nanocomposites were investigated by the analysis of helium (P(He)) and oxygen permeability (P(O 2 )) (see Table S2 for the corresponding permeabilities). Helium is a small, non-interacting molecule with polymers, sensitive on tortuosity effects and having the advantage of needing much shorter experimental times compared to oxygen. The P(He) and P(O 2 ) of PLA amounted to 58 ± 5.10 -18 (m 3 m m -2 . s -1 . Pa -1 ) and 2.6 ± 0.2 10 -18 (m 3 m m -2 . s -1 . Pa -1 ), respectively. Those values were in the range of already published data [START_REF] Courgneau | Effect of crystallization on barrier properties of formulated polylactide[END_REF][START_REF] Guinault | Impact of crystallinity of poly (lactide) on helium and oxygen barrier properties[END_REF][START_REF] Nassar | Multi-scale analysis of the impact of polylactide morphology on gas barrier properties[END_REF]. The low compatibility of PLA and nonmodified ZnO had as a consequence that the samples were mechanically very brittle. As our group has previously shown, the elongation at break decreases from nearly 4.6% for neat PLA to 1.2% for its 5 wt% nanocomposite [START_REF] Lizundia | Physical aging and mechanical performance of poly(L-lactide)/ZnO nanocomposites[END_REF][START_REF] Lizundia | Towards the development of ecofriendly disposable polymers: ZnO-initiated thermal and hydrolytic degradation in poly(L-lactide)/ZnO nanocomposites[END_REF]. This was expected from other reports on noncompatibilized ZnO nanocomposites done by solvent casting [START_REF] Murariu | High-Performance polylactide/ZnO nanocomposites designed for films and fibers with special end-use properties[END_REF][START_REF] Shankar | Incorporation of zinc oxide nanoparticles improved the mechanical, water vapor barrier, UV-light barrier, and antibacterial properties of PLA-based nanocomposite films[END_REF][START_REF] Restrepo | Development and characterization of polyvinyl alcohol stabilized polylactic acid/ZnO nanocomposites[END_REF]. The P(He) could be obtained for samples up to 1 wt% ZnO (for [SI] = 10 -18 (m 3 m m -2 . s -1 . Pa -1 ): ZnO_r 0.5 wt % 74 ± 7, ZnO_r 1 wt% 67 ± 7, ZnO_s 0.5 wt % 69 ± 7, ZnO_s 1 wt % 62 ± 7 (see Table S2 in the Supporting Information). At higher concentrations, the films broke in the measurement device. With regard to the experimental uncertainty, no significant difference among the PLA/ZnO nanocomposites was observed. In the case of P(O 2 ), the measurement times were longer, typically one day compared to 1 h for P(He). Therefore, most of the samples broke during the measurement. In case of successful measurement, the P(O 2 ) of ZnO_r 1 wt% amounted to 2.7 ± 0.2 (see Supporting Information). This confirmed the prediction of P(He) that the distribution of ZnO did not allow for increasing barrier properties. In the case of melt-extruded PLA/ZnO nanoparticles (rod shape), a decrease of about 18% in P(O 2 ) was reported for a nanocomposite having 5 wt% nanoparticles [START_REF] Marra | Polylactic acid/zinc oxide biocomposite films for food packaging application[END_REF]. Studies on water vapour permeability of nanocomposites of PLA and rod-shaped ZnO nanoparticles showed either very small or no gains [START_REF] Pantani | PLA-ZnO nanocomposite films: water vapor barrier properties and specific end-use characteristics[END_REF][START_REF] Shankar | Incorporation of zinc oxide nanoparticles improved the mechanical, water vapor barrier, UV-light barrier, and antibacterial properties of PLA-based nanocomposite films[END_REF]. This result can be expected, because nanospheres or rods have a small shape factor and they do therefore not induce much tortuosity to the diffusive pathway.

In the case of permeates, which interact with the polymer and can induce swelling, a decrease of the overall mass transport can be achieved by the hindering of the polymer relaxation. Espino et al. [START_REF] Espino-Pérez | Designed cellulose nanocrystal surface properties for improving barrier properties in polylactide nanocomposites[END_REF] showed that the anisole transport, which is an aroma compound and has high affinity to PLA, could be decreased by the inclusion of neat and surface-grafted cellulose nanocrystals. Here, the anisole mass uptake of the PLA/ZnO nanocomposites was measured with the help of a microbalance. This method has the advantage that small samples can be used and no mechanical stress exerted. The experimental data shown in Fig. 6a report on the mass uptake kinetics of the sample set. The spikes on the curves originate from the measurement method and can be cancelled out. Fig. 6b shows the transport data on the time scale normalized with the sample thickness in the aim to easier compare the mass uptake rates. Except PLA at long times, all curves linearized on the squared root of time, which indicates an essentially Fickian mass transport mechanism. Because of the sample thickness of the nanocomposites, the non-Fickian transport mechanisms could not be observed within the experimental time scale. The curves show furthermore that within the experimental time, the equilibrium mass uptake was not reached. In the case of the PLA blank, it was possible to obtain a thinner sample, which allowed to approach the mass uptake of the Fickian equilibrium (m eq,app = 0,34%) and to calculate an apparent diffusion coefficient (D app = 6.0.10 -14 m 2 s -1 ), which was consistent L: half sample thickness, s: slope.

with Espino's work [START_REF] Espino-Pérez | Designed cellulose nanocrystal surface properties for improving barrier properties in polylactide nanocomposites[END_REF]. Because of the high uncertainty of the value of the equilibrium mass uptake, in the following only the kinetics of the transport obtained from the initial slopes is discussed. The computed data are shown in Table 4. The comparison of the slopes showed faster mass transport in the nanocomposites. A similar result was obtained upon the study of water mass transport in PLA/ZnO composites [START_REF] Pantani | PLA-ZnO nanocomposite films: water vapor barrier properties and specific end-use characteristics[END_REF]. Low interfacial adhesion bringing supplementary voids to the material might be a reason for this behaviour. PLA/ZnO_s nanocomposites had, at the same concentration, higher diffusion coefficients than PLA/ZnO_r nanocomposites. The ZnO_s particles were more aggregated compared to the ZnO_r particles which can be observed in the TEM micrographs (Fig. 1) and inferred from the low UV shielding effect of ZnO_s. In conclusion, the rod shape seems to be favourable to obtain a more homogenous dispersion of ZnO particles in PLA. The higher dispersion quality brings about better properties.

Conclusions

Here we report on the structural relaxation, gas transport and anisole (an aroma compound) sorption kinetics of PLA/ZnO nanocomposites having both spherical and rod-shaped nanoparticles. Results are analyzed in terms of nanoparticle morphology, concentration and dispersion within the amorphous hosting PLA matrix. Morphological analysis reveals a homogeneous distribution of ZnO_r within the hosting PLLA matrix, while ZnO_s nanoparticles remain highly aggregated. This better dispersion of rod-shaped nanoparticles in regard with spherical ones leads to improved UV-Vis properties, where UV light is markedly blocked while the samples remains transparent in the visible region. Higher physical aging rates were obtained upon both ZnO_r and ZnO_s incorporation as a result of the low interfacial filler-matrix adhesion, which in turn speeds up macromolecular conformation towards the thermodynamic equilibrium. Finally, oxygen and helium permeability tests and anisole sorption tests were carried out as a model gases for the determination of the gas transport properties of developed nanocomposites. These results confirmed the poor interfacial adhesion observed during the structural relaxation analyses, which together with a low tortuosity effect, yield increased permeability values after nanoparticle addition. The results here reported may serve to raise awareness on the relevance of nanoparticle morphology and crystalline/ amorphous character of the hosting matrix on the resulting relaxation and transport properties of the nanocomposites.

Fig. 1 .

 1 Fig. 1. Representative TEM images showing bare ZnO-s (a) and ZnO-r (d) nanoparticles and their dispersion within the PLA matrix: PLA/ZnO_s 1 wt% (b); PLA/ZnO_s 10 wt% (c); PLA/ZnO_r 1 wt% (e) and PLA/ZnO_r 10 wt% (f).

Fig. 2 .

 2 Fig. 2. A) Thermogravimetric traces and b) weight lost rates of PLA/ZnO_r nanocomposites; c) thermogravimetric traces and d) weight lost rates of PLA/ZnO_s nanocomposites.

Fig. 3 .

 3 Fig. 3. UV-Vis transmittance spectra of PLA/ZnO_r films (a) and PLA/ZnO_s films upon ZnO concentration.

Fig. 4 .

 4 Fig. 4. Increase of the enthalpy of recovery with the aging time normalized to the PLA content of the nanocomposites. A) enthalpy of recovery of PLA and PLA/ZnO_r nanocomposites, b) enthalpy of recovery of PLA and PLA/ZnO_s nanocomposites.

Fig. 5 .

 5 Fig. 5. Enthalpy of recovery functions of PLA and PLA/ZnO nanocomposites. A) PLA/ZnO_r nanocomposites at different ZnO_r concentrations, b) PLA/ZnO_s nanocomposites at different ZnO_s concentrations, c) shifted enthalpy of recovery function of PLA/ZnO_r nanocomposites to the curve of PLA, d) shifted enthalpy of recovery function of PLA/ZnO_s nanocomposites to the curve of PLA.

Fig. 6 .

 6 Fig. 6. Anisole sorption kinetics in PLA and PLA/ZnO nanocomposites, a) experimental sorption data, b) sorption data on the square root of experimental time normalized by the sample thickness. The perturbation in the mass uptake curves correspond to the re-regulation of the anisole activity in the measurement cell to avoid balance drift.

Table 1

 1 Characteristic thermodegradation temperatures (T 5% , T 10% and T peak ) and maximum degradation rates (α max ) of neat PLA and its ZnO_r and ZnO_s nanocomposites.

	Sample T PLA ZnO (wt. %) 0 330.8	340.9	366.2	-0.00632	21.8
	PLA/ZnO_r 0.2	257.6	267.3	300.5	-0.00534	28.8
	PLA/ZnO_r 0.5	250.9	260.9	293.5	-0.00596	24.2
	PLA/ZnO_r 1	250.4	259.3	288.8	-0.00682	20.6
	PLA/ZnO_r 5	242.4	252.6	281.9	-0.00618	21.6
	PLA/ZnO_r 10	237.1	247.1	276.8	-0.00565	23.2
	PLA/ZnO_s 0.2	277.6	287.5	335.2	-0.00346	45.1
	PLA/ZnO_s 0.5	262.6	271.6	308.4	-0.00550	24.4
	PLA/ZnO_s 1	254.9	264.2	297.3	-0.00573	25.3
	PLA/ZnO_s 5	250.2	260.7	289.1	-0.00682	20.9
	PLA/ZnO_s 10	243.3	252.8	281.7	-0.00616	20.8

5% (°C) T 10% (°C) T peak (°C) α max (%/°C) FWHM (°C)

Table 2

 2 Transmitted light (%) in the UV Region (360 nm) and visible region (540-560 nm) through nanocomposites films having ~10 μm thickness.

		PLA/ZnO_r		PLA/ZnO_s
		360 nm	540-560 nm	360 nm	540-560 nm
	0 wt.%	94.8	96.0	94.8	96.0
	0.2 wt.%	88.4	90.6	87.9	91.7
	0.5 wt.%	74.5	86.7	86.7	89.2
	1 wt.%	64.9	83.2	79.6	81.4
	5 wt.%	38.3	80.7	72.1	75.9
	10 wt.%	21.5	73.3	67.7	70.9

Table 3

 3 Summary of thermal properties of PLA and PLA/ZnO nanocomposites.

	Sample	ZnO (wt.%)	T g (°C)	Width (°C)	C Δ p (J g -1 min -1 )	∞ ΔH (J g -1 )	β (J g -1 /lg (min))
	PLA	0	56 ± 1	5.1 ± 0.5	0.554 ± 0.012	9.2	2.3 ± 0.1
	PLA_ZnO_r	0.2	50 ± 2	6.4 ± 1	0.562 ± 0.005	6.2	2.8
	PLA_ZnO_r	0.5	50 ± 2	5.9 ± 0.5	0.554 ± 0.009	6.2	2.6
	PLA_ZnO_r	1	52 ± 1	5.5 ± 0.5	0.556 ± 0.016	7.0	2.5
	PLA_ZnO_r	5	49 ± 1	6.5 ± 0.5	0.574 ± 0.039	6.0	2.1
	PLA_ZnO_r	10	49 ± 1	6.7 ± 0.5	0.520 ± 0.005	5.1	2.0
	PLA_ZnO_s	0.2	57 ± 1	5.4 ± 2	0.526 ± 0.005	9.3	2.4
	PLA_ZnO_s	0.5	56 ± 1	4.0 ± 0.5	0.536 ± 0.005	9.2	2.3
	PLA_ZnO_s	1	51 ± 1	5.0 ± 0.5	0.557 ± 0.019	6.7	2.5
	PLA_ZnO_s	5	50 ± 1	5.5 ± 0.5	0.550 ± 0.008	6.2	2.0
	PLA_ZnO_s	10	51 ± 1	7.2 ± 0.5	0.549 ± 0.005	6.7	2.1
	T

g : Mid-point glass transition temperature; C Δ p : heat capacity step at T g normalized by the PLA content; ∞ H Δ : enthalpy of recovery for an infinitive duration of aging after Eq. (

4

).

Table 4

 4 Transport properties of anisole mass uptake in PLA and PLA/ZnO nanocomposites.

	Sample	ZnO (wt.
		%)
	PLA		0.556	10.9	6.1
	PLA/ZnO_r 0.2	1.52	3.53	5.4
	PLA/ZnO_r 0.5	1.67	7.47	12
	PLA/ZnO_r 1	1.98	2.44	4.8
	PLA/ZnO_r 5	1.82	4.96	9.0
	PLA/ZnO_r 10	2.64	5.43	14
	PLA/ZnO_s 0.2	1.58	4.94	7.8
	PLA/ZnO_s 0.5	1.70	5.43	9.2
	PLA/ZnO_s 1	1.77	5.61	9.9
	PLA/ZnO_s 5	1.71	6.11	10
	PLA/ZnO_s 10	No film

L (m) × 10 -4 s (s 0.5 ) × 10 -6 s × L (s 0.5. m) × 10 -10

Acknowledgements

Technical and human support provided by SGIker (UPV/EHU, MICINN, GV/EJ, EGEF and ESF) is gratefully acknowledged. Authors thank the Basque Country Government for financial support (IT718-13).

Data availability

All the data used to support the findings of this study are included within the article.

Appendix A. Supplementary data

Supplementary data to this article can be found online at https:// doi.org/10.1016/j.polymertesting.2019.02.009.


Mohamed Besher Baradi

Camilo Cruz

Thomas Riedel

Gilles Régnier
email: gilles.regnier@ensam.eu

Mechanical and microstructural characterization of flowing weld lines in injection-molded short fiber-reinforced PBT

Keywords:

The aim of this work is an extensive experimental mechanical and microstructural characterization of flowing weld lines (WLs) in injection-molded short glass fiber-reinforced polybutylenterephthalate (PBT) using X-ray computed tomography and digital image correlation (DIC). It was found that the induced fiber orientation (FO) in a flowing WL is similar to that induced by flow along a wall. In this test case, the impact of the flowing WL on the FO did not vanish after a flow length of 70 mm. The shape of the inserts, which originated the flowing WLs, only had a marginal effect on the induced FO gradient. By reducing part thickness, the erasing of the FO gradient induced by the WL is reached at shorter flow distances. At the WLs, there is a reduction of the fiber volume fraction in comparison to the regions far from the WL plane. DIC results show a pronounced strain localization at the WL, which can be explained by the FO gradient induced by the WL.

Introduction

In injection-molded thermoplastic parts, weld lines (WLs) form during the filling phase at the interface where two melt fronts collide. Multiple injection gates, inserts or pins inside the mold cavity lead to the separation of the melt and WL formation. WLs induce a significant reduction of the mechanical properties, especially in the case of short fiber-reinforced polymers. The occurrence of WLs is nearly unavoidable in complex industrial parts and they often cause the failure of those parts [START_REF] Fellahi | Weldlines in injection-moulded parts: a review[END_REF]. Current industrial and environmental needs demand a reliable prediction of the mechanical properties at the WLs for an improved economical design. This requires comprehensive physical and mechanical understanding of the process-induced microstructure at the WLs. The effect of WLs on the mechanical properties of fiber-reinforced thermoplastics has been studied extensively [START_REF] Fellahi | Weldlines in injection-moulded parts: a review[END_REF][START_REF] Criens | The influence of knit-lines on the tensile properties of injection molded parts[END_REF][START_REF] Nguyen-Chung | Effect of an obstacle during processing on the weld line of injection-molded glassy polystyrene: microhardness study[END_REF][START_REF] Seldén | Effect of processing on weld line strength in five thermoplastics[END_REF][START_REF] Oh | Terahertz time-domain spectroscopy of weld line defects formed during an injection moulding process[END_REF][START_REF] Yamada | Evaluation of mechanical properties of adjacent flow weldline[END_REF][START_REF] Malguarnera | The effects of processing parameters on the tensile properties of weld lines in injection molded thermoplastics[END_REF][START_REF] Hashemi | Thermal effects on weld and unweld tensile properties of injection moulded short glass fibre reinforced abs composites[END_REF][START_REF] Baradi | Mechanical characterization of frontal and flowing weld lines in injection-molded short fiber-reinforced thermoplastics[END_REF][START_REF] Liu | Experimental matrix design to optimize the weldline strength in injection molded parts[END_REF][START_REF] Kim | Morphology and mechanical properties of injection molded articles with weld-lines[END_REF][START_REF] Thi | Measurement of fiber orientation distribution in injection-molded short-glass-fiber composites using X-ray computed tomography[END_REF][START_REF] Hannesschläger | Fibre structure characterisation of injection moulded short fibre-reinforced polymers by X-ray scatter dark field tomography[END_REF][START_REF] Ayadi | Determination of orthotropic properties of glass fibre reinforced thermoplastics using X-ray tomography and multiscale finite element computation[END_REF].

Most of the research on WLs shows mechanical characterizations considering the influence of process parameters, filler content and angle of WL formation. According to literature, the WL strength is improved slightly by setting higher melt temperatures [START_REF] Malguarnera | The effects of processing parameters on the tensile properties of weld lines in injection molded thermoplastics[END_REF][START_REF] Liu | Experimental matrix design to optimize the weldline strength in injection molded parts[END_REF]. It has been shown that the severity of a WL increases with the filler content [START_REF] Hashemi | Thermal effects on weld and unweld tensile properties of injection moulded short glass fibre reinforced abs composites[END_REF]. On the other hand, some authors reported an increase of WL strength by enlarging the insert size for flowing WLs made of unreinforced polystyrene [START_REF] Liu | Experimental matrix design to optimize the weldline strength in injection molded parts[END_REF]. The weakness of a WL can be attributed to the following phenomena [START_REF] Fellahi | Weldlines in injection-moulded parts: a review[END_REF][START_REF] Criens | The influence of knit-lines on the tensile properties of injection molded parts[END_REF][START_REF] Nguyen-Chung | Effect of an obstacle during processing on the weld line of injection-molded glassy polystyrene: microhardness study[END_REF][START_REF] Seldén | Effect of processing on weld line strength in five thermoplastics[END_REF][START_REF] Oh | Terahertz time-domain spectroscopy of weld line defects formed during an injection moulding process[END_REF][START_REF] Yamada | Evaluation of mechanical properties of adjacent flow weldline[END_REF][START_REF] Malguarnera | The effects of processing parameters on the tensile properties of weld lines in injection molded thermoplastics[END_REF][START_REF] Hashemi | Thermal effects on weld and unweld tensile properties of injection moulded short glass fibre reinforced abs composites[END_REF][START_REF] Baradi | Mechanical characterization of frontal and flowing weld lines in injection-molded short fiber-reinforced thermoplastics[END_REF]:

1 Incomplete molecular diffusion (insufficient healing of the polymer matrix): Due to the rapid cooling rates in injection molding, the polymer chains might have less time than required to reptate and reach the entanglement density of the bulk. 2 Strong fiber orientation (FO) gradient: Fibers align parallel to the WL surface, substantially reducing their role in strengthening the part across the WL plane. 3 High molecular orientation: Because of the fountain flow phenomenon, the molecules also orient parallel to the WL surface. The molecular orientation however has a negligible contribution to the strength reduction in comparison with the FO gradient [START_REF] Kim | Morphology and mechanical properties of injection molded articles with weld-lines[END_REF]. 4 V-notches: The formation of V-notches at the surface of WLs has been documented elsewhere [START_REF] Fellahi | Weldlines in injection-moulded parts: a review[END_REF][START_REF] Yamada | Evaluation of mechanical properties of adjacent flow weldline[END_REF]. They act as stress concentrators at the WLs, but their formation can be avoided by using ventilation channels and controlling the process parameters, especially the packing (holding) pressure.

Investigations of the microstructure near the WLs and its influence on the mechanical properties is less common in the literature. In this work, the focus is put on the microstructural aspect of WLs. Vaxman et al. [START_REF] Vaxman | Weld-line characteristics in short fiber reinforced thermoplastics[END_REF] used reflected light microscopy to show that in a frontal WL the fibers tend to align parallel to the WL surface. Other researchers used scanning electron microscopy for studying the FO at the WLs [START_REF] Baradi | Mechanical characterization of frontal and flowing weld lines in injection-molded short fiber-reinforced thermoplastics[END_REF] and found that the fibers in flowing WLs orient in the flow direction through the thickness, eliminating the typical skin-core microstructure.

The use of X-ray computed tomography (CT) enables a three-dimensional reconstruction of the microstructure of fiber-reinforced thermoplastics. Prade el al. [START_REF] Prade | Nondestructive characterization of fiber orientation in short fiber reinforced polymer composites with X-ray vector radiography[END_REF] used X-ray vector radiography on fiber-reinforced parts with flowing WLs to validate the characterization method. The authors showed a qualitative assessment of the FO at the WL, but the detailed microstructure of the WL does not appear in their publication. In other work [START_REF] Thi | Measurement of fiber orientation distribution in injection-molded short-glass-fiber composites using X-ray computed tomography[END_REF], based on several CT scans at different positions along a flowing WL, an increase of FO in the out-of-plane direction at the core of the WL was reported. Another group used X-ray scatter dark field imaging and X-ray CT to describe the FO at a flowing WL [START_REF] Hannesschläger | Fibre structure characterisation of injection moulded short fibre-reinforced polymers by X-ray scatter dark field tomography[END_REF]. They confirmed that in a flowing WL the main FO direction is parallel to the flow. Microstructural analysis with X-ray CT at a flowing WL by Ayadi et al. [START_REF] Ayadi | Determination of orthotropic properties of glass fibre reinforced thermoplastics using X-ray tomography and multiscale finite element computation[END_REF] showed a complex FO across the thickness after the insert. The authors noticed a FO gradient around the WL and random FO at the weld region. The higher strain localization at the WL was attributed to the existence of out-of-plane fibers at the core layer of the WL zone. Using a larger scale analysis at the WL, Ayadi et al. [START_REF] Ayadi | Large-scale X-ray microtomography analysis of fiber orientation in weld line of short glass fiber reinforced thermoplastic and related elasticity behavior[END_REF] arrived to the same conclusion. In addition, they analyzed the FO and fiber volume fraction (VF) quantitatively, showing changes both in WL direction and transversally to the WL plane. According to the publication, the variations in VF are also responsible for the variations of the mechanical properties.

Yamada et al. [START_REF] Yamada | Evaluation of mechanical properties of adjacent flow weldline[END_REF] showed that in a flowing WL (made of an unreinforced thermoplastic) the mechanical behavior is highly influenced by the molecular morphology at the surface of the WL. The main factor reducing the mechanical properties at the WL was the V-notch, which induced a stress concentration over the polymer layer beneath it, which was not completely healed. Furthermore, they noticed a high molecular orientation in the flow direction at the WL. The use of DIC to study the strain field during the tensile test of a bar containing a flowing WL was reported by Kalus et al. [START_REF] Kalus | Measuring deformation and mechanical properties of weld lines in short fibre reinforced thermoplastics using digital image correlation[END_REF]. They showed the development of a strong strain localization at the WL. In the case of a 30 wt% glass fiber-reinforced PA12, the strain was localized in a region of 2 mm around the WL. Based on the convergence of the estimated local elastic modulus at the WL for the fiber-reinforced and unreinforced PA12 tensile bars, they suggested that there is no fiber crossing at the WL in the fiber-reinforced case.

In general, the available characterizations of the microstructure at WLs are mainly qualitative and limited to relative small regions around the WL. In this work, a specific mold to perform a comprehensive microstructural study of flowing WLs was developed. Firstly stresses and strains at break were measured to mechanically quantify the effect of WLs. Both FO and VF at different positions along the WL were quantified by analyzing X-ray computed tomography data. The effect of the insert shape and part thickness on the induced FO gradients was studied. The plate thickness changes at the WL were also measured in order to correlate it with the induced FO. The induced FO and the mechanical tensile tests results were analyzed in front of the strain fields measured with digital image correlation (DIC).

Experimental methods

Sample preparation

A polybutylene-terephthalate reinforced with 30 wt% glass fibers (average VF of 0.18) was used. The glass fibers in the composite had a mean fiber length of 200 μm, a diameter of 10 μm and cylindrical shape. The flowing WL mold was designed with the assistance of rheological simulations using Moldflow ® 2016. It was found that the bigger the insert the greater the influence on the FO at the WL. A plate-form mold cavity of 125 × 125 mm with two thicknesses (1.5 and 3.0 mm) and two insert shapes (slit with 1 mm width and cylindrical with 15 mm radius) was used, as shown in Fig. 1.

The injection molding was carried out in a DK-Codim 175/410 injection molding machine. The process parameters were fixed as follows:

Packing pressure of 40 MPa, injection time of 2.3 s, mold temperature of 60 °C and melt temperature of 250 °C. Plates without insert (i.e. without WL) were injected in order to extract tensile test specimens without a WL at a distance of 25.5 mm after the insert tip. The aim of using the same injection time for both thicknesses was to obtain similar velocities at the flow front, as usually encountered in industrial parts. Using water jetting, tensile test specimens were cut from the plates with WL at six different flow distances from the insert tip: 3.5, 14.5, 25.5, 47.5, 69.5 and 91.5 mm. The flowing WL plane was in the middle of each tensile test specimen, which had a width of 5 mm at this location. The total length of the specimen was 85 mm (Fig. 1).

Mechanical testing and digital image correlation

Quasi-static tensile tests were carried out using an Instron 5966 machine with an optical extensometer, with a gauge length varied between 10 and 13 mm. The clamp speed was fixed to 10 mm/min. The strain fields under tensile loading of some selected specimens were also evaluated with DIC using a Q-400 system from LIMESS combined with ISTRA4D (4.4.4) software from Dantec Dynamics. In this case, the tensile testing machine was a Schenck POZ 160. The strain fields were analyzed using a pixel size ranging between 17 and 21 μm.

Thickness measurements

Thickness measurements were made by means of a 3D profilometry technique using a Veeco Dektak 150 with a 12 μm stylus.

X-ray computed tomography

Some tensile test specimens, cut at flow distances of 3.5, 25.5 and 69.5 mm after the insert, were selected for microstructural analysis and were labelled as positions 1, 2 and 3, respectively, (see Fig. 1). The CT scan was performed directly on the tensile test bar by fixing it on a holder. Volumes of 5 × 10 × 3 mm 3 or 5 × 10 × 1.5 mm 3 for 8 cases were scanned, as explained in Table 1.

A CT system GE v|tome|x m 300/180 Metrology Edition with a Nanofocus x-ray tube was employed using the following parameters: Mode 1, x-ray tube voltage 100 kV, x-ray tube current 135 μA, exposure time per projection 500 ms, 3300 projections, no pre-filtration and a voxel size of 2.9 μm. The system used a 16 MPixel detector dynamic 41/ 100 without any binning. Each scan took ~3 h and resulted in a data set of about 20 GB size.

Determination of fiber orientation and volume fraction

The CT data were analyzed using VG Studio Max 3.1 software from Volume Graphics. The evaluation of FO and VF was based on a local filtering method, which uses gradient threshold values and computes a local FO for each voxel. This method does not perform a segmentation of single fibers. A regular hexahedral mesh was used for averaging the computed voxel values inside each mesh element. The element size was fixed at 0.2 mm, which is statistically representative for the given fiber content. In fact, it was estimated that on average nearly 70 fibers lie inside each mesh element. The mesh for microstructure reconstruction was placed in such a way that the three main FO tensor components correspond to the directions of the Cartesian coordinates shown in Fig. 1. This means that the a 11 component corresponds to the FO in x direction (flow direction), a 22 is the FO in y direction (transversal to flow) and a 33 is the FO in z direction (thickness direction). These three directions are very close to the main principle directions given the very low values of the off-diagonal tensor components. In the evaluation, the outer layers in flow direction were not considered, i.e. at the water jet boundaries. For each mesh element, VG Studio Max 3.1 computed a second-order FO tensors, as defined by Advani and Tucker [START_REF] Advani | The use of tensors to describe and predict fiber orientation in short fiber composites[END_REF]. Fig. 2 shows an example of the results obtained on the mesh. The regular mesh can be divided along a given direction into layers containing the elements at the given direction coordinate. In the following, the FO results are presented by averaging the values of each layer according to the direction of analysis. The standard deviation of the mean value depends on the number of layers under consideration, however their maximal values was of 0.022 for a single evaluation point.

Results and discussion

Mechanical properties of flowing weld lines

Based on previous rheological tests, it was known that the healing ability of the PBT matrix is very high. Indeed, at the selected injection temperature the relaxation time of the unreinforced PBT matrix is approximately 4 ms [START_REF] Baradi | Mechanical characterization of frontal and flowing weld lines in injection-molded short fiber-reinforced thermoplastics[END_REF].

This means that the FO is the dominant factor controlling the mechanical properties of the flowing WLs for the material at hand. In fact, V-notches did not appear at the WLs and failure always took place at the WL. In general, a strength reduction factor of 0.85 (reduction of 15%) was found for flowing WLs, which is higher than the factor of frontal WLs [START_REF] Baradi | Mechanical characterization of frontal and flowing weld lines in injection-molded short fiber-reinforced thermoplastics[END_REF]. In Fig. 3, the evolution of the mean failure strain and stress along the flowing WL is plotted for all the cases of the experimental work. Breaking of the specimens with WL always took place exactly at the WL plane. Results show that mostly, a slow increase of failure strength and a failure strain towards the values without WL takes place along the distance of the flowing WL. Given the high healing ability of the PBT matrix, the increase of failure strain can be attributed to the changes in FO along the WL. However, while the measured failure stress reflects the WL strength, the measured engineering failure strains do not reflect necessarily the strain levels at the WL.

The measured failure strain with the optical extensometer is in fact the sum of both failure strain at the WL and strain outside of the WL region (Fig. 3). Using DIC during the tensile tests, a local failure strain at the WL of around 4-8% was determined. This means that the increase of engineering failure strain towards the failure strain values of without WL observed in Fig. 3b is related to the changes of FO at and around the WL. Failure stress just after the cylindrical insert for the 1.5 mm plate was the highest compared to other positions along the flowing WL; it even surpassed the quasi-static strength of the case without WL.

These results show the severity of the WLs and the necessity to study the influence of the process-induced microstructure on the mechanical properties at the WL. No V-notches at the WL were found in the samples

Table 1

Positions of the computed tomography scans for each case of flowing WL. studied, a fact that rules out their role on the mechanical properties obtained.

Fiber orientation

In general, the FO of the flowing WLs exhibits some common aspects. The near the WL plane the more the fibers orient in flow direction reaching a maximum of alignment at the WL plane. The FO tensor at the WL plane was similar in all cases, and the tensor component in flow direction (a 11 ) had an average value of 0.8 across the thickness (Fig. 4a). A slight increase of the FO tensor component was also observed in the thickness direction (a 33 ) at the WL. This FO gradient in y direction, transversally to flow direction, can explain the phenomenon of strain localization at the WL observed with DIC. This strain localization could be explained by a higher orientation in the x direction that is perpendicular to the loading direction (y). The width of the FO gradient reached ~8 mm, i.e. ~4 mm from each side of the WL plane. However, the width of the FO gradient varied depending on the part thickness and the flow distance after the insert.

Fiber orientation across the thickness

Fig. 5 presents the FO tensors divided in different layers across the thickness for the 3.0 mm and 1.5 mm plates in order to better understand the induced FO at the WL. Seven layers across the thickness can be identified: Skin, shell, intermediate, core, intermediate, shell and skin. Table 2 shows the thicknesses of the mentioned layers. The skin layer, which is in contact with the mold wall, was the least influenced by the WL, whose FO gradient had a width of ~2 mm (Fig. 5a). The nearer to the core, the greater the FO gradient around the WL. In fact, the FO gradient widened for the shell and inner layers (Fig. 5(b andc)). The width of FO gradient was the largest at the core and reached more than 8 mm in the plate with a thickness of 3 mm. The core layer was highly oriented in the flow direction and showed a slight tendency to orient in the out-of-plane direction, which reached a maximum exactly at the WL plane (Fig. 5d).

Thickness measurements at the WL on the same samples scanned by CT are presented in Fig. 6. An increase of thickness was found at the WL of around 9 μm and 15 μm for the 1.5 mm and 3.0 mm plate, respectively. In the thicker plate, the region of increased thickness was wider. This region of increased thickness was persistent along the WL until the end of the plate.

The increase of FO in the out-of-plane direction at core can explain the changes in plate thickness at the WL presented in Fig. 6. The fact of having more fibers oriented in the thickness direction reduced the shrinkage at the WL in comparison with regions far apart from the WL plane. The typical skin-core microstructure induced by flow between parallel walls disappeared progressively as the WL plane was approached, where fibers in the core layer were predominately oriented in the flow direction. This shows the huge influence of the WL on the FO, especially in the inner layers of the plate. The later solidification of the inner layers in comparison to the skin layer can explain the higher influence of the flowing WL on the fiber orientation at the core. In the plates of 1.5 mm, the WL had a high influence on the FO only at the core layer, where the width of the FO gradient was ~6 mm. In the other layers, the width of FO gradient was reduced to 2 mm. Fig. 7 compares the width of averaged FO in thickness direction between the two thicknesses under study. The WL in the thicker plate, which has larger shell and core regions, exhibited a larger FO variation in transversal direction than that of the thinner plate. The difference in FO variation in transversal direction can be explained in fact by the results in Fig. 5. The FO in the thicker plate was more influenced by the weld line compared to the FO in the thinner one. The higher FO in the flow direction across the thickness for the 1.5 mm plate (Fig. 7) explains the lower tensile strength (loading in transversal direction) of the corresponding flowing WLs (Fig. 3). The preferential FO in the flow direction at the core and the width of FO gradient indicate that a flowing WL has a comparable effect on FO as a mold wall. In fact, similar to a mold wall, the fibers oriented parallel to the WL surface and the FO gradient spread some millimeters from the wall.

In Fig. 8, the averaged FO at the WL across the normalized thickness is compared for the two thicknesses under study. The FO averaging at the WL was performed using all the FO evaluations, which lay at normal distances shorter than 2 mm from the WL plane (between y = -2 mm and y = 2 mm). In both slit and cylindrical inserts, the averaged FO profile along the thickness is similar for the two thicknesses under study. The typical core region disappeared with high orientation in the transversal-to-flow direction (y).

Influence of the insert shape

To compare the role of the insert shape on the induced FO, Fig. 9 (a and b) compares the slit with the cylindrical insert by averaging, across the thickness, two mesh layers 0.5 mm after the insert and two mesh layers 5 mm after the insert. The induced FO was different at short flow distances (0.5 mm after the insert), but then the induced FO reached a comparable state independently of the insert type. Actually, the FO induced by the cylindrical insert required around 3.5-4 mm after the insert to reach a similar FO state as in the slit insert (Fig. 10). This means that the insert shape has the main influence on the induced FO at the meeting point of the two flow fronts. Simulations of the flow behavior around the cylindrical insert showed that the two flow fronts meet frontally at the beginning of the WL and then the two flows are oriented quickly in the flow direction. The changes in FO just after the insert could be related to the increase of the shearing deformation caused by the insert walls of the cylindrical insert and to the change of the velocity direction of flow front: it is very difficult to quantify the influence of these two phenomena. A slight difference in the FO gradient (variation in transversal direction) was noticed according to the insert used; in fact, the zone of FO distortion around the WL for the slit insert was narrower than that for the cylindrical insert. The insert shape also had an impact on the failure strain along the WL length, as presented in Fig. 3. For example, the specimens extracted from the 3 mm plates with cylindrical insert exhibited a higher failure strain than the specimens extracted from the 3 mm plates with slit insert. This is in agreement with the higher FO gradients in the flowing WLs produced by the slit insert.

In order to get more information about the influence of the insert on the induced FO around the WL, an additional CT scan was performed at position 3 of a 3.0 mm plate injected with a cylindrical insert. The scanned volume was asymmetrical with respect to the WL plane (y = 0) in order to include a larger area transversally to the WL. Fig. 11 shows that the cylindrical insert induced a smaller FO gradient in the transversal direction than the slit insert. As mentioned before, this difference in FO can explain the difference in failure strain presented in Fig. 3, given the fact that a smaller FO gradient in transverse direction would lead to larger failure strains. The microstructural information in Fig. 11 also reveals that the width of the FO gradient can be larger than 10 mm, even after 70 mm of flowing distance. Fig. 12 shows the changes of induced FO along the WL for the different insert types. The induced FO tensor at the WL plane and 0.5 mm away from that plane did not change substantially along the WL. However, the FO gradient seems to flatten gradually with the distance from the insert tip. In the 1.5 mm plate, the FO gradient induced by the WL was strongly smoothed after 70 mm.

In order to have a better understanding of the changes along the WL, Fig. 13 presents the induced FO per layers (as in Fig. 5). For the 3 mm plates, the FO gradient changed strongly at short flow distances in the skin and shell layers until it is hardly detectable 65 mm after the insert. The FO gradient in the intermediate layer exhibits, however, progressive decrease with the flow distance. The fibers at core, on the contrary, kept their orientation along the WL with only a slight change 65 mm after the insert tip. In the case of 1.5 mm, after 65 mm of flow distance, the FO gradient is only detectable at skin and core. Unlike at the core, where there was no reorientation along the WL; at the skin, shell and intermediate layers the induced FO at the WL evolved substantially with the flow distance. The existence of a FO gradient along the WL explains why failure always took place at the WL, even when the WL failure stress (90 mm after the insert tip) is comparable to the failure stress of a sample without WL.

The smaller FO gradient for the 1.5 mm plates could explain the quicker increase of the failure stress towards the value of the bulk strength after the insert (Fig. 3), where it can be postulated that the FO gradient disappeared at shorter flow distances from the insert tip (compared to that one in 3 mm plates), thus, progressively minimizing the WL localization influence. Because of the higher FO in the flow direction obtained in the thinner parts, the FO variation in transversal direction caused by the WL became smaller, which in turn shortened the flow distance for recovering the bulk material strength along the WL.

The part thickness influenced the FO profile along the flow direction. In the 1.5 mm plate, the FO gradient transverse to flow direction across the layers (Fig. 13) was more affected by the flow distance along the WL than that in the 3 mm plate. The similar injection time of both plates induced a higher shear rate in the thinner plate. On one hand, the higher shear rates oriented the fibers predominantly in the flow direction around the WL, reducing the FO gradient at the WL region and enabling a higher increase of the mechanical properties (failure stress and failure strain) towards the values without weld line for the 1.5 mm plates. On the other hand, in the 3 mm plate, only the two outer layers (skin and shell) were influenced by the proximity to the cavity walls and show a reduction of the FO gradient. The lower influence on the intermediate and core layers preserved, in turn, the FO gradient for the 3 mm plate and explains the low increase of mechanical properties along the flow distance.

Using scanning electronic microscopy on the same samples, an FO gradient was found at the core layer, even 91 mm after the insert [START_REF] Baradi | Mechanical characterization of frontal and flowing weld lines in injection-molded short fiber-reinforced thermoplastics[END_REF]. The decrease of the FO gradient (transversal to the WL plane) in combination with the strain increase observed in Fig. 3 suggests that the impact of a flowing WL on the mechanical properties would vanish only when the induced FO gradient disappears completely.

Correlation with DIC results

In order to achieve a better understanding of the mechanical behavior at the WL, the strain field during a quasi-static tensile test of the same specimens used for the CT scans was tracked using DIC. In Fig. 14, the strain field in the loading direction (transversal to the WL on one line) is presented for different positions along the WL and for different types of insert. Each strain profile represents the same macroscopic stress level (approximately 60 MPa). A strain gradient, which reaches a width of around 8 mm and has a maximum at the WL plane can be noted. At a given stress level, there was a reduction of the strain gradient with the flow distance from the insert tip. The strain localization at the WL can be correlated directly with the FO gradient shown in Figs. 5 and13a. The strong strain localization at the WL is, in fact, the result of the unfavorable fiber orientation around the WL (fibers are mainly oriented in flow direction, transversally to the loading direction); which in turn concentrates the stresses in the polymer matrix. The difference in engineering failure strain (measured with extensometer) of the 3 mm samples extracted in position 3 between the slit and cylindrical inserts (Fig. 3b) can be explained in the light of the DIC results in Fig. 14b, where higher strain levels are observed for the cylindrical insert, both at the WL plane and far from it. The strain profiles in 14b correspond to a stress level of ~60.5 MPa, which is close to the failure stress of the sample extracted from the plate with cylindrical insert. The difference of measured strains between the cylindrical and slit insert correlates with the results of FO in Fig. 11. In fact, in the plates with a cylindrical insert the fibers away from the WL plane (2-3 mm from plane) were more oriented in the flow direction than their counterparts in the plates with a slit insert. This would lead to higher deformations when the sample was loaded transversally to the flow direction.

The results from DIC were also used to estimate the local Young's modulus (E) at the WL region and at zones far away from the WL plane. The moduli were calculated by taking the difference of global stress corresponding to a change of local strain going from 0.15% to 0.8%. It was found that the local Young's modulus diminishes with proximity to the WL plane. For all flowing WL cases, E at the WL ranged between 3650 and 3900 MPa. For points in a section 4 mm away from the WL plane, E increased and ranged between 4500 and 5000 MPa. In order to check if the measured local Young's moduli are plausible, the local Young's moduli were estimated analytically by mean field homogenization using the FO measured via CT. The Mori-Tanaka model [START_REF] Mori | Average stress in the matrix and average elastic energy of materials with misfitting inclusions[END_REF] was chosen for the calculations. In terms of the polymer matrix, a Young's modulus of 2230 MPa and a Poisson's ratio of 0.4 were used. These values were obtained from tensile tests performed on injectionmolded samples of the non-reinforced PBT compound. For the E-glass fibers, a Young's modulus of 72 GPa, a Poisson's ratio of 0.22 [START_REF] Kaiser | Modified mean-field formulations for the improved simulation of short fiber reinforced thermoplastics[END_REF] and an aspect ratio of 20 were set. The analytical calculations were done in two regions, one exactly at the WL plane and the other 4 mm away from the WL plane for the case of 3 mm plate at position 1 with a slit insert. A multilayer representative volume element was generated using the same layers and the FO tensors given in Fig. 13. The glass fiber volumetric fraction was defined equal to 0.18 in all layers. In this framework, estimated theoretical E values at the WL and 4 mm away from the WL plane were equal to 3600 MPa and 4500 MPa, respectively. In front of those theoretical values, the E values determined experimentally via DIC at the WL and 4 mm away from the WL plane were approximately 3730 MPa and 4400 MPa, respectively. These results indicate that, for the material used in this study, the decrease of the Young's modulus in the direction of the WL is related to the FO gradients presented in the Fig. 13.

Volume fraction

The changes in VF at the flowing WL were very similar, independently of the insert type, position and thickness. To eliminate the uncertainty of the VF analysis at the boundaries of the scan, only a region of 8 mm was considered. Results in Fig. 15 show that, in the shell layer and more notably in the core layer, the VF decreased as the WL plane was approached. One would expect an increase of VF at the core layer in the vicinity of the WL plane due to the higher FO in flow direction, which would facilitate the fiber compaction. This decrease in VF can be explained by the change in flow front velocity at the WL interface. As mentioned in Sec. 3.2, a WL acts as a wall, especially at the core layer. In fact, lower VFs near the cavity walls have been reported in the literature, where the phenomenon was related to high shear rates and fiber migration [START_REF] Vélez-García | Unambiguous orientation in short fiber composites over small sampling area in a center-gated disk[END_REF][START_REF] Tseng | Progress on fiber concentration for injection molding simulation of fiber reinforced thermoplastics[END_REF]. The same behavior can explain the decrease of fiber content at the WL surface.

Conclusions

A flowing WL acted similar to a mold wall in terms of the induced FO. The largest influence on FO was located at the core layer of the WL in the injection-molded part. A flowing WL induced a large FO variation in transversal direction that reached a maximum at the WL plane. The width of the FO gradient (variation in transversal direction) induced by a flowing WL can be larger than 8-10 mm. The FO gradient slowly decreased along the flow distance, which explains the increase of failure strain and failure stress along the WL. At the flowing WL, there was an increase of fibers oriented in the thickness direction, which modified the shrinkage behavior at the WL and can lead to an increase of the part thickness at the WL. The insert shape did not influence the local FO at the WL, but had an impact on the FO gradient transversally to the WL plane. A cylindrical insert (diameter of 30 mm) induced a larger FO gradient than a slit insert (width of 1 mm). The flowing WLs in thinner parts exhibited shorter flow distances for vanishing the FO gradient induced by the WL. The induced FO gradient at the intermediate and shell layers disappeared with the flow distance, but the FO gradients at core and skin layers were persistent, even after 90 mm of flow distance.

Local strain analysis assisted by DIC during a quasi-static tensile test proved that the FO gradient at the flowing WL was responsible for the strong strain localization in the vicinities of the WL plane, which led finally to the failure of the specimen.

Fig. 1 .

 1 Fig. 1. a-) Plate dimensions of the flowing WL mold cavity. The positions of the X-ray computed tomography scans are marked in red. b-) Dimensions of the inserts used in our study for inducing a flowing WL.

Fig. 2 .

 2 Fig. 2. a-) Visualization of the FO analysis (component a 11 ) in a scanned region via X-ray computed tomography. b-) CT-based image of a section at core (z = 1.5 mm) that shows typical fiber shape and fiber-fiber distance.

Fig. 3 .

 3 Fig. 3. Evolution of the tensile mechanical properties of the WL along the flow direction (x) for plates with different insert shapes (cylindrical and slit) and different thicknesses (1.5 mm and 3 mm). a-) Engineering failure stress and b-) engineering failure strain.

Fig. 4 .

 4 Fig. 4. a-) FO averaged over thickness and flow direction for a flowing WL induced by a slit insert at position 2 (WL at 0 mm). b-) Strain field in loading direction (y) measured with DIC just before failure of the same specimen loaded transversally to the WL.

Fig. 5 .

 5 Fig. 5. FO gradient around the WL in the 4 characteristic layers across the thickness for a plate with slit insert at position 1 (3 mm to the left and 1.5 mm to the right): a-) Skin layer, b-) shell layer, c-) intermediate layer and d-) core layer.

Fig. 6 .

 6 Fig. 6. Thickness variations at the WL area induced in a-) a 1.5 mm plate with slit insert at position 1, b-) a 3 mm plate with a slit insert at position 1 and c-) a 3 mm plate with cylindrical insert.

Fig. 7 .

 7 Fig. 7. Influence of the plate thickness on the averaged FO at a flowing WL induced by a slit insert at position 1.

Fig. 8 .

 8 Fig. 8. Averaged FO around the WL region (between y = -2 mm and y = 2 mm) as a function of the normalized thickness at position 1 from a plate with a-) slit insert and b-) cylindrical insert.

Fig. 9 .

 9 Fig. 9. FO gradient around the WL at different distances from the insert tip (a-0.5 mm and b-5 mm) for plates (3 mm on the left and 1.5 mm on the right) with slit and cylindrical insert.

Fig. 10 .

 10 Fig. 10. Evolution of the mean FO along the flow direction for plates (3 mm on the left and 1.5 mm on the right) with slit and cylindrical insert.

Fig. 11 .

 11 Fig. 11. FO gradient at 65 mm after the insert tip for 3 mm plates with slit and cylindrical insert.

Fig. 12 .

 12 Fig. 12. FO gradient around the WL at different distances from the insert tip in a-) 3 mm plate with slit insert, b-) 3 mm plate with cylindrical insert and c-) 1.5 mm plate with cylindrical insert.

Fig. 13 .

 13 Fig. 13. FO around the WL at different distances from the insert tip in the 4 characteristic layers across the thickness for 3 mm plates with slit insert (left) and 1.5 mm plates with cylindrical insert (right): a-) Skin layer, b-) shell layer, c-) intermediate layer and d-) core layer.

Fig.

  Fig.14bcorrespond to a stress level of ~60.5 MPa, which is close to the failure stress of the sample extracted from the plate with cylindrical insert. The difference of measured strains between the cylindrical and slit insert correlates with the results of FO in Fig.11. In fact, in the plates with a cylindrical insert the fibers away from the WL plane (2-3 mm from plane) were more oriented in the flow direction than their counterparts in the plates with a slit insert. This would lead to higher deformations when the sample was loaded transversally to the flow direction.The results from DIC were also used to estimate the local Young's modulus (E) at the WL region and at zones far away from the WL plane. The moduli were calculated by taking the difference of global stress corresponding to a change of local strain going from 0.15% to 0.8%. It was found that the local Young's modulus diminishes with proximity to the WL plane. For all flowing WL cases, E at the WL ranged between 3650 and 3900 MPa. For points in a section 4 mm away from the WL plane, E increased and ranged between 4500 and 5000 MPa. In order to check if the measured local Young's moduli are plausible, the local Young's moduli were estimated analytically by mean field homogenization using the FO measured via CT. The Mori-Tanaka model[START_REF] Mori | Average stress in the matrix and average elastic energy of materials with misfitting inclusions[END_REF] was chosen for the calculations. In terms of the polymer matrix, a Young's modulus of 2230 MPa and a Poisson's ratio of 0.4 were used. These values were obtained from tensile tests performed on injectionmolded samples of the non-reinforced PBT compound. For the E-glass fibers, a Young's modulus of 72 GPa, a Poisson's ratio of 0.22[START_REF] Kaiser | Modified mean-field formulations for the improved simulation of short fiber reinforced thermoplastics[END_REF] and an aspect ratio of 20 were set. The analytical calculations were done in two regions, one exactly at the WL plane and the other 4 mm away from the WL plane for the case of 3 mm plate at position 1 with a slit insert. A multilayer representative volume element was generated using the same layers and the FO tensors given in Fig.13. The glass fiber volumetric fraction was defined equal to 0.18 in all layers. In this framework, estimated theoretical E values at the WL and 4 mm away from the WL plane were equal to 3600 MPa and 4500 MPa, respectively. In front of those theoretical values, the E values determined experimentally via DIC at the WL and 4 mm away from the WL plane were approximately 3730 MPa and 4400 MPa, respectively. These results indicate that, for the material used in this study, the decrease of the Young's modulus in the direction of the WL is related to the FO gradients presented in the Fig.13.

Fig. 14 .

 14 Fig. 14. True strain profile (WL at 0 mm) measured by DIC at a global stress of 60 MPa a-) for 2 positions in the 3 mm plate with slit insert and b-) for specimens in position 3 from 3 mm plates with different inserts.

Fig. 15 .

 15 Fig. 15. Fiber volume fraction in a flowing WL induced by a slit insert in a 3 mm plate at position 1. Profiles of fiber VF are given for 4 characteristic layers across the thickness and as average through the whole thickness.

Table 2

 2 Thicknesses in z-direction of the identified FO layers at the WL zone.

		Skin	Shell	Intermediate	Core
	3.0 mm plate	0.2 mm	0.6 mm	0.6 mm	0.2 mm
	1.5 mm plate	0.215 mm	0.215 mm	0.215 mm	0.215 mm