# In Situ High Aspect Ratio Development of Phosphate Glass Fillers in Polybutylene Succinate-Based Composites

Simone Fulchini, Olivier Leclerc, Noëllie Ylla, Joshua U Otaigbe, Véronique Bounor-Legaré, René Fulchiron

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Simone Fulchini, Olivier Leclerc, Noëllie Ylla, Joshua U Otaigbe, Véronique Bounor-Legaré, et al. In Situ High Aspect Ratio Development of Phosphate Glass Fillers in Polybutylene Succinate-Based Composites. Polymer Composites, In press, 10.1002/pc.70684. hal-05403630

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Submitted on 8 Dec 2025

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Distributed under a Creative Commons Attribution 4.0 International License
In Situ high aspect ratio development of
phosphate glass fillers in polybutylene
succinate-based composites.

Simone Fulchiniª, Olivier Leclercª, Noëllie Yllaª, Joshua U. Otaigbeᵇ, Véronique Bounor-Legaréª,
René Fulchironª,\*

a) Universite Claude Bernard Lyon 1, INSA Lyon, Université Jean Monnet, CNRS, UMR 5223, Ingénierie des Matériaux Polymères, F-69622 Villeurbanne Cedex, France.

b) School of Polymer Science and Engineering, University of Southern Mississippi, Hattiesburg, MS 39406, USA.

<footnote>\*Corresponding author: rene.fulchiron@univ-lyon1.fr</footnote>
Reinforcing polymers by incorporating fibers is a common practice in the polymer materials industry. However, the processing stages are sometimes complicated because of the flowing difficulties of the melt polymer filled with solid fibers of which the fraction is, in fact, limited.

In this work, the feasibility of in situ creating a fibrillar phosphate glass (P-glass) reinforcing phase in a polybutylene succinate (PBS) during the processing stage is demonstrated. This is ensured using a P-glass of peculiarly low glass transition temperature. Thus, during the processing stages, in this case, injection molding and spinning, nodules are converted into fibrils or needles with a very high aspect ratio. However, particular caution must be taken to avoid the possible premature P-glass crystallization that prevents its easy deformation afterwards. In practice, it is shown that the previous blending of P-glass in PBS at ambient temperature in a powder form is much more efficient than melt blending to avoid this crystallization. As an important size issue, the possibility of producing yarns of PBS/P-glass composite is also shown. The mechanical behavior of the obtained injected parts and yarns clearly exhibits the stiffening effect of the P-glass filler, especially if the filler is of fibrillar shape. Nevertheless, whether on yarns or injected specimens, the presence of fillers unfortunately led to an important reduction of the deformation at break. To complete, the Halpin-Tsai model was successfully applied to describe the Young modulus using filler aspect ratio values consistent with the morphological observations.
**Highlights:**

1) Composites of polybutylene succinate filled with low Tg P-glass are developed

2) The composites were processed by injection molding and spinning

3) Avoiding P-glass crystallization, high aspect ratios of fillers are obtained

4) The Young modulus is controlled by the filler fraction and the aspect ratio

**Keywords:** Polybutylene Succinate, P-glass, Processing, Aspect Ratio
# 1. Introduction

The use of alumino-borosilicate E-glass fibers as reinforcement for thermoplastic polymers is very common, whether with long or short fibers. For conventional thermoplastic processing (extrusion, injection, etc.), short fibers are used for obvious flow reasons. However, when the filler content is high, the flow can become difficult and require high pressures or even make impossible the use of these types of processing techniques. Otherwise, it can lead to the breakage of the fibers, making them less effective for reinforcement because of a lower aspect ratio [1]-[3]. Other types of glass, namely Phosphate (P-glass), have been used as fillers [4]-[10] for different reasons. Their interaction with polar polymers is beneficial to their dispersion [11]-[13], they are potentially well bioabsorbable [14],[15], they can be efficient flame retardant [16] and they have interesting optical properties [16],[17]. Moreover, depending of their chemical structure, they can present a low glass transition temperature ($T_g$) [18]-[20] and a viscosity that drops notably for increasing temperatures above $T_g$ [21]-[23]. These last points are crucial for the development of the present work. Furthermore, by adjusting the composition of these P-glass, the value of $T_g$ can be easily tuned between 100°C and 300°C [24]-[26]. This work will explore the feasibility of using facile strain-induced methods to create ‘in situ’ polymer-based composites containing fibrillar inorganic phosphate glass (P-glass) as reinforcing phase. The main idea is that, if the inorganic P-glass phase has a sufficiently low glass transition temperature, initial particles will be in the liquid state at the processing temperature and will then be converted ‘in situ’ into fibers or elongated structures with a high aspect ratio directly by the highly deforming processing step. In that respect, a few works mention that P-glass filaments, or at least, oblong particles, can be generated in a polymer matrix. This has been reported for injection molding of polyphenylene sulfide [5] polyetherimide [6] and polyetheretherketone [7], thus with a high processing temperature
(330-350 °C) but also with a polylactic acid and a rubberized acrylonitrile butadiene styrene blend for additive manufacturing [27] with processing temperatures higher than 200 °C. In a previous study, it was shown that this can also be performed in a polyethylene matrix composite by simply applying a relatively moderate shearing at 130°C (shear rate of 20 s⁻¹) [28]. Therefore, after cooling, the material will be nothing but a polymer reinforced by inorganic glass fibers even though it has easily processed by means conventionally used for thermoplastics not or only lightly filled. This will be the major difference with traditional 'ex situ' alumino-borosilicate (E-glass fiber reinforced composites) for which the filler content is limited for processes like injection molding and excluded for spinning for example. In addition, the inorganic P-glass fiber phase and the interphase can be potentially tuned to optimize aspect ratios and interfacial interactions via specific chemical compatibilization reactions [29],[30]. Therefore, a detailed understanding of the role of these effects on the behavior of the materials is needed. As described in the literature [28]-[32], the chemistry and dynamics of organic-inorganic polymer blends and composites is a subject of intense study. Moreover, it is worthy to note that the proposed strain-induced strategy is different from the reported advances in sol-gel technology that targets to generate in situ the inorganic phase whether in solution [33],[34] or even in the polymer at the molten state [35]-[37].

In principle, any melt processable synthetic organic polymers could have been used to carry out the present study. The chosen polymer matrix is the polybutylene succinate (PBS) which is a competitor for polyethylene (PE) in terms of its characteristic temperatures but which can advantageously be synthesized from bio-based monomers. Its Tg of approximately -30°C is lower than ambient temperature and its melting point is very similar to that of PE (~115°C) [38]. Moreover, even if PBS exhibits a relatively low crystallization rate which can limit its application [39],[40], it allows a broad range of industrial processing techniques. Besides,
reported results demonstrate that, additionally to the stress transmitted via the matrix, the
morphology of the P-glass particles also depends on the interactions between the P-glass and
the polymer [11],[19],[41]. Thus, the polar polymer PBS with highly functional ester groups is
promising. Furthermore, PBS is completely biodegradable [42],[43] thereby providing a novel
approach for developing 100% biodegradable composites containing in situ generated
potentially biodegradable P-glass reinforcements. It can be mentioned that, in this context of
completely biodegradable polymer composites, some works were reported [14],[44] on the
efficiency of P-glass to enhance the mechanical properties of biodegradable PLA.

In this investigation, P-glass of low T<sub>g</sub> was synthetized and then ground to be blended by
different ways with PBS at different contents. The obtained materials were processed by
injection molding and spinning. The obtained samples were then characterized by mechanical
and morphological analyses.

**2. Experimental**

2.1. *Materials*

The used PBS in this work was the BioPBS™FZ91 supplied by Mitsubishi and produced from
polymerization of bio-based succinic acid and 1,4-butanediol. Its density is equal to 1.26 g.cm⁻³
and the Melt Flow Rate (MFR) is indicated to be 5 g/10 min at 190°C with a mass of 2.16 kg.

While a number of special water durable phosphate glass compositions reported in the
literature [4],[45] could have been used, the P-glass synthesized for this study was of the
targeted following molar composition: 50% of tin fluoride (SnF<sub>2</sub>), 20% of tin oxide (SnO) and
30 % of phosphorus pentoxide (P<sub>2</sub>O<sub>5</sub>). This composition leads to a P-glass with a very low T<sub>g</sub>
(around 120°C) [46],[47], a value consistent with the processing temperatures that will be
used for the material processing. It should be noted that the P₂O₅ is not directly available.

Therefore, it is obtained from the ammonium dihydrogen phosphate NH₄H₂PO₄ as the phosphorus precursor. The P-glass is thus obtained by reaction given in Equation (1):

$$2(1-X-Y) \text{NH}_4\text{H}_2\text{PO}_4 + X \text{SnO} + Y \text{SnF}_2 \rightarrow (\text{SnO})_x(\text{SnF}_2)_y(\text{P}_2\text{O}_5)_{(1-x-y)} + 3(1-X-Y) \text{H}_2\text{O} + 2(1-X-Y) \text{NH}_3 \quad (1)$$

All the reagents of high purity were provided by Sigma-Aldrich (NH₄H₂PO₄: ≥98.5%, SnO: 97%, SnF₂: 99%). Regarding the respective proportions, the components were thoroughly mixed together in a vitreous carbon crucible and subsequently converted into a homogeneous melt by heating the crucible contents in a furnace at 430°C for 30 minutes. The melt was then cooled and annealed at 130°C for 1.5 h. The obtained P-glass block was then removed from the crucible and grinded using a grinder of a Pulverisette 14 (from Fritsch) allowing the use of liquid nitrogen if needed. The obtained powder was then analyzed using a Mastersizer 3000 (from Malvern Panalytical) to determine the particle size distribution. For the sake of simplicity, only typical values obtained are given here: D50 = 48 µm, D10 = 10 µm and D90 = 216 µm.

## 2.2. Material Characterization

### 2.2.1. Size extrusion chromatography (SEC):

The molar mass distribution of PBS was measured, using a system SEC Shimadzu equipped with a multiple detection system of a multi-angle light scattering (TREOS from Wyatt/Waters), a generic refractometer of Shimadzu and a viscosimeter (ViscoStar from Wyatt/Waters). Three columns PLgel MIXED-C - 5 µm (Agilent) were used. The samples were dissolved at 1.5 g.L⁻¹ in chloroform. Approximately 100 µL were injected and the flow rate of eluent was of 1 mL.min⁻¹.
### 2.2.2. Differential scanning calorimetry (DSC):

DSC experiments were performed under nitrogen flow using a Perkin Elmer Diamond DSC apparatus, daily calibrated both in temperature and heat flow with indium. Sample mass was in the range 10 to 15 mg. For all experiments, two heating runs at 10°C.min⁻¹ were carried out, separated by a cooling at 10°C.min⁻¹. For P-glass, the starting and the final temperatures were respectively 50°C and 250°C. For PBS the first heating was between ambient temperature and 190°C. The cooling was then down to -60°C and the second heating was again up to 190°C.

## 2.3. Blending methods

The composite materials were prepared from the desired P-glass powder mass fractions. Moreover, since the volume fraction of filler is often relevant when analyzing the mechanical properties of composites, it may be useful here to recall the relationship between mass and volume fractions (Equation (2)):

$$w_g = \frac{\emptyset_g \times \rho_g}{\emptyset_g \times \rho_g + (1 - \emptyset_g) \times \rho_p} \quad (2)$$

where $w_g$ and $\emptyset_g$ are respectively the mass and the volume fractions of P-glass, $\rho_g$ and $\rho_p$ are respectively the densities of the P-glass (3.75 g.cm⁻³, [27]) and the PBS (1.26 g.cm⁻³).

### 2.3.1. Internal Mixer

The formulations were prepared using a 76 cm³ batch internal mixer, Haake Rheomix OS, equipped with counter-rotating Banbury rotors. Depending on the wanted P-glass fractions the materials quantities were first weighed for a total of around 60 g. First, the PBS was introduced in the mixer where the temperature was set to 160 °C and the rotors speed was 50 rpm. After 3 minutes the P-glass was introduced. To ensure a good dispersion of the P-glass,
4 minutes after its introduction, the rotation speed was set at 100 rpm and the temperature
at 170°C for a further five minutes before the mixer was stopped and the material collected.

2.3.2. Powder blending

Another way of blending was investigated with the aim of reducing the high-temperature
processing time of P-glass (for reasons detailed further). Therefore, the materials were
prepared by mixing PBS and P-glass in powder state, enabling the blends to be used directly
in subsequent processing techniques. The pellets of PBS were cryogenically ground into
powders in liquid nitrogen using the grinder (Pulverisette 14). The size of the obtained
particles was around 80 µm. Before mixing, the P-glass and the polymer matrix were dried in
vacuum oven for 5h at a temperature of around 60°C. The dry blends were obtained thanks
to a speed mixer apparatus (Hauschild SpeedMixer® DAC 150.1 FVZ-K) with the required
quantities of PBS and P-glass powders.

2.3.3. Rheological behavior:

In order to verify material stability during the different processing steps, some rheological
measurements were performed using a rotational rheometer ARES from TA instruments with
a parallel plate geometry of 25 mm diameter and a gap of around 1 mm. First, frequency
sweeps experiments were carried out at 160 °C between 100 and 0.1 rad.s⁻¹ (in decreasing
order) with a strain amplitude of 10 %, on initial PBS and PBS with different P-glass contents
(blends obtained with the internal mixer). A frequency sweep measurement was also
performed on the P-glass alone (with a strain amplitude of 0.5%) but the material appeared
to evolve during the measurement. Thus, different rheological measurements for P-glass were
carried out in transient mode that takes much less time. These measurements consist in
recording the shear stress over time when a constant shear rate is applied (γ̇ = 0.1 s⁻¹). From
these experiments, the stress growth viscosity function η⁺(t) is obtained. They were
conducted after cumulative treatments at 200 °C to analyze the material evolution. Other
tests were done in oscillatory mode to follow the complex viscosity change over time. These
tests were performed at different temperatures (130, 160 and 200 °C) with a frequency of
1 rad.s⁻¹ and a strain amplitude of 1 %. For these experiments, the diameter of the cell and the
gap were adapted to the measured viscosity range (respectively 13 or 25 mm and 1.7 to
2.3 mm).

## 2.4. Processing

2.4.1. *Injection Molding:*

In a first approach, the processing method studied was injection molding. The tests were
carried out using a Babyplast 6/10P mini-press (Cronoplast S. L., Spain) equipped with a single-
cavity mold for dog-bone shaped tensile specimens. Injections were carried out at a melt
temperature of 160°C. The injections of composites containing up to 20 vol% of P-glass were
carried out using blends obtained from the internal mixer or by powder mixing.

2.4.2. *Spinning:*

In a second approach, spinning tests were carried out using the haul-off system adapted to
the capillary rheometer (Rosand RH7 from Malvern). This capillary rheometer is equipped with
a 15 mm diameter cylinder (D_c). The tests were carried out for composites with 5 and 10 vol%
of P-glass. The spinning conditions were as follows: A 1 mm or 1.5 mm diameter die (D_d) was
used (1.5 for 10 vol% of P-glass). Different drawing ratios (DR) were tested. The drawing ratio
is the ratio between the winding speed (V_w) and the average velocity of the material in the die
(V_d), itself obtained by dividing the volume flow rate (Q) by the die cross-section. Thus, DR can
be varied by modifying either the piston speed ($V_p$), directly linked to $Q$, or $V_w$. This results in Equation (3):

$$DR = \frac{V_w}{V_d} = \frac{V_w \pi D_d^2}{Q 4} = \frac{V_w D_d^2}{V_p D_c^2} \qquad (3)$$

In addition, the final filament diameter can be estimated by Equation (4):

$$D_f = 2 \frac{\rho_{liq}}{\rho_{sol}} \left( \frac{Q}{\pi V_w} \right)^{1/2} \qquad (4)$$

where $\rho_{liq}$ and $\rho_{sol}$ are respectively the density of the liquid material at the die exit and the density of the solid filament that is winded (It is interesting to note that this diameter does not depend on the die diameter).

For these tests, the drawing ratios were generally between 4 and 12 with a few up to 30. In addition, as for the injection molding tests, the spinning experiments were carried out with materials prepared either with the internal mixer or by mixing the powders. However, for reasons that will be explained later, the temperature of the melt for the spinning experiments was set at 160°C in the former case and only 130°C in the latter.

## 2.5. Morphological Analysis

### 2.5.1. Wide angle X-ray diffraction (WAXD):

The presence of crystalline structures for the P-glass was investigated by WAXD. The powder structural study was performed in reflection mode (CuKα radiation, $\lambda$ = 1.5418 Å, 40 kV, 40 mA) using a BRUKER D8 Advance diffractometer. The diffraction patterns were obtained in the $2\theta$ range from 5° to 55° by steps of 0.02°.
### 2.5.2. Scanning electron microscopy (SEM):

The microstructure of blends was observed by SEM using either a FEI Quanta 250 microscope or a Zeiss Merlin Compact VP microscope, depending on the case, at an accelerating voltage of 10 kV. The powder of P-glass was directly observed after being spread on a standard support covered with a double-sided carbon tape. For injected parts, the observed areas were obtained by immersion in liquid nitrogen and breakage using pliers. For filaments, especially to obtain sections in the longitudinal plane (drawing direction), embedding in a cross-linkable epoxy resin was used prior to gingerly cut with an ultramicrotome. After this preparation, the composite samples were metallized with copper (10 nm).

## 2.6. Mechanical Characterization

The mechanical characterization of the obtained materials was carried out using a tensile testing machine of type Autograph AGSX (Shimadzu) equipped with a 10 kN capacity load cell. For injected samples of dog-bone shape (type: H2, with a cross-section of 4x2 mm² and an effective length of 20 mm), standard clamps were used and the crosshead velocity was set to 10 mm.min⁻¹. Five specimens of each type of injected parts were tested. For filaments obtained with the haul-off device of the capillary rheometer, a specific clamping system was used: the High Capacity Wire / Rope Grip G1092 (from Mark-10). For these measurements, the stretching velocity was 50 mm.min⁻¹. The tests were repeated 10 times for each type of filaments.
# 3. Results and Discussion

## 3.1. Specific properties of materials and their evolution during thermal treatment

The first part of the present study was to deeply characterize both the initial P-glass and the PBS and the influence of the processing parameters on some specific composites characteristics of the composites such as matrix molar mass, crystallinity or viscosity evolution.

First of all, a SEM picture of the obtained powder of P-glass after synthesis and grinding is given in Figure 1 showing good agreement with the size distribution results ($D_{50}$ = 48 µm).

<figure><img src="image_3.png" /><figcaption>Figure 1: SEM picture of P-glass after synthesis and grinding</figcaption></figure>

DSC analysis of both the synthesized P-glass and PBS are respectively shown in Figure 2 and Figure 3. The glass transition temperature $T_g$ of P-glass appears to be 108°C (midpoint). It should be noted that no change appears between 1st and 2nd heating, indicating that the P-glass synthesized according to the procedure described in paragraph 2.1, is relatively stable, at least from a chemical point of view.
<figure><img src="image_4.png" /><figcaption>Figure 2: DSC traces for P-glass (heating and cooling at 10°C.min⁻¹)</figcaption></figure>

PBS initially shows a main melting peak at around 119°C (with a first small peak at around 90°C). On cooling, an exothermic peak at around 68°C indicates crystallization. On the second heating from -60°C, the glass transition is first observed at around -36°C, as expected [48], followed by the melting peak at 115°C after a small exothermic peak at 97°C. The latter is clearly due to a cold crystallization or melting-recrystallization phenomenon.

<figure><img src="image_5.png" /><figcaption>Figure 3: DSC traces for PBS (heating and cooling at 10°C.min⁻¹)</figcaption></figure>
From these DSC measurements, the crystallinity is obtained using Equation (5):

$$X_c(\%) = \frac{\Delta H_m - |\Delta H_{cc}|}{\Delta H_{100\%}} \times 100 \quad (5)$$

where $\Delta H_m$ is the measured melting enthalpy, $\Delta H_{cc}$, the enthalpy due to possible cold crystallization during the heating run and $\Delta H_{100\%}$ the melting enthalpy for a 100% crystalline PBS that was taken to be 200 J.g⁻¹ [49].

Hence, from the results given in Figure 3, where $\Delta H_m - |\Delta H_{cc}|$ is 71.44 J.g⁻¹, the obtained crystallinity is 35.7% for this PBS.

The blends obtained by melt blending in internal mixer were first characterized in terms of rheology. The complex viscosity curves are shown in Figure 4, and compared with PBS 5 alone and with P-glass alone.

<figure><img src="image_6.png" /><figcaption>Figure 4: Complex viscosity versus frequency for PBS and P-glass alone and for blends obtained with the internal mixture (Measurement temperature 160 °C)</figcaption></figure>

The first point to notice in Figure 4 is that, for the 1 and 2 vol% P-glass mixtures, the viscosity is surprisingly lower than that of PBS alone. As the viscosity of the P-glass filler is the highest
one, this can only be explained by assuming that the molar mass of PBS in the presence of P-glass was reduced. In this regard, number (Mn) and weight (Mw) average molar masses obtained by SEC are given in Table 1 for the initial PBS, the PBS alone that was spun with the haul-off system and the PBS that was previously blended in the internal mixer before being spun. From these results, it can be concluded that the mixing operation leads to a reduction of the PBS molar mass, and this, more significantly than the spinning process alone.

<table><caption>Table 1: Molar masses of PBS before and after blending in internal mixer and spinning operation at 160 °C.</caption><thead><tr><th>Material</th><th>Mn (g.mol<sup>-1</sup>)</th><th>Mw (g.mol<sup>-1</sup>)</th><th>Mw/Mn</th></tr></thead><tbody><tr><td>Initial PBS</td><td>54760</td><td>118900</td><td>2.17</td></tr><tr><td>PBS after spinning at 160 °C</td><td>55100</td><td>113000</td><td>2.05</td></tr><tr><td>PBS / P-glass 5 vol% after spinning at 160 °C</td><td>48400</td><td>100300</td><td>2.07</td></tr></tbody></table>

The molar mass distribution of PBS, like generally for polyesters, is known to be affected by the melt processing. A decrease of the molar mass is usually encountered but its magnitude depends on different parameters as the previous drying process, the temperature [50], the melt-state time [51]-[53], and even the presence of inorganic fillers [54]. Nevertheless, depending on temperature and melt-state time, branching leading to Mw increases is also reported [50],[53].

From a more quantitative point of view, the obtained ratio between Mw for the PBS in the blend and Mw of the initial PBS is equal to 0.84. Recalling that this Mw ratio is raised to the power 3.4 for the Newtonian viscosity [55], this result appears consistent with the viscosity curves in Figure 4 for low P-glass fractions. For mixtures with a greater fraction of P-glass, the PBS degradation certainly also exists, but the effect is balanced by the presence of P-glass and its higher viscosity, leading to a higher viscosity than that of PBS alone.
With regard to the complex viscosity curve of the P-glass alone, it should be noted that the behavior shows a surprising variation with a decreasing viscosity as a function of the frequency up to 3 rad.s⁻¹ and then a stabilization. However, it should be remembered that the tests were carried out in the decreasing frequency order. Moreover, the lower the frequency is, the longer is the time necessary for the measurement. Generally, this type of results is explained by an evolution of the material during the experiment. Thus, additional rheological measurements have been performed to explain this evolution and quantify the phenomenon kinetics. First, successive transient stress growth experiments (fast measurements) were carried out after different treatment times at 200°C and the obtained stress growth viscosity functions (η⁺(t)) are plotted in Figure 5. From these experiments, the P-glass viscosity clearly changes with time, showing an increasing by a factor around 5 between 20 min and 30 min at 200 °C.

<figure><img src="image_7.png" /><figcaption>Figure 5: Stress growth viscosity functions for P-glass after different treatment times at 200 °C (Measurement temperature 200 °C).</figcaption></figure>

Additionally, time sweep experiments were performed at different temperatures (130, 160 and 200 °C) as shown in Figure 6. Again, the increase of P-glass viscosity is noted, but the
interesting aspect is that the kinetics of this evolution is much more pronounced at 200 °C compared to 130 °C.

Obviously, various reasons can be advanced to explain this increase of viscosity, but since the increase is faster when the temperature rises and no variation of $T_g$ was observed in DSC between the first and second heating, the phenomenon that appears most likely is the crystallization of P-glass. To verify this hypothesis, the samples used for the rheological time sweep experiments were rapidly cooled at the end of the experiments (2 h) and analyzed by WAXD and the results are shown in Figure 7. Clearly, the crystallization of P-glass is confirmed, especially at 200 °C but it has started even at 160 °C with the observation of few diffraction peaks. On the contrary, the rheological and WAXS analyses performed at 130°C confirmed that even after 2 h of thermal treatment, the P-glass remains amorphous.

<figure><img src="image_8.png" /><figcaption>Figure 6: Complex viscosity of P-glass versus time for different temperatures.</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>Figure 7: WAXD diffractograms of P-glass after treatments at different temperatures for 2 h. (For better clarity, the curves have been vertically shifted).</figcaption></figure>

However, the outstanding feature is that the kinetics of the P-glass rheological evolution due to its crystallization is much pronounced at 200 °C compared to 130°C. As the aim of this study is to provide a high aspect ratio to the dispersed P-glass particles, this phenomenon is of importance. Indeed, it is classically considered that the higher is the process temperature, the lower is the P-glass viscosity (as indicated by the initial viscosities in Figure 6) which should be beneficial to reach a high aspect ratio during process. However, since the viscosity of the P-glass increases a lot and faster when the temperature is higher due to its crystallization, while counterintuitive, it is much better to remain at a low temperature, but obviously, above the P-glass $T_g$, to avoid this evolution (in Figure 6, at 130 °C the initial viscosity is higher but it remains almost constant).

## 3.2. Morphology and mechanical properties of injected composites

Concerning the injection molding process first, examples of SEM pictures of slices of the parts cut in the flow direction are shown in Figure 8 for materials prepared either with the internal mixer or by powder blending.
<figure><img src="image_10.png" /><figcaption>Figure 8: SEM Pictures of injected tensile test pieces longitudinally sliced. a) PBS / P-glass 10 vol% prepared with the internal mixer, b) PBS / P-glass 10 vol% prepared by powder blending, c) PBS / P-glass 20 vol% prepared by powder blending.</figcaption></figure>

From these pictures, it is obvious that the blending method has a crucial effect on the composite morphology in the injected parts. Clearly, when the composite has been prepared with the internal mixer, that means at more than 160 °C for tens of minutes, the P-glass particles are very difficult to elongate. In that case, it is very likely that the P-glass has crystallized, at least sufficiently to prevent the deformation during injection molding. On the contrary, following the powder blending method avoids this preliminary crystallization allowing the injection flow to confer a high aspect ratio of the P-glass particles. Actually, in this case, the thermal treatment at 160 °C is only during the injection molding operation and therefore remains much less time in the crystallization temperature window.

In conjunction with these morphological analyses, the mechanical properties of the composite injected parts have been scrutinized. First, for a rapid comparison, Figure 9 shows exemplary tensile curves and Table 2 gathers the values of Young modulus, tensile strength and elongation at break, averaged on 10 trials.
<figure><img src="image_11.png" /><figcaption>Figure 9: Typical tensile curves of injected parts. a) Full scale, b) Zoom in on the early part of the curves.</figcaption></figure>

<table><caption>Table 2: Characteristic values obtained by tensile tests on the injected parts.</caption><thead><tr><th>Material</th><th>Young Modulus (MPa)</th><th>Tensile Strength (MPa)</th><th>Elongation at Break (%)</th></tr></thead><tbody><tr><td>PBS</td><td>284 ± 6.9</td><td>39.1 ± 0.77</td><td>403 ± 13</td></tr><tr><td>PBS (Internal Mixer)</td><td>344 ± 7.9</td><td>39.7 ± 0.67</td><td>233 ± 42</td></tr><tr><td>PBS / P-glass 5 vol%, Internal Mixer</td><td>352 ± 14.7</td><td>34.8 ± 0.56</td><td>197 ± 28</td></tr><tr><td>PBS / P-glass 10 vol%, Internal Mixer</td><td>422 ± 15.2</td><td>33.2 ± 1.65</td><td>42.4 ± 8.3</td></tr><tr><td>PBS / P-glass 15 vol%, Internal Mixer</td><td>474 ± 20.2</td><td>30.1 ± 1.89</td><td>23.3 ± 4.9</td></tr><tr><td>PBS / P-glass 20 vol%, Internal Mixer</td><td>434 ± 9.8</td><td>26.8 ± 0.31</td><td>36.4 ± 2.9</td></tr><tr><td>PBS / P-glass 10 vol%, Powder Blending</td><td>505 ± 14.0</td><td>35.6 ± 0.58</td><td>25.1 ± 1.3</td></tr><tr><td>PBS / P-glass 20 vol%, Powder Blending</td><td>614 ± 54.3</td><td>34.6 ± 1.06</td><td>20.1 ± 1.6</td></tr></tbody></table>

From these experiments, various rules of thumb can be driven. First, as expected, the Young modulus clearly increases with the P-glass fraction while, unfortunately but not surprisingly, the elongation at break is drastically reduced. Second, the P-glass is more effective to increase the Young modulus when the material was prepared by the powder blending route (but still with reducing the elongation at break). Such a reduction in fracture properties is generally linked to too weak bonds between the filler and the matrix, and this poor interfacial adhesion rapidly leads to cracking initiation. In the present work, this issue has not been tackled yet to improve these filler matrix interactions. In this regard, Figure 10 shows SEM pictures obtained on sections of injected parts before and especially after the tensile test. The very poor quality of the filler matrix interface is very clear and this point has to be improved in future works by, for example, implementing compatibilizer elements.
<figure><img src="image_12.png" /><figcaption>Figure 10: SEM Pictures of injected tensile test pieces of PBS / P-glass 5 vol% prepared with the internal mixer a) before tensile test, b) after tensile test.</figcaption></figure>

Nevertheless, for a deeper analysis of the mechanical behaviors, the obtained Young modulus values were then confronted with the Halpin-Tsai model that is very commonly considered for fibers reinforced composites. The Halpin-Tsai equations are derived from Hill's generalized self-consistent micromechanics model which assumed a composite in which the embedded phase is made of continuous and perfectly aligned cylindrical fibers without interactions between them. The extension of the model to short fibers composites led to semi- empirical equations needing to adjust a parameter for a variety of reinforcement geometries (of different aspect ratios and orientations)[56],[57]. Apart from assuming that the fibers are uniform in shape and dimension and that they are perfectly oriented, the model also considers an ideal bonding between both phases (no interfacial slip or micro cracking). Basically, considering that the fillers have a cylindrical shape and that they are oriented in the traction direction, the model is written in Equation (6) [58],[59]:

$$E_c = E_p \frac{(1+2pk\emptyset_g)}{1-k\emptyset_g} \quad \text{with} \quad k = \frac{\left(\frac{E_g}{E_p}-1\right)}{\left(\frac{E_g}{E_p}+2p\right)} \qquad (6)$$

where $E_c$, $E_p$ and $E_g$ are the Young moduli of the composite, the polymer matrix and the P-glass respectively, and $p$ is the aspect ratio, length over diameter, of the fillers ($p = 1$ for spherical particles).
Figure 11 shows the values of the measured Young modulus for various P-glass fractions and both blending methods (different experimental campaigns). The Halpin Tsai model predictions are also plotted for two aspect ratios (1 for spheres and et 3, for example, for small cylinders oriented in the tensile direction). For these predictions it was necessary to provide values of the Young modulus of PBS ($E_p$) and P-glass ($E_g$). The former (314 MPa) was taken as the average between the obtained measured values for $\emptyset_g = 0$. It must be mentioned that this obtained value agrees with values from the literature [60],[61]. For $E_g$, unfortunately, no data was available for exactly the same composition of P-glass. Therefore, a mean value of 40 GPa was taken from the literature concerning the phosphate glasses gathering various compositions ([62]-[64]). In any case, it must be mentioned that, since this modulus is much higher than this of the PBS, the predictions of the Halpin-Tsai equation is not very sensitive to this parameter in the considered volume fractions of P-glass.

<figure><img src="image_13.png" /><figcaption>Figure 11: Young modulus of the composites as a function of the P-glass fraction. Symbols: Experimental data of different experimental campaigns, solid lines: Halpin-Tsai predictions for two aspect ratio values.</figcaption></figure>
At this stage, the conclusion that can be drawn is that the Young modulus is very well predicted by the Halpin-Tsai equation considering an aspect ratio of 1 for the material prepared in the internal mixer and an aspect ratio of 3 for the materials prepared from the powder. This is fully consistent with the SEM observations revealing high aspect ratios for the dispersed P-glass when prepared from powders. However, it must be admitted that this prediction is less accurate for the P-glass fraction of 20 vol%.

### 3.3. Morphology and mechanical properties of filaments

Regarding now the results for the process that mainly involves elongation deformation, such as spinning, the SEM observations are shown in Figure 12 for filaments obtained with different compositions, different *DR* and prepared and spun differently. Again, these observations highlight the effect of the composite preparation method on the obtained morphology. Moreover, it must be specified that, in this case, not only the blending methods were different but also, in the case of blends prepared from powders, the spinning process was performed at only 130 °C to further reduce the risk of P-glass crystallization before being drawn along the thread line (it was 160 °C for the composites prepared with the internal mixer). From Figure 12 b), c) and d), it can be deduced that this strategy is successful given the very elongated shapes of P-glass phase obtained. It can be noted that for Figure 12 d), the filament was cut longitudinally but the cross-section is also distinguishable, revealing the real needle shape of the dispersed P-glass.
<figure><img src="image_14.png" /><figcaption>Figure 12: SEM Pictures of filaments longitudinally sliced. a) PBS / P-glass 5 vol% prepared with the internal mixer, *DR* = 8.8, b) PBS / P-glass 5 vol% prepared by powder blending, *DR* = 4.4, c) PBS / P-glass 10 vol% prepared by powder blending, *DR* = 30, d) PBS / P-glass 10 vol% prepared by powder blending, *DR* = 10.</figcaption></figure>

Concerning the mechanical characterization, the results analyzed below concern only the filaments obtained from the powder blending method. The effect of the P-glass volume fraction and of *DR* are discussed. First, raw results of tensile tests are shown in Figure 13 a) revealing very large behavior differences between the filaments. Again, these curves show that the elongation at break is reduced for the composites compared to the neat PBS. It must be admitted that, from Figure 12, determining the experimental aspect ratio of the dispersed P-glass phase for an input in the Halpin-Tsai modelling would not really be appropriate because of a very large uncertainty. That is why the model has been tested by adjusting the aspect ratio values. In Figure 13 b) the measured Young modulus is compared with the prediction of the Halpin-Tsai model showing that, for moderate *DR* (4.4), an aspect ratio of 4.5 is needed for well predicting the experimental value but that, for high *DR* (10 or 30) the
agreement between data and the model is obtained for an aspect ratio of 15. This evolution
of the adjusted aspect ratio is well correlated with the difference of morphology observed
between Figure 12 b) and Figure 12 c) and d). Additionally, it is shown that, for the high values
of DR, the measured Young modulus does not increase with DR anymore. Nevertheless, this
is not very surprising when comparing the needle shapes in Figure 12 c) and d) that are very
similar even if the DR values are different and this justifies the use of the same aspect ratio
value.

<figure><img src="image_15.png" /><figcaption>Figure 13: a) Typical tensile curves of filaments, b) Young modulus of the composites as a function of the P glass fraction. Symbols: Experimental data (Composites prepared by powder blending), solid lines: Halpin-Tsai predictions for 3 aspect ratio values.</figcaption></figure>

**4. Conclusion**

The feasibility of in situ creating a fibrillar reinforcing phase of phosphate glass (P-glass)
during the processing stage to obtain new PBS-based composite materials is shown in this
work. The concept is based on the particular low glass transition temperature of the selected
P-glass formulation. Thus, nodules can be converted into fibrils or needles with a very high
aspect ratio during material processing. However, for this to be possible, it is essential to avoid
premature P-glass crystallization and thus to control the thermal treatment. Otherwise, high-
temperature treatment will prevent the P-glass from being easily deformed afterwards. In
practice, this has been concluded by studying different methods of mixing P-glass with the PBS
matrix.

For example, when the blends were prepared by the melt route in the internal mixer
subsequent processing operations did not lead to the production of materials with elongated
P-glass fillers. However, as a first interesting result, yarns could be processed by spinning this
PBS/P-glass composite even if the P glass filler appeared here in the form of quasi-spherical
nodules. This was also observed after injection molding. On the other hand, when the
composites were prepared by grinding the PBS and P glass powders and mixing them at room
temperature, the morphologies of the yarns and of the injected parts showed very long P-
glass fibrils.

Following these processing experiments, the mechanical behavior was characterized for
injected parts and yarns. The stiffening effect of the P-glass filler is clearly demonstrated, and
to a much greater extent when the filler is in the form of fibrils. On the other hand, whether
on yarns or injected specimens, the presence of fillers unfortunately led to a very sharp
reduction in deformation at break.

The obtained moduli of the materials were successfully modeled using the Halpin-Tsai
equation that considers the filler aspect ratio of which the used values were consistent with
the morphological observations.

The next steps of these works would be to improve the mechanical deformability of the
composites. This would need to study and improve the P-glass/matrix interactions in order to
reduce the possibility of fracture initiation at the interface and increase the elongation at
break. Some preliminary experiments have shown that it was possible to graft (aminopropyl)-
triethoxysilane (APTS) on the P-glass or maleic anhydride to PBS in order to compatibilize both
materials.

Moreover, further works have to be carried out to analyze the biodegradability of such
composites. In this regard, if the PBS matrix itself appears suitable, a difficulty may arise from
the composition of the P-glass, which, it should be noted, was used for its low T<sub>g</sub>. At this stage,
this composition contains a rather large amount of tin (albeit in the form of a metal oxide,
which seems to be considered by the specialists of the domain, less harmful than in an
organometallic state).

Acknowledgements:

This work was supported by the region Auvergne-Rhône-Alpes (Project PhoBi, Pack
Ambition Recherche 2019). The authors also acknowledge the European program Erasmus +
for financial support to student exchange. The authors would like to thank the "Centre de
Diffractométrie Henri Longchambon" and the "Centre Technologique des Microstructures
(CTµ)" of University Claude Bernard Lyon 1 respectively for WAXS measurements, conducted
by Ruben Vera, and for SEM experiments, conducted by Pierre Alcouffe and Eloi Chamczyk.
The authors would also like to thank the company MDB Texinov®, in particular Ruddy Breleur
and Bruno Mougin, for supplying the PBS and for the fruitful scientific discussions. For the
purpose of Open Access, a CC-BY-4.0 public copyright license has been applied by the authors
to the present document and will be applied to all subsequent versions up to the Author
Accepted Manuscript arising from this submission
(https://creativecommons.org/licenses/by/4.0/).
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