Impact-modified polyamide 11/carbon nanotubes 3D printing and potential enhancement of toughness in fiber reinforced polymers

Abstract

Fused filament fabrication (FFF) of semi-crystalline polymers, such as polyamide 11 (PA11), is challenging due to significant shrinkage that occurs during cooling. This issue can be mitigated by blending the polymer with chemically compatible additives, which alter its rheological, thermal, and mechanical properties, thereby improving processability. In this study, we demonstrate the successful FFF printability of PA11 filaments modified with an ethylene–acrylic acid copolymer grafted with maleic anhydride (EAA-g-MAH) as an impact modifier. Carbon nanotubes (CNTs) were incorporated to compensate for the loss in terms of stiffness caused by the addition of EAA-g-MAH compound. Bulk compositions were prepared via melt compounding and compression molding and characterized, revealing that the addition of 20 wt% EAA-g-MAH increased the specific impact energy of PA11 by 282%. Two optimized impact-modified PA11 formulations—with and without CNTs—were subsequently extruded into filaments and, for the first time, successfully 3D printed through FFF. Mechanical testing demonstrated substantial improvements in ductility for the 3D-printed specimens, with increases of up to 56% for unfilled PA11 and 105% for CNT-reinforced PA11, relative to their bulk counterparts. This work opens new possibilities for directly printing tough polymers onto reinforcing fibers to create self-healing composites. Additionally, strong interfacial adhesion was achieved between glass fibers and the printed CNT-reinforced PA11, suggesting a potential increase in interlaminar toughness in fiber reinforced polymers (FRPs). These multifunctional materials processable via FFF represent a promising step toward the development of smart, self-healing structural composites.

Keywords

polyamide 11 (PA11)carbon nanotubes (CNTs)fused filament fabrication (FFF)toughness self-healing

Introduction

In recent years, additive manufacturing (AM), also known as 3D printing, has emerged as a remarkably versatile manufacturing technology, aiming to overcome the limitations of subtractive methods through rapid prototyping, improved product performance, and cost-effectiveness.^1,2 AM technologies have garnered attention, with applications primarily found in the aerospace,³ automotive,⁴ and medical fields.⁵ Among the various techniques, fused filament fabrication (FFF) has significantly contributed to the production of complex geometric parts, due to its straightforward processing and widespread accessibility.⁶ Patented by Stratasys in 1989,⁷ FFF entails the construction of a three-dimensional object, which is easily built up layer by layer guided by a computer-aided design (CAD) model.^8–10 Polymers and reinforced plastics are suitable materials for producing lightweight and cost-effective 3D printed products. Amorphous polymers like acrylonitrile butadiene styrene (ABS), thermoplastic polyurethane (TPU), and specific grades of polylactic acid (PLA) are among the most widely employed in this field. This is due to their relatively low softening point and thermal expansion coefficient, which make them suitable for this process.^11,12 However, the primary limitation of such polymers lies in the insufficient mechanical strength of the manufactured components, which is influenced by both material properties and printing parameters. A key approach to address this challenge is the use of high-performance thermoplastics, typically semi-crystalline polymers with high melting points, which exhibit enhanced mechanical properties.¹³ From an environmental perspective, another crucial consideration is the impact of the polymers used in FFF, but achieving the dual goals of high mechanical performance and environmental sustainability, particularly through the use of biopolymers, remains a significant challenge.¹⁴

Polyamide 11 (PA11) is a semi-crystalline polymer, characterized by high intermolecular forces and high crystallinity.¹⁵ Additionally, in contrast to most of the petroleum-based materials used for FFF, PA11 is biobased, as it can be synthesized from castor oil.¹⁶ Moreover, PA11 does not release toxic substances during printing, unlike styrene evolution during the printing of ABS.¹⁷ Due to these positive aspects, combined with good thermomechanical properties, ease of processability, high chemical and thermal resistance, and high toughness, PA11 has been widely used in various additive manufacturing techniques.¹⁸ For instance, powder bed fusion (PBF)¹⁹ and selective laser sintering (SLS)²⁰ enable the use of PA11, as the printing process is facilitated by laser interaction, which allows for the achievement of high temperatures in the molten region. Moreover, these techniques enable a controlled cooling process and uniform temperature distribution. This contributes to a more effective layer-layer bonding and higher mechanical properties in the final product.^21,22 In contrast, several issues arise during the deposition of PA11 layers in FFF, limiting print quality and dimensional stability. Printing complex parts with this technique is made difficult by the anisotropic shrinkage of PA11 upon cooling on the printing bed. This phenomenon is associated with the dense packing of the polymeric chains in the crystalline regions and most commonly results in warpage of the 3D printed part. Practical inter-layer welding in additive manufacturing is significantly hindered by the limited diffusion of the material prior to crystallization.²³ A mitigation of differential shrinkage is essential to strictly tailor the printing parameters, such as nozzle and bed temperatures, cooling rate, raster speed, and build orientation.²⁴ Despite these optimizations, warping phenomena that impede the printability of the polymers may persist due to the strong intermolecular forces in PA11.²⁵ A promising strategy for improving the printability of semicrystalline thermoplastic polymers and for enhancing their melt processing is to mix them with compatibilizers and additives.²⁶ This approach aims to increase chain mobility and reduce melt viscosity with the introduction of low molecular weight polymers.^27,28 Additives such as plasticizers and impact modifiers reduce the intermolecular interactions, primarily affecting the amorphous phase and generally preserving the crystalline structure.²⁹ In this way, the ductility of the resulting polymer is enhanced, with a slight reduction in elastic modulus and tensile strength only.^29,30 Thermoplastic polymers can be effectively mixed with elastomers, such as ethylene-propylene-diene monomer rubber (EPDM)^31,32 and ethylene acrylic copolymer modified with maleic anhydride (EAA-g-MAH),^33–35 to enhance impact resistance and plasticity. To counterbalance the reduction in stiffness that may result from the introduction of these impact modifiers, incorporating a filler capable of imparting increased stiffness could be a viable strategy.

A common strategy to improve the mechanical, electrical, and thermal performance of polymers employed in fused filament fabrication (FFF) is the incorporation of nanostructured fillers. Among these, carbon nanotubes (CNTs) are widely used in the production of polymer nanocomposites due to their ability to impart high strength, stiffness, and toughness.^36–38 CNTs are particularly versatile, as they can form a conductive network within otherwise insulating polymer matrices once the percolation threshold is exceeded. Moreover, owing to their high thermal and electrical conductivity, CNTs enable the development of polymer composites capable of localized self-heating, a feature of particular interest for thermally activated self-healing composite systems.^39–43

Generally speaking, self-healing materials emulate biological systems by autonomously repairing damage, operating through several mechanisms such as polymer chain diffusion, capsule rupture, or reversible bonding.⁴⁴ These processes are typically relevant to materials designed to endure significant static and cyclic loads, such as fiber-reinforced polymer composites.⁴⁵ Novel self-healing approaches are increasingly focused on using thermoplastic polymers as healing agents, leveraging their re-processability and thus allowing for multiple healing cycles.⁴⁶ A scalable and effective approach involves directly printing thermoplastic polymers onto the surface of the reinforcing fibers. Upon damage, the composite can be heated to a specified temperature, thus promoting healing through thermo-reversible entanglement of the thermoplastic polymer within the composite matrix.⁴⁷ Currently investigated thermoplastic matrices in this field include low melting point polymers like ethylene-co-methacrylic acid (EMAA)⁴⁸ or polycaprolactone (PCL).⁴⁹ High melting point polymers, such as PA11, can instead serve as healing agents in structural composite materials at elevated mechanical strength and service temperatures.

Based on these considerations, this study develops 3D-printable PA11-based filaments having elevated toughness, which can be 3D-printed directly on the reinforcing fabrics to produce structural composites. This work investigates for the first time PA11/EAA-g-MAH/CNT nanocomposites, examining their mechanical properties and 3D printability. To achieve this aim, different formulations were prepared by melt compounding and hot pressing, then the two most promising compositions were extruded into filaments suitable for FFF. The printability and the mechanical properties of the FFF-processed samples were evaluated and compared with those of the corresponding bulk materials. Moreover, the CNT-filled filament was 3D printed directly on glass fiber plies, and peeling tests were carried out, in view of the future development of epoxy/glass fiber composites with self-healing capability.⁴⁷ Hence, the novelty of this work lies in the first successful demonstration of polyamide 11 3D printability via FFF through coupling with EAA-g-MAH, which, with the additional contribution of CNTs potentially enables the toughening of the composites.

Materials and methods

Materials

Rilsan^® BESNO TL, an extrusion grade polyamide 11 (PA11) supplied in the form of granules, having a density of 1.02 g/cm³ and a melt flow index of 1 g/10 min at 235°C and 2.16 kg, was provided by Arkema Inc. (Colombes, France). The employed impact modifier was an ethylene acrylic copolymer modified with maleic anhydride (EAA-g-MAH), commercially known as Compoline^® CO/PA BA, hereinafter referred to as C. This product was supplied by Auserpolimeri Srl (Lucca, Italy) in the form of white solid granules and used as received. According to the technical datasheet, it had a density of 0.925 g/cm³ and a melt flow index of 2.5 g/10 min at 190°C and 2.16 kg. The selected multi-walled CNTs were NANOCYL^® NC7000, purchased from Nanocyl SA (Sambreville, Belgium). They were characterized by a mean diameter and length, respectively, of 9.5 nm and 1.5 µm, a surface area of 250–300 m²/g, and a density of 50 kg/m³. To evaluate the efficacy of 3D printing of impact-modified PA11 nanocomposites on reinforcing fiber plies in FRPs, 0°/90° glass fiber plies, supplied by Angeloni S. r. l. (Quarto d’Altino, Italy), were selected. These fibers presented a silane sizing suitable for coupling with epoxy resin, an areal density of 320 g/m², and a fiber diameter of 9.8 ± 0.4 µm.

Sample preparation

Bulk samples based on PA11/EAA-g-MAH/CNTs were produced via melt compounding and hot pressing. PA11, C granules, and CNTs were dried overnight at 80°C under vacuum and added at different relative amounts into a Thermo Haake^® Rheomix 600 melt compounder, where they were processed at 200°C for 5 min, with a rotational speed of the counter-rotating rotors of 60 rpm (C was added at minute 1, after complete melting of PA11). The resulting compounds were then ground and hot-pressed using a Carver^® hot plate press at 200°C for 5 min, applying a pressure of 3.4 MPa to a square mold of 120 × 120 × 2 mm³. The nominal composition of the prepared blends is reported in Table 1.

Table 1.

List of the prepared bulk samples with the relative amount of the constituents.

Sample	PA11 content^a[wt%]	EAA-g-MAH content^a[wt%]	CNT content^b[phr]
PA11	100	0	0
PA11_10C	90	10	0
PA11_20C	80	20	0
PA11_2.5CNT	100	0	2.5
PA11_10C_2.5CNT	90	10	2.5
PA11_20C_2.5CNT	80	20	2.5
PA11_5CNT	100	0	5
PA11_10C_5CNT	90	10	5
PA11_20C_5CNT	80	20	5

^aweight fractions of PA11 and C sum up to 100%.

^bphr = parts per hundred resin (grams every 100 g of total polymer mass (PA11 + EAA-g-MAH)).

A PRISM^® PTW 16/25D (Thermo Electron Polylab System Rheocord RC400P) corotating twin-screw extruder (TSE), having a screw diameter of 16 mm, a length-to-diameter ratio of 25:1, and a die diameter of 1 mm, was employed for the manufacturing of filaments suitable for the 3D printing process. Specifically, two compositions were selected from the preliminary analysis on bulk materials, according to their rheological, tensile, and impact properties; PA11_20C and PA11_15C_5CNT (i.e., the composite with 15 wt% of C and 5 phr of CNTs, see Paragraph 4.1). The extrusion was carried out on the previously compounded granules at a screw speed of 15 rpm. The optimized temperature profile in the extruder was set to 150°C, 215°C, 235°C, 230°C, and 220°C from the feed to the metering zone. The extrudate was cooled in water and wound onto a support, with a take-up speed of 6 m/min, to obtain a filament with a 1.75 mm diameter.

For the 3D printing process, a Creality^® Enders 3 printer for FFF was employed. To allow effective printing, the extruder was provided with a 1 kg load cell for bed leveling. After an optimization of the process parameters, a bed temperature of 100°C, a printing temperature of 255°C, and a print speed of 50 mm/s were utilized. To determine the tensile and impact properties of the 3D printed samples, 1BA dumbbell and standard 10 × 10 × 55 mm³ V-notched specimens were produced, utilizing a raster angle of ±45°.

Experimental techniques

Characterization of bulk samples

Rheological properties

Rheological measurements were conducted using an HR-2 Discovery^® Hybrid Rheometer (TA Instruments, New Castle, DE, USA) operating in parallel plate configuration (diameter of the plates = 25 mm, gap = 1.8 mm). Frequency sweep tests were performed at 200°C in air, at a frequency range from 0.05 to 600 rad/s, and applying a strain amplitude of 1%. Through these tests, it was possible to determine the trend of the apparent viscosity (η*) as a function of the shear rate. At least three specimens were tested for each composition.

Microstructural and chemical properties

Field emission scanning electron microscopy (FESEM) micrographs of the cryo-fractured surfaces of compression-molded samples were acquired using a Zeiss Supra 40 microscope, operating at an accelerating voltage of 2.5 kV. A platinum–palladium (80:20) conductive coating was sputtered on the specimens before the observation to make them electrically conductive. Fourier-transform infrared spectroscopy (FT-IR) was conducted in attenuated total reflectance (ATR) mode using a Spectrum Two FT-IR Spectrometer (PerkinElmer Inc., Shelton, Connecticut, USA), equipped with a Diamond crystal, in a wavenumber range of 450–4000 cm⁻¹. To reduce the signal-to-noise ratio, 20 scans were collected for each spectrum, utilizing a resolution of 4 cm⁻¹.

Thermal properties

Differential scanning calorimetry (DSC) tests were performed using a Mettler^® DSC30 instrument under a nitrogen flow of 100 mL/min, at a temperature range from 0 to 220°C, with a heating/cooling rate of 10°C/min. Each specimen was subjected to a first heating scan, a cooling scan, and a second heating scan. From the obtained thermograms, the thermal transitions of the prepared samples were measured, that is, the melting temperature and enthalpy (T_m, ΔH_m) of PA11, and its crystallization temperature and enthalpy (T_c, ΔH_c) values. Equation (1) illustrates how the degree of crystallinity (χ) of the PA11 phase was evaluated

χ = \frac{{Δ H}_{m} - {Δ H}_{c c}}{{Δ H}_{m}^{*} \cdot ω_{P A 11}} \cdot 100

(1)

where ΔH_cc is the enthalpy of cold crystallization of the PA11, ΔH^*_m is the enthalpy of melting of fully crystalline PA11, equal to 206 J/g,⁵⁰ and ω_PA11 is the weight fraction of PA11 in the samples.

Mechanical properties

Quasi-static tensile tests were performed at room temperature by using an Instron^® 5969 tensile testing machine (ITW Test & Measurement and Equipment, USA), equipped with a 1 kN load cell, testing ISO 527 1BA specimens, having a gauge length of 25 mm. The tests were performed at a crosshead speed of 10 mm/min, and at least 10 specimens were tested for each composition. From quasi-static tests, the ultimate tensile strength (σ_uts), and the stress and strain at break (σ_b, ε_b) were determined. To determine the elastic modulus (E), quasi-static tensile tests were conducted using the same machine equipped with an Instron^® 2620–601 extensometer, featuring a gauge length of 12.5 mm, at a crosshead speed of 0.25 mm/min. The secant modulus in a strain range between 0.05 % and 0.25 % was assessed. At least five specimens were tested for each formulation in this case.

Charpy impact tests were performed at ambient temperature using a Charpy pendulum CEAST^® 9050, following the ISO179 standard. The hammer possessed a mass of 2.5 kg and a length of 32.6 cm. An angle of 33° was used to obtain an impact velocity of 1 m/s, the span length was fixed at 62 mm. At least eight specimens were tested for each composition.

The mechanical properties, namely ultimate tensile strength (σ_uts) and Charpy impact energy, were modeled using Response Surface Methodology (RSM), to investigate their dependence on CNT and C content. RSM is a well-established technique in the literature and has proven to be a practical approach for identifying and quantifying the influence of multiple variables on a given response.⁵¹ A full quadratic response surface model was adopted to describe the influence of these factors on the selected mechanical properties. An analysis of variance (ANOVA) was subsequently conducted to assess the statistical significance of each term in the model. Factors with a high F-value and a corresponding p-value below 0.05 were identified as significant contributors to the response, while those with P-values greater than 0.05 were considered statistically non-significant and excluded from the final model.

Characterization of the 3D printed materials

Investigation of the 3D printability

Firstly, the 3D printability of the produced filaments was assessed by manufacturing 3 × 3 × 3 cm³ cubes, by using the printing parameters described in Paragraph 3.2. 3D printed specimens were then subjected to quasi-static tensile tests and Charpy impact tests, applying the same testing conditions described in Paragraph 3.3.1.4, and their properties were compared with those of the corresponding bulk samples. Then, the printability of the filament with composition PA11_15C_5CNT on glass fiber plies was also assessed by increasing the temperature to 285°C to favor the deposition, in view of its potential application as a healing agent in FRPs.

Peeling tests

A T-peel test was performed to assess interfacial adhesion between 3D-printed materials and glass fiber plies, following ASTM D1876 standard.⁵² Five specimens were prepared by 3D printing a polymer strip measuring 12 × 2 cm² directly onto a glass fiber strip positioned onto the print bed, using a nozzle temperature of 285°C and a bed temperature of 100°C. The polymer strip was printed with a longitudinal line infill parallel to the bond line detachment direction, at an infill density of 100%, to achieve a uniform polymer layer. A second glass fiber strip was then placed on top to create a three-layer assembly, which was melt-consolidated using a hot press at 1 kPa and 210°C for 10 min. These conditions match the manufacturing parameters of high service temperature epoxy/glass fiber composites, where this 3D printed material could potentially serve as a healing agent.⁴⁷ After consolidation, the specimens were gradually cooled inside the press to prevent anisotropic shrinkage, and then dried in a ventilated oven at 50°C before testing. During the T-peel test, the two glass fiber layers were gripped at opposite ends and pulled apart under tensile loading, causing the polymer interlayer to peel away from the fibers. The test was conducted at a speed of 100 mm/min. In this way, maximum and average peeling load values were determined.^52,53

Results and discussion

Characterization of bulk samples

Figures 1(a) and (b) report the results of the rheological measurements performed on the prepared bulk samples. As expected, both the impact modifier and the CNTs addition increase the viscosity compared to neat PA11, and this is particularly evident at low shear rate. It is possible to assess that the highest values of viscosity at low shear rates are registered with 5 phr of CNTs (Figure 1(a)), with a viscosity increase of up to two orders of magnitude compared to neat PA11, consistent with what is reported in the literature.⁵⁴ The EAA-g-MAH effect is comparable to that of CNTs, although it is slightly mitigated (Figure 1(b)). This can be explained by considering that the maleic anhydride groups present in the EAA-g-MAH can react with the PA11 terminal groups, thereby creating a more entangled polymer network and increasing the molecular weight.^55,56 Moreover, at elevated temperatures, the reactive groups in EAA-g-MAH can participate in post-condensation reactions, hence restricting molecular mobility and increasing melt viscosity.⁵⁷

Regarding the future application of these materials in 3D printing, viscosity trends should be investigated within a shear rate interval of 30-500 s⁻¹, corresponding to the shear rates typically applied during polymer extrusion from the nozzle. In this range, all specimens portray a decrease in viscosity, consistent with the pseudoplastic behavior of PA11. This is beneficial for the material processing, both during the filament extrusion and the 3D printing process. Viscosity values in a frequency range comprised between 0.01 and 0.1 s⁻¹ are more significant for understanding the rheological behavior of the material upon consolidation on the printing bed.^58,59 Notably, the more abrupt the increase in viscosity of PA11 at low frequencies is, the more severe the crystallization phenomenon will be.⁶⁰ Therefore, according to rheological results, the optimal composition to be printed should have the lowest value of viscosity at low shear rates. However, it must be noted that these results should be compared to those obtained from thermal and mechanical characterization to gain a comprehensive understanding and select the most suitable formulations for the intended application.

Figures 2(a) and (b) report the FESEM micrographs of some selected compositions to highlight the effect of both C and CNTs introduction. Considering the PA11/C blends, the SEM micrographs do not show canonical droplet-like morphology with C domains immersed in a PA11 matrix, typical of an immiscible blend. Instead, the observed homogeneous microstructure suggests good miscibility between C and PA11. This is in good agreement with the dramatic increase in viscosity at low shear rate observed in the rheological tests, indicating the occurrence of chemical reactions between PA11 and C during compounding. In the micrographs of PA11_5C and PA11_20C samples, some porosity is visible, most likely due to the entrapped air during compounding or to slight degradation phenomena caused by the high processing temperature. Moreover, the micrographs of the nano-filled samples reveal the presence of CNT clusters. The agglomeration of CNTs in a polymer matrix poses a significant problem, as it tends to act as a defect, impairing the mechanical properties of the resulting composites. The clustering of CNTs can be attributed to the high viscosity observed in the PA11/CNT nanocomposites, as previously reported in the literature.⁶¹

Figures 3(a) and(b) show the FTIR spectra of some prepared compositions and an inset highlighting the effect of C and CNTs on PA11. The characteristic signals of polyamide 11, that is, N–H stretching (3301 cm⁻¹), C = O (amide I) stretching (1635 cm⁻¹), and N–H (amide II) bending (1535 cm⁻¹), can be detected.⁶² No difference was observed in the FTIR spectra upon the addition of the CNTs, as shown in Figure 3(a). Notably, no shift of the peak at 3301 cm^-1 is shown in the FTIR spectra with varying CNT content, indicating that hydrogen bonding between CNTs and PA11 amide groups is unlikely or below the detection limit of FTIR.^63,64 This observation suggests that the interfacial interaction is likely dominated by weak physical forces, such as Van der Waals interactions, which may contribute to CNT stabilization and interfacial load transfer, as reported for similar PA-based nanocomposites.⁶⁵ According to literature,⁶⁶ the strength and stiffness of CNT–polymer composites are strongly influenced by the interfacial ordering of polymer chains around CNTs. Although this effect was not directly investigated here, it could play a role in reducing CNT aggregation during melt mixing.⁶⁶ Figure 3(b) highlights the effect given by the impact modifier. The presence of the maleic anhydride ring is detected in the region between 1650 and 1850 cm⁻¹, particularly at 1765 cm⁻¹, with an absorption peak associated with the C = O stretching.⁶⁷ The peaks within the 1724–1705 cm⁻¹ interval could be related to the C = O stretching of carboxylic acid groups, mostly pertinent to the maleic anhydride ring opening hydrolysis, which could indicate the occurrence of the reaction between PA11 and the EAA-g-MAH, schematized in Figure 3(c).⁶⁸ These results are consistent with the hypotheses drawn by observing the viscosity trends in rheological measurements.

The main results of the DSC measurements on the prepared bulk samples are shown in Figures 4(a)–(c) and in Table 2. All the DSC scans are reported, however, only the results related to the first heating and the cooling stage are shown, as they are the most significant ones. From the DSC thermograms, it is evident that the melting temperature (T_m) of PA11 is not substantially affected by either CNTs or C introduction. Concerning the impact modifier, a slight decrease in the glass transition temperature is noticeable upon its introduction, likely due to increased chain mobility compared to PA11. Moreover, C does not reduce the crystallinity of PA11, suggesting that the increase in viscosity shown in Paragraph 4.1 is most probably due to the rise of molecular weight and not to chain branching. On the contrary, it is possible to appreciate that the crystallinity of PA11 tends to slightly increase upon the addition of CNTs, especially at elevated nanofiller concentration. This is in agreement with the literature, as CNTs act as nucleating agents and often increase the crystallinity content of the polymer matrices in which they are inserted.⁶⁰ Moreover, by observing the cooling scan, it is possible to assess that the crystallization temperature (T_c) of all compositions containing CNTs is anticipated of 10°C with respect to the unfilled samples, further demonstrating the nucleating potential of CNTs. However, the observed crystallinity increase promoted by CNTs addition is not dramatic, and the 3D printability of the materials would probably be improved upon C introduction, as it is completely amorphous. In these conditions, the shrinkage phenomena that could occur in 3D printed samples during the cooling phase should be limited.

Table 2.

Results of the DSC tests on the prepared bulk samples.

Sample	First scan				Cooling
Sample	T_g [°C]	T_m [°C]	ΔH_m [J/g]	χ [%]	T_c [°C]
PA11	41.2	193.4	40.7	20.2	160.5
PA11_10C	38.9	192.1	47.3	25.9	164.2
PA11_20C	35.6	191.4	38.7	23.1	162.8
PA11_2.5CNT	41.9	191.2	54.4	27.0	170.8
PA11_10C_2.5CNT	36.8	190.8	44.1	24.1	169.8
PA11_20C_2.5CNT	38.3	190.9	39.3	24.3	169.0
PA11_5CNT	43.1	190.2	52.3	27.2	172.4
PA11_10C_5CNT	42.8	189.8	47.1	27.4	170.9
PA11_20C_5CNT	38.1	190.5	43.1	27.0	170.9

The thermal stability of the samples has also been evaluated through thermogravimetric analysis (TGA). However, these results have not been reported since they did not highlight any significant variation promoted by the introduction of C and/or CNTs.

Figure 5 reports representative stress–strain curves of samples containing 5 phr of CNTs, while the numerical results of all the tested compositions are reported in Table 3. By observing the effect given by C introduction, it is possible to see that by increasing its weight fraction, $E$ , $σ_{u t s}$ , and $σ_{b}$ decrease, but, as expected, $ε_{b}$ increases considerably. This increase in ductility confirms the main role of C as a plasticizer and impact modifier. On the contrary, in the samples reinforced with CNTs only, higher values of $E$ , $σ_{u t s}$ , and $σ_{b}$ can be observed with respect to neat PA11, without substantially impairing the deformation at break. Therefore, in the samples with both CNTs and C, the effects of both components are balanced, limiting the reduction in elastic modulus and tensile strength caused by C, while retaining suitable elongation at break values.

Table 3.

Results of tensile and Charpy impact tests performed on the prepared bulk samples.

Tensile properties					Charpy impact properties
Specimen	E [GPa]	σ_UTS [MPa]	σ_b [MPa]	ε_b [%]	E_TOT,SP [kJ/m²]
PA11	1.5 ± 0.1	44.6 ± 1.8	36.9 ± 3.6	45.9 ± 5.8	5.5 ± 1.5
PA11_10C	1.2 ± 0.1	39.4 ± 0.4	32.8 ± 0.7	72.7 ± 2.9	16.2 ± 1.4
PA11_20C	1.0 ± 0.1	32.8 ± 0.5	29.7 ± 1.4	61.4 ± 11.1	15.5 ± 1.8
PA11_2.5CNT	1.7 ± 0.1	49.2 ± 1.2	41.2 ± 3.3	42.1 ± 5.8	5.0 ± 0.1
PA11_10C_2.5CNT	1.4 ± 0.1	39.3 ± 1.4	36.5 ± 0.6	45.2 ± 6.5	8.2 ± 0.1
PA11_20C_2.5CNT	1.2 ± 0.1	32.8 ± 1.4	32.2 ± 1.4	42.5 ± 2.5	7.6 ± 0.1
PA11_5CNT	1.9 ± 0.1	50.1 ± 0.9	43.7 ± 0.6	48.1 ± 5.9	5.0 ± 0.1
PA11_10C_5CNT	1.6 ± 0.1	41.8 ± 1.3	40.3 ± 1.2	45.3 ± 4.3	8.2 ± 0.1
PA11_20C_5CNT	1.2 ± 0.5	32.1 ± 0.8	30.2 ± 0.9	42.6 ± 12.2	8.7 ± 0.1

Regarding the impact properties, thanks to the addition of the impact modifier, the total specific energy absorbed by the specimens is high. By comparing PA11 with PA11_20C, the specific total energy absorbed by the latter is approximately 282 % higher than the former. This increase in the specific total energy absorption can also be connected to the reaction that occurred between PA11 and EAA-g-MAH, as reported in the analysis of the rheological measurements (Figures 1(a) and (b)), FESEM micrographs (Figures 2(a) and (b)), and the FTIR spectra (Figures 3(a) and (b)). On the other hand, the addition of CNTs slightly impairs E_tot, due to CNT agglomeration. However, the addition of C contributes to retaining (and in some formulations to increasing) the original impact properties of PA11. Generally, high impact properties are desired when selecting a polymer for 3D printing, since this reduces layer delamination and cracking, and increases internal stress absorption during cooling and solidification. Therefore, given these results, PA11_20C was selected as a reference composition for filament manufacturing to assess its 3D printability.

To better comprehend the synergistic effect of C and CNTs and to identify the most suitable compositions, a multi-objective optimization method called Response Surface Methodology (RSM) was employed. This process enables the construction of Pareto curves, in which on the y-axis there are σ_UTS and the specific total energy absorbed, and on x and z axes there are CNT and impact modifier content, respectively. Figures 6(a) and (b) illustrate the trends of σ_UTS and E_TOT,SP according to the RSM methodology.

From the RSM analyses, it is possible to identify the potential optimal composition as the one containing 7 wt% of C and 5 phr of CNT (PA11_7C_5CNT), which is characterized by high specific total energy absorbed, high $σ_{U T S}$ , and $ε_{b}$ , according to Figures 6(a) and (b). However, a filament with this composition proved to be non-printable due to its low flowability and flexibility. Therefore, a composition with doubled content of C (i.e., PA11_15C_5CNT) was selected for the filament manufacturing, to ensure printability, while still preserving good mechanical and impact properties.

From the obtained results, a comprehensive understanding of the effects induced by the impact modifier and CNTs on the properties of PA11-based systems has thus been obtained. The addition of C increases viscosity due to chain extension reactions and significantly improves impact resistance without compromising thermal stability and/or printability. Particularly, PA11_20C exhibits an impressive ductility increase compared to neat PA11, confirming its potential as a high toughness material suitable for additive manufacturing. On the other hand, the incorporation of CNTs not only contributes to increasing stiffness and strength but also introduces some issues related to agglomeration, which slightly impairs impact resistance. However, a balanced formulation combining both additives, namely PA11_15C_5CNT, enables the achievement of high toughness and potential printability, while also yielding enhanced strength and stiffness. This synergistic combination provides an optimal trade-off between mechanical reinforcement and ductility, both of which are crucial for high-performance 3D printing applications. Therefore, the following paragraph will be related to 3D printed PA11_20C and PA11_15C_5CNT compositions only.

Characterization of the 3D printed materials

As reported in the previous paragraph, PA11_20C and PA11_15C_5CNT compositions were selected for the manufacturing of filaments for 3D printing, following the methodology described in Paragraph 3.2. Then, their 3D printability was demonstrated through the manufacturing of 3 × 3 × 3 cm³ cubes, as shown in Figure 7. 3D printed PA11_20C cube shows quite precise layer deposition. Typically, the successive stacking of crystalline polymer layers results in anisotropic volumetric shrinkage, rendering polyamides quite difficult to process through FFF. In the present case, the addition of the impact modifier enables effective deposition, making crystallization more difficult, therefore demonstrating the printability of PA11 for the first time.⁶⁹ The cube manufactured with the PA11_15C_5CNT formulation exhibits a more pronounced warping at the interface with the printing plate compared to PA11_20C, attributed to the higher tendency to crystallize upon cooling due to the presence of CNTs (see also the DSC results reported in Paragraph 4.1). Due to this issue, the 3D printing of this filament was performed with the cooling speed kept as low as possible. However, it can be concluded that an acceptable 3D printability was obtained for both formulations.

The numerical results of the quasi-static tensile test performed on 3D printed specimens are reported in Table 4, in terms of elastic modulus (

E

), ultimate tensile strength (

σ_{U T S}

), stress at break (

σ_{b}

), and strain at break (

ε_{b}

). The representative stress–strain curves in this case have not been reported, as they do not exhibit any appreciable difference in the trends compared to those of bulk specimens. It can be noticed that the introduction of carbon nanotubes increases stiffness and strength, while the strain at break is not considerably impaired by the addition of CNTs, because of the toughening effect of C in the PA11_15C_5CNT formulation. From the comparison of the mechanical properties between 3D printed and bulk samples, the higher tensile strength and elastic modulus of bulk PA11_20C with respect to its 3D printed counterparts can be observed, presumably due to the porosities introduced during FFF. On the contrary, in the case of PA11_15C_5CNT, an enhancement in ductility is observed for the 3D printed specimens. For both the compositions, the strain at break increases in 3D printed specimens compared to the corresponding bulk materials. Therefore, it can be stated that the line infill configuration of 3D printed specimens enhances their ductility, thanks to the synergistic effect given by the 3D deposited layers with ±45° orientation. According to these results, it can be hypothesized that an increase of the raster angle up to 180° with infill lines parallel to the loading direction could lead to a further enhancement in strain at break.

Table 4.

Tensile and Charpy impact properties of 3D printed specimens prepared with the two selected formulations.

Properties	PA11_20C	PA11_15C_5CNT
E [GPa]	0.9 ± 0.1	1.3 ± 0.1
σ_UTS [MPa]	27.6 ± 0.2	39.2 ± 2.1
σ_b [MPa]	25.5 ± 1.0	37.3 ± 2.3
ε_b [%]	96.3 ± 14.9	91.3 ± 23.2
Charpy impact energy [kJ/m²]	18.4 ± 1.3	18.8 ± 2.2

From the results of Charpy impact tests reported in Table 4, it can be assessed that the maximum impact energy values are quite comparable between the two compositions, and that the impact energy of 3D printed specimens is significantly higher with respect to that of the corresponding bulk specimens. A possible explanation for this result is that the line infill of the 3D printed specimens allows for higher energy dissipation. Therefore, the obtained enhancement in impact toughness is promoted not only by C introduction but also by the adopted manufacturing parameters.

Additionally, the printability of PA11_15C_5CNT on glass fibers was successfully demonstrated, as shown in Figure 8. This test was conducted to evaluate the feasibility of using a CNT-filled filament as a 3D printed healing agent in epoxy/glass fiber composites. It can be noted that the deposition of the PA11_15C_5CNT filament is not hindered by any shrinkage phenomenon. However, to have an effective deposition onto the plies, it was necessary to increase the nozzle temperature up to 285°C.

In this work, a peeling test for adhesives was adapted to evaluate the adhesion properties of the 3D printed PA11_15C_5CNT filament on glass fibers, as described in Paragraph 3.3.2.2. From the representative load-extension plots reported in Figure 9(a), a stepped peeling curve can be observed, showing a maximum peeling load of 25.0 ± 9.6 N and an average peeling force of 17.2 ± 6.1 N, quite higher with respect to the values for pressure-sensitive adhesive tapes, which are usually in the range of a few N.⁷⁰ The morphology of the curve suggests good adhesion between the polymer and the fibers, since the separation of the two adherents occurs at discrete steps and does not follow a continuous trend, demonstrating the good peel resistance of the 3D printed polymer against the fibers. Good adhesion is also confirmed by the image reported in Figure 9(b), showing the specimen after the test, in which glass fiber residues are visible on the 3D printed polymer surface. This strong adhesion between the polymeric filament and glass fibers is likely attributed to the high affinity between the silane-based sizing layer on the glass fibers and the amide groups of PA11. Interactions may include hydrogen bonding, ionic or covalent coupling of the functional silane tail groups with polymer functional groups.^71,72

In conclusion, the successful printability onto fibers, the observed increase in impact energy for 3D-printed specimens and most importantly, the peeling resistance of PA11_15C_5CNT against glass fibers, represent key findings. These results highlight the potential for enhanced toughening in self-healing glass fiber–reinforced composites, which will be developed in future work.⁷³

Conclusions

In this study, blends based on PA11 were developed by incorporating an EAA-g-MAH impact modifier, thereby enabling FFF printing of PA11 for the first time. Furthermore, carbon nanotubes (CNTs) were incorporated into the blends to increase the mechanical properties of the blends, mainly in terms of stiffness. This research examined the potential application of these 3D printable formulations as a thermoplastic healing agent in self-healing composites, wherein the healing agent is directly 3D printed onto the dry fiber plies before manufacturing the laminate.

The PA11-based formulations, containing 10–20 wt% impact modifier and 2.5–5 phr of CNTs, were prepared and characterized rheologically, thermally, and mechanically to identify the optimal compositions for 3D printing. Adding 20 wt% of impact modifier causes an increase in ductility and impact energy compared to neat PA11. The introduction of CNTs, as expected, improved the elastic modulus and tensile strength, but increased polymer crystallinity due to nucleation phenomena, which may hinder PA11 printability. However, balancing CNTs effect with the impact modifier, makes PA11 3D printable. The optimized compositions selected after these preliminary tests, that is, PA11_20 C and PA11_15C_5CNT, were successfully 3D printed, and the 3D printed specimens also showed enhanced impact properties and ductility compared to bulk compositions.

As a further step, the PA11_15C_5CNT filament was successfully deposited via fused filament fabrication (FFF) onto glass fiber plies, demonstrating its potential use as a healing agent in self-healing epoxy/glass fiber composites. T-Peel tests confirmed good interfacial adhesion between the PA11_15C_5CNT filament and the glass fibers, indicating a promising toughening effect for future composite systems.

In conclusion, this study reports the first successful 3D printing of PA11-based blends via FFF, contributing to the advancement of semi-crystalline polymers within accessible additive manufacturing technologies. The results highlight the potential of CNT-reinforced PA11 blends as 3D-printable healing agents for self-healing glass fiber composites, owing to their enhanced toughness and strong adhesion to glass fiber plies. Building on these findings, future work will focus on the fabrication of self-healing glass fiber composites in which the healing mechanism is driven by thermo-reversible entanglement of the embedded healing agent. The primary objective is to advance self-healing composite technologies by developing tough materials capable of multiple healing cycles, thereby extending their service life.

Footnotes

ORCID iDs

Dalia Di Liberto

Giulia Fredi

Daniele Rigotti

Alessandro Pegoretti

Funding

The authors disclosed receipt of the following financial support for the research, authorship, and/or publication of this article: Funded by the European Union—Next Generation EU—PNRR, Mission 4 Component 2, Investment 1.3—PE MICS Spoke 5—LOLIMAR Project (PE00000004, CUP D43C22003120001).

Declaration of conflicting interests

The authors declared no potential conflicts of interest with respect to the research, authorship, and/or publication of this article.

Data Availability Statement

Data sharing not applicable to this article as no datasets were generated or analyzed during the current study.

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