3D printing strategies and mechanical performance
assessment of continuous basalt-fibre/PA12 composites
with high fibre volume fraction

P.L. Pichard, T. Le Gagne, L. Maheo, Justin Dirrenberger, U. Lafont, A. Le
Duigou

► To cite this version:

P.L. Pichard, T. Le Gagne, L. Maheo, Justin Dirrenberger, U. Lafont, et al.. 3D printing strategies and mechanical performance assessment of continuous basalt-fibre/PA12 composites with high fibre volume fraction. Polymer Testing, 2025, 150, pp.108936. 10.1016/j.polymertesting.2025.108936. hal-05279690

HAL Id: hal-05279690

https://hal.science/hal-05279690v1

Submitted on 16 Oct 2025

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# 3D printing strategies and mechanical performance assessment of continuous basalt-fibre/PA12 composites with high fibre volume fraction

P.L. Pichard<sup>a,∗</sup> ID, T. Le Gagne<sup>a</sup>, L. Maheo<sup>a,b</sup>, J. Dirrenberger<sup>c,d</sup>, U. Lafont<sup>e</sup>, A. Le Duigou<sup>a,d</sup>

## ARTICLE INFO

**Keywords:**

Continuous filament fabrication
PA12
Basalt fibre
Additive manufacturing

## ABSTRACT

The continuous filament fabrication (CFF), based on the fused filament fabrication (FFF), is a technology enabling to produce 3D-printed composite parts with high mechanical performance. Basalt fibres remain rather unused in CFF among fibrous reinforcements, in spite of properties similar to glass fibres. In addition, in the context of in-situ resource utilization for out-of-earth manufacturing, basalt fibres stand as an excellent earth-based analogue material of lunar regolith.

This study aims at bringing deeper insights on the effect of specific printing strategies, by tuning the slicing parameters and by using a customized printer for CFF, as well as on the thermo-elastic properties of a novel high-performance composite material consisting of continuous basalt fibres embedded into a PA12 matrix, with a fibre volume fraction of 55%.

Thermal and morphological properties of the filament are first characterized. Then, the optimal slicing parameters – layer height (LH) and interfilament distance (ID) – are explored through the prism of process-induced porosity; a pore volume content as low as 2.15 % is reported for LH = 0.2 mm and ID = 1.15 mm, testifying to the high material quality and process handling. Finally, the in-plane mechanical properties of PA12-basalt are measured in the quasi-static regime, including longitudinal (0°) tension and compression, transverse (90°) tension, in-plane shear (±45°) and out-of-plane three-point bending. An extreme anisotropic behaviour is revealed. Competitive mechanical properties in the fibre direction are reported, comparable to similar glass-fibre reinforced composites. Conversely, those in the transverse direction are relatively moderate. The latter are attributed to the localized porosity induced by the juxtaposition of adjacent filaments for the interfilament distance selected.

## 1. Introduction

The continuous filament fabrication (CFF) technology is a derivative of the fused filament fabrication (FFF) technology, where the filament is made of a thermoplastic matrix reinforced by continuous fibres. The low mechanical properties of FFF parts remain indeed one of the weaknesses of this technology, and adding reinforcements, such as fillers, short fibres or even long fibres can greatly overcome it [1]. CFF thus combines the mechanical properties of composite laminates manufactured with more traditional techniques, with the low cost and design freedom allowed by additive manufacturing. By controlling the material path deposition, it is not only possible to take advantage of the anisotropic

behaviour of the fibres to manufacture parts adapted to definite loading cases [2], but also to develop novel mechanical metamaterials [3–5].

Several techniques were developed to efficiently produce fibre-reinforced polymer matrices. The initial in-nozzle impregnation, as its name indicates, consists of blending the polymer at the molten state and the continuous fibres directly in the nozzle of the printer [6]. The dual-nozzle printing system, where fibres and matrix are supplied simultaneously by two separate nozzles, arose from further research. The fibres are usually previously coated with a thermoplastic or a thermoset resin [7]. This is currently the only CFF technology for which printers are commercially available from brands such as MarkForged® and Anisoprint®. It is also possible to use a pre-impregnated filament

<footnote><sup>a</sup> IRDL, Université de Bretagne Sud, UMR CNRS 6027, BIONICS Group, 56100, Lorient, France</footnote>

<footnote><sup>b</sup> Saint-Cyr-Coëtquidan, CREC, 56380, Guer, France</footnote>

<footnote><sup>c</sup> PIMM, Arts et Métiers CNRS, Cnam, 75013, Paris, France</footnote>

<footnote><sup>d</sup> Institut Universitaire de France (IUF), Paris, France</footnote>

<footnote><sup>e</sup> ESA, ESTEC, TEC-MSE, 2201 AZ, Noordwijk, the Netherlands</footnote>

<footnote>* Corresponding author.
E-mail address: pierre-louis.pichard@univ-ubs.fr (P.L. Pichard).</footnote>
with a single nozzle. Although this requires an additional process step to produce the continuously-reinforced filament, the fibre impregnation is greatly improved as compared to the in-nozzle impregnation, and out-of-the-shelf FFF printers with minor modifications can be used [8]. Co-extrusion with an extruder by means of a dedicated nozzle, sometimes preceded by a first coating in solution, is usually the standard impregnation method for continuous fibre-reinforced thermoplastic filaments when using a single filament and a single-nozzle printer [9–11].

The range of polymers used in CFF is wide, from basic PLA to engineering thermoplastics such as acrylo-butadiene-styrene, polyamide and poly-ether-ether-ketone [12,13]. On the other hand, the diversity of reinforcements remains more limited. An overwhelming majority consists of carbon fibre [13–21], in a lesser extent glass fibre [21–23] and finally natural fibres [10]. Basalt fibres are also promising manmade candidates, possessing similar strength and stiffness as compared to glass fibres [24]. They particularly stand out in terms of temperature operating range (from –200 °C to +800 °C), chemical resistance to alkalis, wear resistance, resistance to UV and high-energy electromagnetic radiations, low moisture absorption, and are intrinsically tough [25–27]. Basalt fibres also represent a suitable analogue material for lunar regolith fibres, which could be employed in future extra-terrestrial construction [28,29]. In spite of these advantages, continuous basalt fibres remain sparsely employed in composites, mainly due to the poor fibre-matrix bond strength without any sizing agent [30]. The few examples of 3D-printed continuous basalt-fibre polymer composites are essentially processed with commercially-available dual-nozzle printers [12,31,32]. PLA-basalt filaments were however produced in a two-stage impregnation process, reaching 31 % of fibre volume fraction within the subsequently printed parts [9]. Zanelli et al. achieved a similar fibre content (34 %) with an Anisoprint® printer, unlike Zhang and Sun (around 5 %), using a MarkForged® printer but similar fibres and thermoplastic matrices [12,31]. A first tangible application of a TPU-continuous basalt fibre composite was considered, where the 3D-printed green body was further sintered and used as a tailored thermal blanket [33]. Nevertheless, 3D-printed composites with basalt fibre volume fractions above 35 % have not yet been tested, nor the influence of the slicing and printing parameters on the part morphology and the resulting mechanical properties.

The fine-tuning of the slicing parameters, i.e. all the parameters that are related to the geometrical properties of the filament and of the part, is shown to be essential to limit the process-induced porosity and to modulate the mechanical properties [10]. Porosity is undesirable for applications where the bulk material is employed, while others may authorize material deposition at the ply scale which includes voids of controlled dimensions. At a constant printing speed, the slicing parameters that dictate the filament deposition to build the structure are the interfilament distance (ID), the filament width (FW) and the layer height (LH). When ID is superior to FW, tailored voids are generated, which dimensions can be modulated by varying ID [10]. The fraction of the process-induced pores located between adjacent filaments can also be controlled by ID, which is then almost equal to FW. LH similarly influences the inter-layer porosity.

The following research aims at investigating the in-plane quasi-static mechanical properties of 3D-printed parts manufactured with a pre-impregnated composite filament containing a high volume fraction of basalt fibres, using a customized single-nozzle commercial 3D-printer. The morphological and intrinsic properties of the filament are first provided, before the effect of the slicing parameters and printing speed on the process-induced porosity are discussed. The anisotropy of the filament imposes multi-directional testing, hence the quasi-static mechanical properties of PA12-basalt printed parts are measured in longitudinal (0°) tension and compression, transverse (90°) tension and in-plane shear (±45°).

# 2. Materials and methods

## 2.1. Materials

The composite filament is supplied by Suprem SA, Switzerland. It consists of continuous basalt fibres embedded in a PA12 matrix and is referred to as PA12-basalt filament in the rest of the paper. PA12 is used in powder form for the filament production. Basalt fibres are sized at 1 % in mass. The fibre volume fraction announced by the supplier is 58 ± 5 %, with a circular section of 0.5 mm in diameter. The PA12 grade used for comparison is Fillamentum FX256, but the latter does not correspond to the exact same grade as that entering into the composition of PA12-basalt filament.

## 2.2. Sample manufacturing and conditioning

PA12 is printed with a Prusa MK3S+ with the cooling fan duct of the Prusa MK4S, while PA12-basalt is printed with a Prusa MK3S+ adapted for the continuous filament fabrication (CFF). Modifications are carried out following the work described in Refs. [19,34], which allows a much more accurate control over the filament deposition, and thus enables to reach a process-induced porosity as low as about 2 % for continuous carbon fibre-reinforced nylon, compared to around 10 % for off-the-shelf CFF printers [19]. An in-house-manufactured flat-tip brass nozzle with an internal diameter of 1.0 mm is installed, as well as a two-duct fan shroud to provide a more homogeneous cooling field around the nozzle than the default printing fan shroud. Each printer is placed in a thermally isolated enclosure, which is passively heated during printing and the printing volume reaches a threshold of 40–48 °C in steady-state operation.

The continuous nature of the filaments, and above all the absence of cutting device in the printing head, prevents the use of conventional FFF or even CFF slicers. The object must be designed in one single-stroke trajectory. As a result, the nozzle path is programmed directly thanks to a routine coded in a CAD software (Rhino7/Grasshopper, Robert McNeel & Associates). This routine also implements the manufacturing parameters and provides the Gcode script which is then sent to the 3D printer.

The filament spools are stored 24 h at 45 °C in the oven prior printing. The print duration per sample or batch of samples (up to 10 h in this work) is considered negligible regarding moisture absorption [35], the spools are kept at ambient conditions (temperature: 20–26 °C, relative humidity: 40–60 %). The enclosure is pre-heated and the print is launched when the printing volume temperature has reached a steady-state, which requires around 30 min. The printing parameters reported in Table 1 are strictly identical for all the samples produced with each type of material, with a notable exception for a PA12-basalt series printed at 6 mm s⁻¹ in the straight trajectories, in order to evaluate the influence of the printing speed with regards to the process-induced porosity. If any, support structures are carefully removed. A batch of samples refers to a printed part out of which individual samples are extracted via water jet cutting and then dried in a vacuum oven at 45 °C for at least 48 h. All the samples are finally

<table><caption>**Table 1**

Printing parameters used for PA12 and PA12-basalt.</caption><thead><tr><th>Materials</th><th>PA12</th><th>PA12-basalt</th></tr></thead><tbody><tr><td>Nozzle temperature (°C)</td><td>230</td><td>260</td></tr><tr><td>Bed temperature (°C)</td><td>90</td><td>90</td></tr><tr><td>Enclosure temperature (°C)</td><td>40–48</td><td>40–48</td></tr><tr><td>Printing speed (mm. s<sup>-1</sup>)</td><td>10</td><td>2.5 for straight trajectories<br>1.0 for acute-angle (90° and higher) curved trajectories</td></tr><tr><td>Cooling fan power (%)</td><td>100</td><td>25 for the three first layers<br>100 for the others</td></tr></tbody></table>
conditioned in a dedicated enclosure at 23 °C and 56 % RH prior mechanical testing.

## 2.3. Characterization of samples' morphology

Images acquired via optical microscopy (VHX-7000, Keyence) are used to measure the layer height (LH) and the filament width (FW) at the scale of a single-printed filament, while scanning electron microscopy (JSM-IT500HR, JEOL) imaging is performed in order to evaluate the morphology of the filament as retrieved from the spool and of UD0° laminates. All specimens are cut with a diamond saw (SECOTOM-1, Struers), cast in epoxy resin and polished prior image acquisition. Porosity measurements are performed by image analysis with ImageJ, using SEM micrographs of 10-filament-wide UD0° laminates printed with a constant LH of 0.2 mm and a constant width. It is assumed that the pore surface fraction measured is equal to the pore volume fraction in the composite. The single filaments are printed with varying programmed layer height (LH) in the range 0.125–0.400 mm, by step of 0.025 mm. The compression ratio C (Equation (1)) established by Le Duigou et al. [10] and the filament aspect ratio A (Equation (2)) are useful metrics to evaluate the ability of the filament to spread at a given pressure, temperature and speed:

$$C = \frac{LH_{\text{measured}}}{LH_{\text{programmed}}} \qquad \text{Equation 1}$$

$$A = \frac{ID_{\text{measured}}}{LH_{\text{measured}}} \qquad \text{Equation 2}$$

Finally, the influence of the slicing parameters on the morphology of resulting UD0° laminates is investigated. The varying parameters are the interfilament distance (ID = 1.10, 1.15, 1.20, 1.25 mm), the printing speeds (2.5 and 6 mm s⁻¹) and the number of layers (NL = 6 and 12). 5 specimens are printed for each ID and NL at 2.5 mm s⁻¹, and one at 6 mm s⁻¹. Cross-sections are then analysed with ImageJ, where the pore distribution and fraction are measured.

## 2.4. Characterization of thermal properties

TGA (TGA/DSC3+, Mettler Toledo) is carried out in air (80 mL min⁻¹), from 25 °C to 700 °C at 10 K min⁻¹, on 5–15 mg of PA12-basalt filament retrieved from the spool, placed in a cylindrical alumina crucible. The onset temperature and the basalt fibre mass fraction within the filament w<sub>f</sub> are measured.

DSC (DSC3+/700/3884, Mettler Toledo) is carried out on 5–15 mg of PA12 matrix, of PA12 Fillamentum FX256 filament, of PA12-basalt filament retrieved from the spool, and of 3D-printed PA12-basalt filament (printing speed of 2.5 mm s⁻¹, ID = 1.2 mm and LH = 0.2 mm). Samples are placed in 40 µL aluminium pierced-lid crucibles. All tests are carried out at a rate of 10 K min⁻¹ under static nitrogen, and consist of a heating scan, followed by a cooling and a second heating scan, between –30 °C and 210 °C. The glass transition temperature T<sub>g</sub>, the melting temperature T<sub>m</sub> and the crystallization temperature T<sub>c</sub> are determined from the obtained heat scans. The crystallinity ratio, in percent, is calculated from Equation (3), where ΔH<sub>m</sub> is the enthalpy of melting of the filament, in J.g⁻¹, measured as the area under the curve and delimited by the baseline; ΔH<sub>m</sub><sup>0</sup> the enthalpy of melting of a theoretical 100 % crystalline PA12, equal to 209.3 J g⁻¹ [36]; and w<sub>f</sub> the basalt mass fraction calculated from the TGA. Three replicates are performed for PA12-basalt before and after 3D-printing, while commercial PA12 powder and PA12 Fillamentum FX256 are tested only once.

$$\chi = \frac{\Delta H_m}{\Delta H_m^0 (1 - w_f)} \qquad \text{Equation 3}$$

TMA (TMA/SDTA841e, Mettler Toledo) is carried out under nitrogen atmosphere (50 mL min⁻¹), on three cubes of 7x7x7 mm³ per direction.

They are cut with a diamond wire saw from a larger UD0° 3D printed slab. In total, nine samples are prepared. The slicing parameters are ID = 1.2 mm and LH = 0.2 mm (see 3.2 Determination of the slicing parameters and evolution of the corresponding process-induced porosity). The printing parameters are stated in Table 1. The cubes are put in an oven at 50 °C for 3 days, and the dimensions are measured with a micrometer (IP65, Mitutoyo) right prior testing. The test starts at room temperature and consists of a cooling down to –30 °C, and then heating from –30 °C to 140 °C at 3 °C.min⁻¹, with a fixed compressive force of 0.05 N. The displacement of the sample is recorded as a function of the temperature; the coefficient of linear thermal expansion is therefore the slope of the line, according to Equation (4):

$$\bar{\alpha}_i = \frac{\Delta L_i}{\Delta T} \times \frac{1}{L_{0,i}} \qquad \text{Equation 4}$$

where α<sub>i</sub> is the coefficient of linear thermal expansion (CTE) in the direction i, ΔL<sub>i</sub> is the elongation in the direction i and L<sub>0,i</sub> is the initial length of the sample in the direction i; direction i being that in which the load is applied. The theoretical CTE for composites can be calculated by means of the rule of mixtures, following Equation (5).

$$\alpha_{c,i} = \alpha_{f,i} V_f + \alpha_{m,i} V_m \qquad \text{Equation 5}$$

where α<sub>c,i</sub> is the coefficient of linear thermal expansion of the composite in the direction i, α<sub>f,i</sub> and α<sub>m,i</sub> the CTE of the fibre and the matrix in the direction i, and V<sub>f</sub> and V<sub>m</sub> the volume fractions of fibre and matrix, respectively.

## 2.5. Characterization of the mechanical properties

Longitudinal test specimens are printed one after another, while transverse and ±45° samples are extracted from a larger plate following the procedure described in Section 2.2. In addition to the printing parameters (Table 1), the layer height and the interfilament distance are set to respectively 0.2 mm and 1.2 mm. The interfilament distance (ID) deliberately does not correspond to that minimizing the intralayer porosity (see Section 3.2 Determination of the slicing parameters and evolution of the corresponding process-induced porosity) as lower values could not experimentally provide sufficiently repeatable and reliable prints when a representative number of layers (more than 20) is used. In practice, adjacent filaments tend to locally overlap when using ID = 1.15 mm, which creates excessive thickness whose accumulation finally leads to part detachment from the printing bed, filament breakage at the nozzle tip or other kinds of print failure. As a consequence, the mechanical properties measured in this article, especially in transverse loading and in-plane shear, are not expected to be the highest achievable with this material and manufacturing technology. But they are representative of a large range of objects that can currently be manufactured without any issue with the customized printer used in this work. A minimum of three samples per test type and filament orientation are evaluated, their dimensions and the corresponding test parameters are reported in Table 2. Four end tabs per sample are glued with cyanoacrylate (Loctite 406, Henkel) on a previously sanded surface prepared with a polyolefin primer (Loctite SF 770, Henkel), except the ±45° samples which are glued with epoxy (Araldite 420 A/B, Huntsman)

<table><caption>**Table 2**
Dimensions and test parameters for each type of sample.</caption><thead><tr><th>Sample type</th><th>Sample dimensions (Length x width x height, mm³)</th><th>Crosshead displacement speed (mm.min⁻¹)</th></tr></thead><tbody><tr><td>UD0° in tension</td><td>160 × 15 x 0.8</td><td>2</td></tr><tr><td>UD90° in tension</td><td>120 × 20 x 3</td><td>1</td></tr><tr><td>±45° in shear</td><td>120 × 15 x 2.4</td><td>2</td></tr><tr><td>UD0° in compression</td><td>110 × 10 x 2</td><td>1</td></tr><tr><td>UD0° in bending</td><td>60 × 15 x 2</td><td>2</td></tr></tbody></table>
without applying any primer.

Longitudinal tensile tests are performed at room temperature with an Instron 6800 equipped with a 50 kN load cell, while a 10 kN load cell is used for the other tension, shear and compression tests. Three-point bending tests are performed at room temperature on a ZwickRoell Z250 with a 20 kN load cell. For this test, a pre-load of 5 N is applied and the test is stopped once 80 % of the maximal force recorded is reached. Strain measurement for UDO° samples in tension is handled with a bidirectional strain gauge placed on the surface printed on the bed; with a unidirectional strain gauge placed on the same surface for ±45° in-plane shear samples; and with one unidirectional strain gauge per sample side for UDO° samples in compression. Strain gauges are glued to the sample with Loctite 406. An extensometer records the strain for UD90° samples in tension. The flexural strain is calculated from the sample deflection, as indicated in the ISO 14125 standard. According to ISO 527-5, 14125 and 14126 the elastic moduli respectively in tension, three-point bending and compression are calculated between 0.5 % and 0.25 % strain, and between 0.1 % and 0.5 % strain in shear as stated by ISO 14129.

# 3. Results & discussions

## 3.1. Morphology and thermal properties of the filament

The filament exhibits a tan-like colour (Fig. 1A) and has a fairly constant section given the impregnation technique used during the filament manufacturing process. The overall dimensions are 0.41 ± 0.1 mm in the narrowest side and 0.58 ± 0.1 mm in its widest, which roughly corresponds to the 0.5 mm-diameter circular filament expected.

The average filament area determined via image analysis is 0.2080 ± 0.0046 mm². Basalt fibres are homogeneously distributed and dispersed within the PA12 matrix (Fig. 1A–B), although a slight matrix excess can be noticed in the edges of the filament. In general, the impregnation of basalt fibres appears to be good, with only a few areas where fibres remain dry, as highlighted in the red dashed rectangle in Fig. 1C. The pore volume content is 1.43 ± 0.90 %, which is in the range of other commercially available filaments [37]. However, some fractures and splinters within basalt fibres might have occurred during the sample preparation (cutting and polishing) and are evenly scattered over the entire filament cross-section (Fig. 1C and more detailed view on Fig. 1D, purple arrows). This fibre damage gives rise to very localized darkened areas throughout the entire cross-section, but the resulting potential overestimation of the calculated pore volume content is considered negligible.

Basalt fibres have diameters ranging between 7 and 18 µm, this variation originating from the multi-end roving used as continuous reinforcement. Basalt fibres account for 55.02 ± 4.23 % of the filament volume according to image analysis and the relatively high standard deviation reflects the inhomogeneity of the multi-end roving. The fibre volume fraction (FVF) is further confirmed by TGA measurements, at 55.60 % (76.53 % fibre mass fraction), as seen on Fig. S1. Both values from TGA and image analysis are in accordance with the technical datasheet provided by the manufacturer. The FVF of CFF-suited filaments typically ranges from 10 to 40 % and only exceeds 50 % on rare occasions, placing the current PA12-basalt filament with potential improved mechanical properties compared to literature [1,37,38]. The reference value regarding the basalt FVF for all subsequent calculations is 55 %.

<figure><img src="image_3.png" /><figcaption>**Fig. 1.** – Image of the PA12-basalt filament wound on the spool (A). Drawing of a dense printed part with several layers, defining the slicing parameters (B). SEM micrograph of a fraction of the filament cross-section. Basalt fibres appear in white, while PA12 is dark grey, contrasting with the epoxy resin in medium grey (C). Close-up on an impregnation-lacking zone, with fractured or split basalt fibres highlighted by purple arrows (D).</figcaption></figure>
The DSC and TGA analyses enable to set the thermal properties of the material, and therefore to accordingly define the possible range for the bed and nozzle temperatures. The measurements retrieved from the DSC scans of PA12 powder, PA12 Fillamentum FX256, PA12-basalt at the filament state and PA12-basalt once printed are reported in Table 3. The melting point and crystallization temperature of all samples correspond to those of PA12 [36]. The large variation in the glass transition temperature ($T_g$) and crystallinity ratio observed for the first scan are directly linked to the storage conditions and to the processing of the different materials, which are largely unknown for the raw materials. The thermal history of 3D-printed PA12-basalt is conversely better grasped. Here, the 3D-printing process involves a high cooling rate upon material deposition (maximal cooling fan power), and previous work showed that the higher the cooling rate the lower the crystallinity ratio [39]. However, as the printing environment is maintained in the same temperature range as that of the $T_g$ of PA12, the thermal printing conditions promote recrystallization [40]. This in turn results in a slight increase in glass transition temperature (see Figs. S2–S5 and Table S1). There is no evidence of cold crystallization in the first heating scan, indicating that PA12-basalt could reach its maximum crystallinity ratio during the 3D-printing process [39,40].

On the other hand, the second scan shows that the glass transition temperature of PA12 is around 50 °C, with a crystallinity ratio between above 25 %. The difference of $T_g$ is further confirmed by the TMA measurements (Fig. 2A). The lower crystallinity ratio and $T_g$ for PA12-basalt compared to PA12 do not suggest a specific nucleating effect of the basalt fibres, but rather indicates that they tend to delay crystallization, as already noticed for PLA/continuous basalt fibre composites [41]. These findings are supported by the reduction of crystallization enthalpy noticed on the DSC cooling scan of PA12-basalt compared to that of PA12 (Figs. S3 and S5). Although this behaviour appears to be in contradiction with the well-documented nucleating effect of short basalt fibres [42], especially with semi-crystalline polymers such as PA12, a reduction of this effect with the increase of the fibre length was already reported for basalt fibres, up to no noticeable effect for very high fibre aspect ratios [43,44].

The onset of degradation is measured at 270 °C according to ASTM E2550 (Fig. S1) and is attributed to the decomposition of the PA12 matrix. Basalt fibres remain conversely thermally stable up to at least 700 °C which is the maximum applied during the TGA analysis. In this way, the nozzle temperature shall be set between the melting point (180 °C) and the onset of degradation (270 °C), and this range corresponds to other PA12-based materials processing windows [31,45].

Fig. 2A shows representative smoothed curves of the coefficient of thermal expansion (CTE) in the transverse direction, for 3D-printed PA12 (blue) and 3D-printed PA12-basalt (tan). All raw data, for both materials and in the three directions, are displayed in Fig. S6. All CTE curves can be described in three different segments, starting with a fairly linear plateau between -20 °C and +20 °C. Then, a wide range between 20 °C and 65 °C corresponds to the transition from the glassy state to the rubber state, where the CTE is not constant. The glass transition temperature $T_g$ particularly stands out on the first heating scan (local maximum indicated by the red square), and coincides with the values found with the DSC method. This intermediate segment exhibits a large CTE increase, which is rather constant for the 1st cooling and 2nd heating scan, but much more oscillating during the 1st heating scan. It can be explained by the progressive release during the temperature rise of the residual stresses stored during the printing process, the first scan acting as an annealing step for PA12 and PA12-basalt [46]. The last segment begins after 65 °C, and is either a plateau during the cooling scan or a further marked increase during the heating phases. The heating scans after 65 °C superimpose each other (Figure S6A-D and F) except for the transverse CTE of PA12-12(Fig. S6E), for which a pronounced hysteresis is noticed. The reduced measured CTE during the first heating scan can be attributed to changes of stress state that, given the microstructure of the sample (Fig. 3A), might be amplified by presence of the intralayer pores and thus more visible than in other directions where porosity is negligible.

<table><caption>**Table 3**

Thermal properties of PA12 and PA12-basalt (filament and 3D-printed).</caption><thead><tr><th>Material</th><th>T<sub>g</sub> (°C), second scan</th><th>T<sub>c</sub> (°C)</th><th>T<sub>m</sub> (°C)</th><th>χ (%) , second scan</th></tr></thead><tbody><tr><td>PA12, powder</td><td>49.7</td><td>151.5</td><td>178.8</td><td>26.1</td></tr><tr><td>PA12, FX256</td><td>40.4</td><td>148.9</td><td>180.0</td><td>20.2</td></tr><tr><td>PA12-basalt, spool</td><td>44.85 ± 1.3</td><td>150.8</td><td>177.8 ± 0.4</td><td>19.5 ± 4.1</td></tr><tr><td>PA12-basalt, 3D-printed</td><td>41.3 ± 0.9</td><td>150.7 ± 0.5</td><td>178.3 ± 0.1</td><td>20.7 ± 2.2</td></tr></tbody></table>

As $T_g$ is passed during the test, CTE measurements have to be distinguished before and after this transition. CTE of 3D-printed PA12 are similar in all three directions (Fig. S6A-C), in line with previous measurements carried out on non-filled 3D-printed thermoplastic, which tend to be isotropic with regard to the CTE [47]. As a consequence, the average CTE in all three main directions is plotted in Fig. 2B (see Table S2 for explicit values): it almost exhibits a two-fold increase before and after the glass transition zone, respectively around 110 ppm °C⁻¹ and 200 ppm °C⁻¹, which are typical values for PA12 [48]. On the other hand, the presence of basalt fibres induces a high degree of anisotropy, with a constant CTE around 5 ppm °C⁻¹ along the fibre direction, corresponding to the CTE of basalt fibres [25]. CTEs in the transverse and out-of-plane directions are similar, and much higher than along the fibres, attesting the contribution of the PA12 matrix. Theoretical calculations using Equation (5) confirm the experimental results, giving respectively 52 ppm °C⁻¹ and 93 ppm °C⁻¹ before and after $T_g$.

With 55 % basalt fibres volume fraction, the investigated filament for additive manufacturing lies in the same range as composites manufactured with conventional techniques. On a thermal point of view, the presence of basalt fibres does not affect the key temperatures and thermal transitions compared to pure PA12, but conversely imparts an expected anisotropic behaviour with regard to the CTE, with an order of magnitude of difference between the longitudinal and the transverse/out-of-plane directions.

## 3.2. Determination of the slicing parameters and evolution of the corresponding process-induced porosity

When printing a new CFF filament, determining the slicing parameters is of primary importance, as this will condition the print path resolution and the range of attainable geometries. More precisely, the interfilament distance (ID) and the layer height (LH) are the parameters that can be modulated easily and almost independently of the part geometry (Fig. 1C). The filament width (FW) is dependent on the pressure applied by the nozzle, therefore indirectly linked to LH. Basalt fibres are under the form of a bundle that can easily be squeezed and spread, as opposed to the twisted flax fibre yarn which compressibility is limited [10].

There exists however an optimal layer height, at a given printing speed and nozzle temperature, for which the measured layer height is equal to the programmed layer height. Respecting this is important in order not to lose tolerance when stacking multiple layers. For that purpose, an initial LH scanning over the range 0.125–0.400 mm is proposed. Indeed, the filament dimensions are 0.41 × 0.58 mm, the layer height cannot exceed the smallest dimension. From an experimental point of view, it is neither possible to print below 0.125 mm (the filament breaks at the tip of the nozzle, due to excessive compression on the basalt fibres) nor above 0.375 mm (the filament systematically detaches from the printing bed). Consequently, the single filaments are printed in the readjusted range of 0.125–0.375 mm for the programmed layer height, and the resulting measured LH and compression ratio C (Equation (1)) are presented in Fig. 3B.

C is equal to 1 when programmed and experimental values converge,
<figure><img src="image_4.png" /><figcaption>**Fig. 2.** Representative curves of the transverse coefficient of thermal expansion (CTE) as a function of temperature for PA12 (blue) and PA12-basalt (tan) (A) – Mean CTE in the [-20; +20]°C interval (before Tg) and [65; 130]°C interval (after Tg), for three specimens and excluding the 1st heating scan (B).</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>**Fig. 3.** SEM micrograph of a UDO° specimen printed at 2.5 mm s⁻¹ with NL = 12, LH = 0.2 mm and ID = 1.2 mm, detailing the pore morphologies (A). Evolution of the printed layer height measured and the compression ratio, as a function of the programmed layer height. The optimal layer height, corresponding to a compression ratio equal to 1, is indicated by the black circle and arrows (B). Evolution of the process-induced porosity as a function of the interfilament distance (ID), the number of layer in the specimen and the printing speed (C). Numerical values are shown in Table S3.</figcaption></figure>

which occurs for LH = 0.200 mm, (Fig. 3B, black circle). Under this value, the compression ratio is over 1, meaning that the nozzle effectively exerts pressure on the filament, this is the ironing effect [17]. On the other hand, above 0.200 mm there is a lack of compression on the material deposited. It is important to keep in mind that these tests are made at the scale of a single filament, and thus that the effect of adjacent filaments is not taken into account. This is of great importance yet when

manufacturing large, dense parts, as filaments within the same layer restrict each other's ability to expand laterally [10]. It can also be observed on Fig. 3B that the layer height is constant every two measurements from 0.200 mm, reflecting the Z-axis resolution of the printer, which is indicated to be approximatively equal to 50 µm.

The filament aspect ratio A (Equation (2)) reflects its ability to spread out when subjected to the compressive force of the nozzle during
printing (Fig. S7). As expected, the filament width decreases when increasing the layer height – since the filament volume remains constant. An aspect ratio as high as 8.8 for LH = 0.125 mm means that the filament is squeezed down by the nozzle to have a filament width 8.8 larger than the layer height. Similar measurements conducted on PEKK-CF filament with the same fibre volume fraction printed with the same layer height demonstrate a slightly larger aspect ratio of 10.7–11.0 [38]. These findings suggest that the printing parameters (nozzle temperature, printing speed and fan cooling rate), by affecting the filament flowability and its solidification rate after deposition, also have a direct influence on the slicing parameters. On the other hand, the higher the layer height and the lower the filament aspect ratio, reaching a threshold around 4.5. This is generally the maximum value reachable by short-fibre reinforced filaments, whose aspect ratio varies between 2 and 4, depending on the layer height and the extrusion ratio [49]. However, this last parameter is not accessible for the CFF process, as the material deposition is calculated as a function of the filament length, and not of a material flow.

Finally, the compression ratio allows the correct layer height (LH) to be selected, enabling layers to be stacked smoothly and LH to be maintained over a large number of layers. The filament width (FW) depends on the applied LH, and the measurements at the single filament scale provide an estimation of the optimal interfilament distance (ID) for dense parts. Indeed, in theory when ID is equal to FW, the pore volume content should be minimum as filaments are exactly juxtaposed (as depicted in Fig. 1B). Reducing ID below FW results in filament overlapping and is not achievable for numerous layers; increasing ID conversely allows numerous design possibilities with tailored pore sizes. For LH = 0.200 mm, the measured FW at the scale of one printed filament is 1.31 ± 0.08 mm (Fig. S7). Nevertheless, this value is not exactly representative of the filament width within a larger part. The presence of adjacent filaments is expected to reduce the ability of the material to spread out upon printing and to reach a threshold from a dozen of adjacent filaments; in this case, setting ID lower to FW enables transverse filament compaction, enhancing the in-plane coalescence and molecular interdiffusion [10].

Porosity measurements for various interfilament distances (ID) slightly below 1.3 mm are therefore carried out, considering the influence of the number of layers (Fig. 3C). Logically, increasing ID increases the process-induced porosity in a linear trend, mainly driven by the intralayer pores (Fig. 3A). However, ID = 1.1 mm is experimentally unsustainable for parts thicker than around fifteen layers, the filament overlapping progressively leading to over-compaction. At a printing speed of 2.5 mm s⁻¹, the lowest porosity content is therefore reached for ID = 1.15 mm at 2.15 ± 0.24 % (Fig. 3C, red bars and Table S3). This is largely lower than most of the parts printed with the CFF technology, generally comprised between 7 and 11 % for commercial printers [18, 50]. Nonetheless, it lies in the same range as PA6i-CF printed with the same hardware [19], as PEKK-CF with the same fibre volume fraction printed without filament tension [38], but exceeds the remarkable 1.4 ± 0.2 % reported for PA12-basalt printed with a Markforged printer [31].

It can also be observed that the process-induced porosity for parts with 6 layers is always lower than that with 12 layers. Indeed, the first three/four layers are generally less porous than the others, as noticed in Fig. 3A. For the first layer, given the proximity with the heated printing bed, the solidification rate is sufficiently low to allow the polymer matrix to fill in all the gaps until complete coalescence and interdiffusion. However, the deposited PA12-basalt layers progressively insulate the top of the part from the printing bed, as both PA12 and basalt fibres have a poor thermal conductivity [51,52]. This results in a faster solidification rate and to pore formation. The porosity content at the layer scale then tends to become constant when stacking layers [10]. Moreover, rising the printing speed to 6 mm s⁻¹ leads to at least a two-fold increase in the measured porosity (Fig. 3C, yellow bars). With ID = 1.15 mm, it reaches 5.1 %. This behaviour can be attributed to a modification of the

temperature field and applied nozzle pressure on the filament during deposition, resulting in insufficient coalescence between adjacent filaments [53,54].

Three different pore morphologies can be noticed within the parts printed out of PA12-basalt, at 2.5 mm s⁻¹, with ID = 1.2 mm and NL = 12 (Fig. 3A). First, the intrafilament pores (blue arrows), in the range of 20 µm or lower, probably come from the initial pores present in the filament before printing, and remained trapped during the material deposition. They tend to have a circular or elliptical shape, and are essentially located in the middle of the filament cross-sections. Secondly, interlaminar pores are visible at the junction between two layers (orange-framed zones). They result from insufficient bonding when stacking the layers, and tend to have an elongated shape. Intrafilament and interlaminar pores account for a minor fraction of the pores present in the material, as noticed for other CFF-printed parts [55]. Most of them are intralayer pores (green-framed zones), which as their name indicates are located at the interface between adjacent filaments, and can run over several layers. As seen on Fig. S8A, they have an elongated shape and can be as long as 5 mm in the fibre direction. On the other hand, the typical dimensions in width and height do not exceed a few hundred micrometres (Fig. 3A). The shape of the intralayer pores is irregular and depends on the precision of the deposition trajectory, the nozzle pressure as well as on the local cooling kinetics. Their edges correspond to those of the printed filament, which are shown to present many partly debonded, loose and broken fibres. The proportion of material effectively in contact between adjacent filaments is thus considerably reduced compared to the out-of-plane direction, and even more to the longitudinal direction. The efficiency of the load transfer with such a discontinuous is therefore questionable, and it is sure that they act as major stress concentration areas, as reported for FFF-printed parts [56].

The major drawback of a such highly-filled filament is the very low printing speed of 2.5 mm s⁻¹, barely higher than the 1.7 mm s⁻¹ reported for a PEKK-CF filament with the same FVF [38], still yet around an order of magnitude lower than other CFF-printed parts [1]. This also involves a high rigidity, which restrains the range of attainable curvature radii and thus limits the manufacturing of low-resolution intricate objects [57]. The low fraction of matrix makes the bonding between adjacent layers difficult, hence a lower printing speed to ensure a complete melting of PA12 through the filament, as well as enough time for coalescence between adjacent layers. In case of dense parts, the partial re-melting of the already deposited filaments prompted by the nozzle within a single layer helps improving the intra-layer bonding in the transverse direction [58]. Increasing the printing speed yet rapidly yields to a drastic loss in print quality. It is therefore clear that the current technology, using a pre-impregnated filament with a single-nozzle printer – although claimed as the most efficient CFF technique for high-fibre volume fraction composites [19]–, is limited for some industrial applications where high production rate is of first importance. CFF should then be primarily considered for highly technical parts where intricate geometries are required. For larger-dimension parts, the upscaled version of CFF, automated fibre placement (AFP) technology, is more suitable, allowing a much higher material deposition rate, despite lower deposition resolution. Overall, setting the optimal ID for dense parts is a trade-off between the targeted lowest amount of process-induced pore, the printer resolution and the printability over a large range of geometries.

### 3.3. *In-plane mechanical properties*

Parts produced by additive manufacturing show an anisotropic mechanical behaviour, induced by the in-plane juxtaposition of material paths and by the layer stacking. For FFF, the coalescence of adjacent polymer filaments is driving the interfacial bond strength and therefore greatly impacts the resulting mechanical properties of the printed part [59]. They are generally much higher in the longitudinal direction, i.e. the filament deposition direction, than in the transverse and
out-of-plane directions where process-induced voids occur [59,60]. This
difference is further exacerbated in CFF-printed parts, as the fibre/ma-
trix interface introduces a supplementary order of anisotropy, and is
reported as preferential failure zones [19]. Zanelli et al. measured a
considerable 18-fold loss in strength in PA12-continuous basalt fibres
composites, between the longitudinal and the transverse directions, the
latter being even lower than that of the polymer matrix [31]. In longi-
tudinal tension, basalt fibres break progressively, in a scattered manner
and with extensive debonding, compared to carbon fibres which present
a more concentrated fracture surface [61].

Here, representative stress-strain curves of longitudinal tensile (UD0°), transverse tensile (UD90°), longitudinal compressive (UD0°), out-of-plane three-point bending (UD0°) and in-plane shear (±45°) are provided in Fig. 4A (see Figs. S9–S10 and Table S4 for exhaustive data). Under longitudinal tension, 3D-printed PA12-basalt behaves as a linear elastic and brittle material, with successive and progressive basalt fibre failures, until final part rupture, as indicated by the jagged signal on the marine blue curve of Fig. 4A. More precisely, each filament breaks one after another, resulting in the scattering of the filament throughout the sample thickness, as if it had exploded (Fig. 5A). Similar extensive damage appearance is observed for basalt-fibre/epoxy composites [62]. At the scale of the filament, basalt fibres are detached and dispersed, as already noticed for various basalt-based composites [61,62]. The tensile elastic modulus and the tensile strength, evaluated respectively at 38 GPa and 700 MPa, are in the same range as CFF-printed glass-fibre composites, and largely outperform the tensile properties of other CFF

basalt fibre-based materials [12,31,32]. PA12-basalt even rivals a few
carbon composites in terms of density-normalized tensile strength
(Fig. 4B).

It is worth noting yet that the printing process largely degrades this
property. While the elastic modulus remains similar for PA12-basalt at
both stages, 3D-printed UD0° composites show a 36 % reduction in
tensile strength compared to the as-received filament (700 MPa vs 1100
MPa), which stress-strain curves are provided in Fig. S12B. An even
more pronounced trend was already reported for carbon-fibre/
polyamide composites, presenting a drastic 60 % loss of tensile
strength after printing [63]. Passing through the extruder gear and the
heating elements is found to promote abrasion, induces some carbon
fibre breakage and lowers the strength of the entire composite, but most
of the damage occurs during the filament deposition, resulting from the
simultaneous compressive force exerted by the nozzle tip and bending
force generated by the 90°-angle at the nozzle outlet. In the present
work, fibre breakage is noticed both on top and bottom surfaces of
3D-printed UD0° laminates (Fig. S8A–B), but the identified damage re-
mains moderate. However, additional basalt fibres surely broke within
the PA12 coating, meaning that calcination or chemical dissolution
would be required to fully characterize the fibre damage extent induced
by the 3D-printing process [63]. Fig. S8A also displays numerous loose
fibres, and general fibre misalignments which is a typical defect of the
CFF process, especially in non-linear trajectories [64]. Finally, Fig. S8C
clearly shows abrasion marks on the edges of the broken fibre, attesting
that although the ironing effect provided by the nozzle may be beneficial

<figure><img src="image_6.png" /><figcaption>**Fig. 4.** Representative stress-strain curves for each of the mechanical tests performed. The three-point bending curve was truncated at the flexural strength, but complete curves are provided in Figure S10 (A); Performance diagrams of density-normalized strength versus density-normalized elastic modulus of various 3D-printed fibre-reinforced polyamides [12,14,16,18–20,22,23,31,65,66], in longitudinal tension (B), transverse tension (C) and in-plane shear (D). Composite density was either directly retrieved from the literature, or calculated from the given constituents density and fraction. Coloured areas provide visual guidance but are not representative of the exhaustive range of mechanical properties for each type of reinforcement.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>Fig. 5. Samples after mechanical testing - Top view of burst-like UD0°, tension (A) - Close-up of the sharp fracture surface of UD90°, tension (B) – Top view of ±45°, in-plane shear (C.1), close-ups of the fracture location (top view, C.2 and side view, C.3). Rupture occurs at the ply interface, but also within the plane between adjacent filaments – Localized buckling and partial delamination of UD0°, compression (top view, D.1 and side-view, D.2) – Localized buckling, fibre fracture and general delamination area of UD0° in three-point bending. (side-view, E.1, close-up on the damaged surface in contact with the three-point bending fixture, E.2, and close-up on localized buckling and fibre fracture on the top surface, E.3).</figcaption></figure>

to bonding between adjacent filament, the compressive strength of basalt fibres is locally and regularly exceeded, which is conversely detrimental to the mechanical properties in the longitudinal direction.

In three-point bending on UD0° samples, PA12-basalt shows an initial elastic behaviour until maximal strength is reached around 200 MPa at a strain of 1.3 %, and then fails immediately in a brittle manner, with losses of more than 50 MPa (Fig. S10). This initial failure is linked to delamination, originating from the area where the force is applied, and usually propagating along one of the sides of the sample (Fig. 5E1, black area in the lower side of the sample). The flexural strength and modulus are respectively equal to 200 MPa and 20 GPa, which are slightly lower than that for basalt-epoxy laminates [67], but are explained by the reduced shear strength of the interface between adjacent layers, given the additive manufacturing technique [68], and the use of a softer polymer matrix. After initial failure, the stress may reach a plateau, gradually decrease as the fibres break in a compressive manner on the upper surface (Fig. 5E2), or also drop abruptly in case more delamination propagates, in the same way as for basalt-epoxy laminates [67]. The fibre breakage also leads to localized delamination and debonding between adjacent filaments, as seen on Fig. 5E3.

As for longitudinal compression, the general behaviour is alike longitudinal tension, with an elastic linear phase followed by catastrophic failure, attributed to the formation of a localized buckling along with partial delamination (Fig. 5D1 and D2). The material is slightly stiffer in longitudinal compression than in longitudinal tension, but its strength is conversely around 3.5 times lower (200 MPa). A similar trend in behaviour and properties is observed within basalt fibre-epoxy composites but the failure mode reported is ply splitting [62].

Nevertheless, this impressive behaviour in the fibre direction gives way to much weaker properties in the transverse and in-plane directions. A 75-fold decay in the elastic modulus and an even more dramatic 580-fold decay in strength are observed for the transverse properties. The latter are even lower than those of PA12 itself. The in-plane shear properties are barely higher than transverse properties, but the mechanical behaviour is different (insert in Fig. 4A).

On the one hand, the 3D-printed composite fails in a brittle manner in transverse tension, at the interface between adjacent filaments, revealing a sharp fracture surface aligned in the fibre direction. Similar findings are reported for basalt-epoxy laminates, where matrix failure is reported in transverse tension, giving a flat rupture surface with a few debonded fibres; the strength is divided by around 26 compared to tensile loading in the fibre direction [62]. For 3D-printed PA12-basalt UD90° tested under transverse tension, the fracture surface morphology follows the irregular contours of the intralayer pores, and the whiter areas indicate matrix-rich areas (Fig. 5B and Fig. S11). The small amount of PA12 matrix is particularly striking, and it is clear that fracture originates from the intralayer pores previously described, confirming their detrimental role at providing sufficient transverse mechanical properties. In addition, the fracture surface presents apparent and numerous scattered exposed basalt fibres. Such a fibre debonding and dispersion therefore suggests a poor PA12-basalt fibre bond strength, which means that the homogeneous fibre impregnation observed in the filament cross-section (Fig. 1C) does not compulsory ensure a good load transfer. Due to their chemical inertness and smooth surface, it is notoriously difficult to have a strong interface between basalt fibres and most of the existing polymers [26,69].

On the other hand, in-plane shear is characterized by an elastoplastic behaviour, indicating PA12 matrix shear. Failure also occurs at the junction between two adjacent filaments in a similar manner as for transverse tensile loading, but the damage zone extends further to other filament interfaces, as seen in Fig. 5C1 and C.2. No fibre breakage is observed. However, partial delamination is reported, resulting in a staggered fracture profile when observed from the sample side (Fig. 5C3), given the alternation of plies oriented at +45° and -45° from the test direction.

To conclude with the in-plane mechanical properties, PA12-basalt present remarkable strength and modulus in the fibre direction that are comparable to other continuous glass-fibre reinforced thermoplastics in tensile, compressive and flexural loadings. Nevertheless, properties collapse when testing is performed in another direction, the worst
case being transverse tension, and the strength is then lower than the other PA12-basalt filament reported in literature [31](Fig. 4B and C). It is expected that slightly shifting the interfilament distance down to 1.15 mm instead of the selected 1.2 mm would reduce the volume of intralayer pores identified as the major weakness and cause of failure within the 3D-printed parts in the transverse direction as well as in-plane shear. Fig. S13 shows the almost total disappearance of intralayer pores, which is beneficial at reducing the stress concentration areas.

More investigation is then required to obtain a better trade-off between printability and mechanical properties in case the latter are of major importance in the transverse direction. Several solutions, already been explored in literature, exist to reduce the number of pores. The first lever is the improvement of the fibre-matrix interfacial bond strength, as basalt fibres by their chemical inertness and smooth surface are difficult to embed in polymer matrices [70]. Chemical additives such as sizing agents, or plasma surface treatment, could respectively address these drawbacks [11,30]. In addition, a finer tuning of the slicing parameters, especially selecting an interfilament distance of 1.15 mm, is recommended for future investigation aiming at lowering the intralayer pores, while verifying the printability for thick parts. Modifying the 3D-printer to enhance the filament bonding via better interdiffusion is also a possible solution. It may include an active thermal control in the enclosure, a larger nozzle to increase the ironing effect, or the addition of a heated roller or of an infrared heater in the nozzle vicinity [50]. Moreover, many post-processing steps could be performed, including annealing and hot pressing [71,72]. Regarding the very low printing speed, it is believed that the solutions listed above may offset the higher porosity content resulting from a higher printing speed while saving significant manufacturing time. Another area for improvement could address a better control on the accelerations, as the low printing speed is essentially due to curved trajectories. Finally, manufacturing laminates with various orientations, whenever possible, was also showed to partly offset the high anisotropy caused by the material and the process [55].

## 4. Conclusion

The present research provides in-depth insights regarding the characterization of a highly-filled PA12-basalt filament for the continuous filament fabrication (CFF), both on the thermal and quasi-static mechanical point of view. The main following conclusions can be drawn:

(1) All the thermal properties are shown to be influenced by the 3D-printing process. Compared to 3D-printed PA12, the presence of basalt fibres tends to reduce both the glass transition temperature and the crystallinity ratio, and imparts an anisotropic behaviour to the coefficient of thermal expansion, which is significantly lower along the fibres than in the transverse and out-of-plane directions.

(2) With 55 % volume fraction of continuous basalt fibres, this filament is one of the most reinforced ones currently available, all types of fibres combined. The printability is nevertheless adversely affected, with a printing speed for linear trajectories as slow as $2.5 \text{ mm s}^{-1}$ to ensure optimal layer bonding and minimal porosity, which leads to very long manufacturing duration, compared to FFF parts. Besides the printing parameters, the slicing parameters are of primary importance when manufacturing objects, especially in CFF. The optimal layer height is found at 0.2 mm, and the lowest pore volume content is measured at 2.15 %, for an interfilament distance (ID) of 1.15 mm, 12 layers, and printed at $2.5 \text{ mm s}^{-1}$. As general guidelines, reducing the printing speed and ID are beneficial to limit the porosity. Nevertheless, ID reaches a threshold under which adjacent filaments start overlapping, which jeopardizes the smooth printability of the part.

(3) Mechanical testing reveals a pronounced anisotropic behaviour, with excellent properties in longitudinal tension, similar to glass

fibres. The strength-to-density ratio is equal to $405 \text{ MPa cm}^3 \text{ g}^{-1}$ for PA12-basalt, which makes it compete with some carbon-fibre composites printed with the same technology ($340–735 \text{ MPa cm}^3 \text{ g}^{-1}$). Transverse properties are in return very limited, lower than the matrix itself, and are attributed to low matrix fraction and poor bonding quality – referring to insufficient fibre/matrix interfacial strength and extensive presence of intralayer pores.

The limitations regarding CFF and more generally additive manufacturing are also observed here, including anisotropy induced by the layer stacking, layer delamination, porosities induced by the process. Two major concerns are raised regarding the fibre-matrix and filament-filament interfaces, as they are shown to greatly affect the transverse properties of PA12-basalt. It is finally worth mentioning that some specific applications requiring high anisotropic properties could benefit from the findings presented in this study, such as sound absorption, energy dissipation or thermal insulators.

### CRediT authorship contribution statement

P.L. Pichard: Writing – review & editing, Writing – original draft, Methodology, Investigation, Formal analysis, Data curation. T. Le Gagne: Writing – review & editing, Investigation, Formal analysis, Data curation. L. Maheo: Writing – review & editing, Supervision, Methodology, Formal analysis, Conceptualization. J. Dirrenberger: Writing – review & editing, Supervision, Methodology, Formal analysis, Conceptualization. U. Lafont: Writing – review & editing, Supervision, Resources, Funding acquisition, Conceptualization. A. Le Duigou: Writing – review & editing, Supervision, Resources, Project administration, Methodology, Funding acquisition, Formal analysis, Conceptualization.

### Data availability statement

The data that support the present study are available from the corresponding author upon reasonable request.

### Author agreement statement

We the undersigned declare that this manuscript is original, has not been published before and is not currently being considered for publication elsewhere.

We confirm that the manuscript has been read and approved by all named authors and that there are no other persons who satisfied the criteria for authorship but are not listed. We further confirm that the order of authors listed in the manuscript has been approved by all of us.

We understand that the Corresponding Author is the sole contact for the Editorial process. He is responsible for communicating with the other authors about progress, submissions of revisions and final approval of proofs.

Pierre-Louis PICHARD on the behalf of the co-authors (Théo LE GAGNE, Laurent MAHEO, Justin DIRRENBERGER, Ugo LAFONT, Antoine LE DUIGOU).

### Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

### Acknowledgments

The authors wish to thank the Institut Carnot ARTS, the European Space Agency (contract number: 4000139402/22/NL/PA/pbe) and the Collège Doctoral de Bretagne (international mobility grant) for the financial support. Rodolphe HENRI from Suprem SA, Switzerland, is warmly thanked for the supply of the PA12-basalt filament. Bryan
LAUWERS, Hervé BELLEGOU, Antoine KERVOELEN, Isabelle PILLIN and Anthony MAGUERESSE from the Université de Bretagne Sud are also acknowledged for the technical support provided for the material characterization, as well as Julien DESCHAMPS and Quentin LE PAIH for the printing of the cubic samples for CTE measurements. Antoine LE DUIGOU and Justin DIRRENBERGER would also like to gratefully acknowledge Institut Universitaire de France (IUF) for financial support.

## Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.polymertesting.2025.108936.

## Data availability

Data will be made available on request.

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