Fernandes, E. M., Correlo, V. M., Chagas, J. A. M., Mano, J. F., and Reis, R. L. (2010). Cork based composites using polyolefin's as matrix: morphology and mechanical performance. *Compos. Sci. Technol.* 70, 2310-2318. doi: 10.1016/j.compscitech.2010.09.010

Fernandes, E. M., Correlo, V. M., Chagas, J. A. M., Mano, J. F., and Reis, R. L. (2011). Properties of new cork-polymer composites: advantages and drawbacks as compared with commercially available fibreboard materials. *Compos. Struct.* 93, 3120-3129. doi: 10.1016/j.compstruct.2011.06.020

Fletes, R. C. V., López, E. O. C., Sánchez, F. J. M., Mendizábal, E., Núñez, R. G., Rodrigue, D., et al. (2020). Morphological and mechanical properties of bilayers wood-plastic composites and foams obtained by rotational molding. *Polymers* 12, 1-19. doi: 10.3390/polym12030503

Fortes, M. A., and Nogueira, M. T. (1989). The poisson effect in cork. *Mater. Sci. Eng. A* 122, 227-232.

Gibson, L., Easterling, K., and Ashby, M. (1981). The structure and mechanics of cork. *Proc. R. Soc. London A* 377, 99-117. doi: 10.1007/s10853-011-5914-9

Gibson, L. J. (2005). Biomechanics of cellular solids. *J. Biomech.* 38, 377-399. doi: 10.1016/j.jbiomech.2004.09.027

Gibson, L. J. (2003). Cellular solids. *MRS Bull.* 28, 270-274.

Gil, L. (2007). *Cork as a building material Technical Manual*. Santa Maria de Lamas: APCOR-Associação Portuguesa de Cortiça.

Greco, A., Ferrari, F., Buccoliero, M. G., and Trono, G. (2019). Thermal and mechanical analysis of polyethylene homo-composites processed by rotational molding. *Polymers* 11:528. doi: 10.3390/polym11030528

Gul, J., Saleemi, A. R., Mirza, S., Feroze, N., Mansha, M., Gul, J., et al. (2013). Thermal and mechanical characteristics of cork filled insulation for aerospace applications. *Plast. Rubber Compos.* 39, 28-32. doi: 10.1179/174328910X12608851832010

Hansen, D., and Bernier, G. A. (1972). Thermal conductivity of polyethylene: the effects of crystal size, density and orientation on the thermal conductivity. *Polym. Eng. Sci.* 12, 204-208. doi: 10.1002/pen.76012 0308

Jansri, E., and O-charoen, N. (2018). Polypropylene / polyethylene two-layered by one-step rotational molding. *J. Polym. Eng.* 38, 685-694. doi: 10.1515/polyeng-2017-0367

Kim, H. S., Kim, S., Kim, H. J., and Yang, H. S. (2006). Thermal properties of bio-flour-filled polyolefin composites with different compatibilizing agent type and content. *Thermochim. Acta* 451, 181-188. doi: 10.1016/j.tca.2006.09.013

Mano, J. F. (2002). The viscoelastic properties of cork. *J. Mater. Sci.* 37, 257-263. doi: 10.1023/A:1013635809035

Nugent, P. (2017). "Rotational Molding," in *Applied Plastics Engineering Handbook: Processing, Materials, and Applications*, ed. M. Kutz (Reading, PA: Elsevier).

Oliveira, M. J., Cramez, M. C., and Crawford, R. J. (1996). Structure-property relationships in rotationally moulded polyethylene. *J. Eng. Appl. Sci.* 2, 2096-2099. doi: 10.1007/BF01152932

Oliveira, M. J., Cramez, M. C., Oliveira, M. J., and Cramez, M. C. (2007). Rotational molding of polyolefins: processing, morphology and properties. *J. Macromol. Sci. Part B* 40, 457-471. doi: 10.1081/MB-1001 06170

Ortega, Z., Monzón, M. D., Benítez, A. N., Kearns, M., McCourt, M., and Hornsby, P. R. (2013). Banana and abaca fiber-reinforced plastic composites obtained by rotational molding process. *Mater. Manuf. Process.* 28, 879-883. doi: 10.1080/10426914.2013.792431

Pereira, H. (1988). Chemical composition and variability of cork from *Quercus suber* L. *Wood Sci.Technol.* 218, 211-218. doi: 10.1007/bf003 86015

Pereira, H. (1992). The thermochemical degradation of cork. *Wood Sci.Technol.* 269, 259-269.

Pereira, H. (2007). *Cork: Biology, Prodution and Uses*. Lisboa: Elsevier.

Pereira, H. (2015). Rotational molding: some basics. *Pop. Plast. Packag.* 10, 26-28.

Pereira, H., and Velez Marques, A. (1988). The effect of chemical treatments on the cellular structure of cork. *IAWA J.* 9, 337-345. doi: 10.1163/22941932-90001093

Pick, L. T., and Harkin-Jones, E. (2006). Effect of mould pressurisation on impact strength of rotationally moulded polyethylenes. *Plast. Rubber Compos.* 35, 324-330. doi: 10.1179/174328906X143868

Pontes, A. J., Ferreira, E. C., Mendes, J., Fonseca, J., Cramez, M. C., and Pouzada, A. S. (2014). Inovação na moldação por injeção. *Rev. da Cefanol Assoc. Nac. da Indúst. Moldes* 103, 23-29.

Ramkumar, P. L., Kulkarni, D. M., and Chaudhari, V. V. (2014). Parametric and mechanical characterization of linear low density polyethylene (LLDPE) using rotational moulding technology. *Sadhana* 39, 625-635. doi: 10.1007/s12046-013-0223-4

Ramkumar, P. L., Waigaonkar, S. D., and Kulkarni, D. M. (2016). Effect of oven residence time on mechanical properties in rotomoulding of LLDPE. *Sadhana* 41, 571-582. doi: 10.1007/s12046-016-0489-4

Rosa, M. E., and Fortes, M. A. (1988a). Temperature-induced alterations of the structure and mechanical properties of cork. *Mater. Sci. Eng.* 100, 69-78. doi: 10.1016/0025-5416(88)90240-6

Rosa, M. E., and Fortes, M. A. (1988b). Termogravimetric analysis of cork. *J. Mater. Sci. Lett.* 7, 1064-1065. doi: 10.1007/BF00720828

Şen, A., Marques, A. V., Gominho, J., and Pereira, H. (2012). Study of thermochemical treatments of cork in the 150-400°C range using colour analysis and FTIR spectroscopy. *Ind. Crops Prod.* 38, 132-138. doi: 10.1016/j. indcrop.2012.01.018

Şen, A., Van Den Bulcke, J., Defoirdt, N., Van Acker, J., and Pereira, H. (2014). Thermal behaviour of cork and cork components. *Thermochim. Acta* 582, 94-100. doi: 10.1016/j.tca.2014.03.007

Silva, M. D., Oleskovicz, M., and Coury, D. V. (2005). Cork: properties, capabilities and applications. *Control. y Autom.* 16, 345-358. doi: 10.1179/174328005X41168

Soares, B., Reis, L., and Sousa, L. (2011). Cork composites and their role in sustainable development. *Procedia Eng.* 10, 3214-3329. doi: 10.1016/j.proeng. 2011.04.531

Subedi, D. P. (2011). Contact angle measurement for the surface characterization of solids. *Himal. Phys.* 2, 1-4. doi: 10.3126/hj.v2i2.5201

Torres, F. G., and Aragon, C. L. (2006). Final product testing of rotational moulded natural fibre-reinforced polyethylene. *Polym. Test.* 25, 568-577. doi: 10.1016/j. polymertesting.2006.03.010

Vignali, A., Iannace, S., Falcone, G., Utzeri, R, Stagnaro, P., and Bertini, F. (2019). Lightweight poly-(ε-caprolactone) composites with surface modified hollow glass microspheres for use in rotational molding: thermal, rheological and mechanical properties. *Polymers* 11:624. doi: 10.3390/polym11040624

Vilela, C., Sousa, A. F., Freire, C. S. R., Silvestre, A. J. D., and Pascoal Neto, C. (2013). Novel sustainable composites prepared from cork residues and biopolymers. *Biomass Bioenergy* 55, 148-155. doi: 10.1016/j.biombioe.2013.01.029

Wei, L., McDonald, A. G., Freitag, C., and Morrell, J. J. (2013). Effects of wood fiber esterification on properties, weatherability and biodurability of wood plastic composites. *Polym. Degrad. Stab.* 98, 1348-1361. doi: 10.1016/j. polymdegradstab.2013.03.027

Xu, L., and Crawford, R. J. (1993). Analysis of the formation and removal of gas bubbles in rotationally moulded thermoplastics. *J. Mater. Sci.* 28, 2067-2074. doi: 10.1007/BF00367563

Yuan, X., Easteal, A. J., and Bhattacharyya, D. (2007). Mechanical performance of rotomoulded wollastonite-reinforced polyethylene composites. *Int. J. Mod. Phys. B* 21, 1059-1066. doi: 10.1142/S0217979207036916

**Conflict of Interest:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2020 Martins and Gil. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
<figure><img src="image_351.png" /></figure>

Check for
updates

# Statistical Fatigue Investigation and Failure Prediction of a Healable Composite System

Nathan Hostettler, Amaël Cohades and Véronique Michaud*

Laboratory for Processing of Advanced Composites (LPAC), Institute of Materials, Ecole Polytechnique Fédérale de Lausanne (EPFL), Lausanne, Switzerland

Fiber reinforced polymers are massively used as an alternative to metals in structural applications. The brittle nature of their matrix, however, makes them more susceptible to crack formation and propagation resulting in costly repair operations and increased environmental impact. Intrinsic healable composites provide a good alternative to these conventional composite materials, whereas their mechanical properties in static solicitation or impact testing are well documented, only few studies address fatigue testing. This research focuses on 3-point bending fatigue tests of polymer-blend based healable E-glass composite materials. The S–N curve was first built to compare the fatigue behavior of the healable system to a conventional epoxy composite. A statistical approach based on Weibull statistics was developed to predict the failure probability as a function of the applied stress amplitude, to compare both systems at equivalent probability of failure. The healable system showed a higher fatigue resistance at high cycle fatigue. Furthermore, a full stiffness recovery was obtained and a life extension of at least five times compared to the reference system when healed after reaching a 90% chance of survival. The healable system thus opens new perspectives for more sustainable load-bearing composites.

**Keywords:** self-healing, composites, fatigue, failure prediction, Weibull statistics

**OPEN ACCESS**

**Edited by:**

Patricia Krawczak,
IMT Lille Douai, France

**Reviewed by:**

Ming Qiu Zhang,
Sun Yat-sen University, China
Nancy Sottos,
University of Illinois at
Urbana-Champaign, United States
Ruben Dirk Sevenois,
Ghent University, Belgium

**\*Correspondence:**

Véronique Michaud
veronique.michaud@epfl.ch

**Specialty section:**

This article was submitted to Polymeric and Composite Materials, a section of the journal Frontiers in Materials

**Received:** 13 May 2020

**Accepted:** 24 August 2020

**Published:** 14 September 2020

**Citation:**

Hostettler N, Cohades A and Michaud V (2020) Statistical Fatigue Investigation and Failure Prediction of a Healable Composite System. Front. Mater. 7:561852.
doi: 10.3389/fmats.2020.561852

## INTRODUCTION

The use of fiber reinforced polymers as a stiff and lightweight alternative to metal in structural applications enables to reach higher performance especially in transport, energy, and aerospace industry. However, the brittle nature of these materials makes them more prone to damage during manufacturing, handling, maintenance activities or during operation. Repair is not straightforward and often requires to remove part of the damaged zone and adhesively bond composite patches (Bond et al., 2008; Katnam et al., 2015; Préau et al., 2016). The overall budget for composite materials repair is thus significant and expected to reach not less than 22 Bn$ worldwide by 2026 (MarketsAndMarkets, 2017); in parallel, tons of unrecycled composite waste are produced due to unrepaired damaged structures, reaching 40,000 tons in France for the year 2015 (JEC-GROUP, 2017). Self-healing and more especially healable composites are a promising alternative to conventional composites to reduce repair costs and waste because of their ease of repair, demanding no additional material, which sometimes goes along with increased recyclability using moderate manpower and energy.
Healable polymer systems have been under investigation in the recent years, in particular integrating healing functionalities into thermoset polymers for the autonomous repair of subcritical damage in composite materials (Cohades et al., 2018). Two approaches are usually identified: a) extrinsic healable systems, where an extrinsic healing agent is integrated in the matrix through loaded microcapsules or vasculs, and is released after damage events so as to flow, fill, and heal the crack after polymerization (Blaiszik, et al., 2010; Kling and Czigány 2014), b) intrinsic healable systems, based on a modification of the matrix chemistry itself providing autonomous healing capabilities to the system (Garcia, 2014). Extrinsic systems generally show a reduction in composite mechanical properties by the integration of a large amount of capsules or vasculs necessary to reach good healing capabilities; alternatively, the volume fraction of structural fibers is reduced, and the textile structure modified to make sufficient space for the healing agent containers; finally, these are limited to one healing cycle, and/or require the use of external pumping systems to circulate the healing fluids (Kessler and White, 2001; Patel et al., 2010; Norris et al., 2013; Patrick et al., 2014; Manfredi et al., 2015; Zhu et al., 2016). The intrinsic healing approach solves most of these issues providing the initial matrix with healable capabilities through reversible covalent (Park et al., 2009; Heo and Sodano 2015) bonds (Post et al., 2017; Shabani et al., 2020), supramolecular chemistries (Sordo and Michaud, 2016) or polymer blends (Pingkarawat et al., 2013; Cohades and Michaud 2017). Even if these intrinsic systems demonstrate good healing capabilities, and allow composites with a high load of reinforcing fibers, their mechanical properties are usually below those of conventional epoxy in term of stiffness, strength and damage tolerance. This is particularly true for composites based on matrices that heal at room temperature, which tend to show high mobility of bonds, and thus low mechanical properties (Deng et al., 2015; Sordo and Michaud 2016).

Healing efficiency is generally determined in composite materials through static fracture mechanics tests, in mode I or II, bending tests or compression after impact (Tsangouri et al., 2015); however, few studies are carried out on cyclic damage progression, as found in fatigue, although this case is one where small matrix cracks initially develop, that may be healed if caught at an early stage. Several studies were conducted on the fatigue resistance of extrinsic self-healing systems showing a retardation of failure due to the increased tortuosity of the crack path (Brown et al., 2005; Hamilton et al., 2012; Luterbacher et al., 2016; Ye, et al. 2017; Kim et al., 2019); however, little work is reported on the fatigue performance of intrinsically healable composites, although it is a key parameter for their long term integration as a structural material. To the best of our knowledge, only Pingkarawat and co-workers (Pingkarawat et al., 2012; Pingkarawat et al., 2013; Pingkarawat et al., 2014; Ladani et al., 2019) investigated the fatigue behavior of carbon-epoxy composites, using thermoplastic additives and only in double cantilever beam forced crack propagation mode, under cyclic loading. They found that thermoplastic additives in general helped retard the fatigue crack growth in mode I due to the presence of bridging thermoplastic ligaments formed in the

healing operation at the crack location. Polyethylene-co-methacrylic acid (EMAA) was found to be the most effective, and was shown to also delay crack growth under cyclic loading when under the form of stitches, provided the quantity was high enough, or when they were combined with carbon fiber tows (Pingkarawat et al. 2014; Ladani et al., 2019). In this research, a commercial healable system based on epoxy/thermoplastic polymer blends was tested in 3-point bending fatigue to investigate its performance compared to a conventional epoxy-based composite. This loading method was chosen to concentrate the damage at the center of the specimen and to mimic potential applications where the part is loaded in flexural mode. Furthermore, a statistical approach based on Weibull statistics was developed to predict the probability of failure of the samples as a function of any applied stress amplitude. This approach was used as a damage criterion to reach a similar extent of damage in the two systems before healing, and re-testing to assess the healing efficiency.

# MATERIALS AND METHODS

## Materials and Samples Preparation

Two types of resin were used to produce the samples. A first set of plates was produced with the healable system from CompPair Technologies Ltd. This novel resin system is based on epoxy/thermoplastic blends capable of autonomous healing by heat treatment at 150°C for 30 min. The second set was produced with the EPIKOTE<sub>TM</sub> MGS RIMR 135 from Hexion, a difunctional bisphenol A/epichlorohydrin derived liquid epoxy resin commonly used in rotor blades industry for wind turbines or boat building and cured with EPIKURE<sub>TM</sub> Curing Agent MGS<sub>TM</sub> RIMH 137 with a 100:30 weight mixing ratio and used as reference. The reinforcement was a woven twill 2 × 2 E-glass fabric with a nominal areal weight of 390 g/m², 6 end/cm for wrap fibers and 6.7 picks/cm for weft fibers, with a diameter of 9 µm, a yarn thickness of 0.45 mm, warp tex of 68 × 5 and weft tex of 272, from Suter-Kunststoffe AG.

Both systems were produced by vacuum assisted resin infusion molding. Sixteen layers of reinforcements were cut in 350 mm × 250 mm rectangles and stacked following a quasi-isotropic stacking sequence of [(+45/-45);(0/90)]<sub>4s</sub>. The stack was infused at room temperature through the long edge with the resin systems and then cured at 60°C for 6 h. A post-cure treatment was then performed for 4 h at 80°C and 1 h at 100°C to reach the infinite glass transition temperature of the system ($T_{g,inf}$ = 97°C). The resulting composite had a target fiber volume fraction ($V_f$) of 50 vol% with a thickness of ~5 mm.

Fractured samples from each system were transversally cut, embedded in EPOFIX resin (Struers, Inc., Cleveland, United States) and polished to a 1 µm diamond paste finish. The crack profiles were then observed with an optical microscope (Olympus BH-2).

## Three-Point Bending Fatigue Testing

Rectangular samples were cut from the infused plates with dimensions following 3-point bending ASTM D7264 (ASTM
International, 2015) with a length, width and span-to-thickness ratio of respectively, 120, 13 mm and 16:1 (80 mm for 5 mm thick samples). The analysis method following this norm implies treating the material as homogeneous which is incorrect for a stack of laminates. In the special case of the symmetrical quasi-isotropic layup studied here, a preliminary static stress analysis with Altair ESAcompTM, using a woven 2-2 Twill glass and an equivalent epoxy resin, showed that it can be approximated as a homogeneous material in 3-point bending, allowing the use of this simple approach for comparison purposes. Nonetheless, the values provided in this work are thus an apparent maximal stress, apparent maximal strain, and apparent flexural secant modulus as defined in the norm, and are also based on a small strain approximation. The tests were performed at a frequency of 5 Hz on a 100 kN servo-hydraulic Schenk testing machine. Preliminary tests were performed at 0.1, 2, and 5 Hz to confirm that the heat generation due to micro-friction in the crack regions was negligible up to 5 Hz, in order to ensure that no healing could take place due to local heating during fatigue testing (Herman et al., 1990; Gagel et al., 2006; Adam and Horst 2017; Movahedirad et al., 2018). These preliminary tests at 0.1 and 2 Hz were also integrated in the final S-N curve.

The tests were conducted at room temperature under load control with a loading ratio of $R = 0.1$ in compression. At least three ($N_{cycle} > 106$) and usually five ($N_{cycle} < 106$) replicates were carried out for each stress conditions to establish the S-N curve of the two systems. Note that because of the difference in volume fraction of fibers ($V_f$) between the healable and reference systems (Table 1), all the results are presented normalized to an equivalent volume fraction of $V_f = 50\%$ (which was the theoretical target volume fraction) for better comparison. This difference is due to the higher viscosity of the healable system that leads to a lower fiber compaction during infusion.

To investigate the healing efficiency of the system, a damage criterion was defined through Weibull statistical analysis of the fatigue experiments allowing direct access to a survival probability of the sample at a specific stress amplitude as a function of the number of cycles. This approach was necessary to take into account any difference in thickness and stiffness between the two systems and be able to compare and reach a similar amount of damage prior healing, since there is not direct access to damage levels in these bulk fatigue experiments, as compared to forced crack growth as performed in the literature.

The healing performance tests were performed at the stress amplitude of 220 MPa ($V_f = 50\%$ equivalent) at 5 Hz. Once the sample had reached a number of cycles corresponding to a survival probability of 90% (determined from the Weibull statistical study), the tests were stopped, the samples were healed at 150°C for 30 min in an oven, but under no applied pressure, then cooled down for an additional 30 min at room

temperature (approximately 20 min from 150°C to ambient plus an additional 10 min for temperature equilibration) and again tested in fatigue at the exact same stress amplitude. This test-healing procedure was repeated several times, and at each step the change of the apparent flexural secant modulus calculated from the estimated maximal strain value at the maximal load of each cycle to determine the lifetime extension of the materials.

## RESULTS AND DISCUSSION

### S-N Curve Determination

The results of the fatigue tests, in the form of conventional S-N curves (stress amplitude vs. number of cycles to failure), are given in Figure 1 for the two studied systems, along with a standard power law fit. The stress amplitude was corrected to an equivalent volume fraction of $V_f = 50\%$ to take into account the large difference in thickness due to the processing of the healable system. This correction was carried out simply multiplying the stress amplitude by the ratio $0.5/V_f$ explaining the scattering of the point regarding the applied stress. The ultimate strength in static mode of the reference system (437 ± 20 MPa) is slightly higher than that of the healable systems (361 ± 23 MPa) due to the presence of the thermoplastic phase accordingly reducing the flexural strength. Note that this diminution follows a rule of mixture considering the amount of thermoplastic present in the healable system, confirming the good quality of the composites. Damage, however, progresses faster in the pure epoxy system with a fatigue curve decreasing faster and even crossing the healable systems endurance curve at a stress amplitude of ~200 MPa. The healable system is thus more durable than the reference one in particular for a stress amplitude lower than 200 MPa, i.e., for high cycle fatigue. This case of fatigue is typically the solicitation observed in structural parts during their lifetime. In both systems, damage mainly occurred by interlaminar crack propagation through the fiber tows and at the resin/fiber interface

<table><caption>TABLE 1 | Volume fraction of fiber for the healable and reference system.</caption><thead><tr><td>System</td><td>V<sub>f</sub></td><td>SD</td></tr></thead><tbody><tr><td>Healable</td><td>0.39</td><td>0.0206</td></tr><tr><td>Reference</td><td>0.48</td><td>0.0052</td></tr></tbody></table>

<figure><img src="image_352.png" /><figcaption>FIGURE 1 | S-N curve for the healable and reference system.</figcaption></figure>
<figure><img src="image_353.png" /><figcaption>FIGURE 2 | (Left) Micrograph of a reference sample after failure (σa = 191.8 MPa). (Right) Micrograph of a healable sample after failure (σa = 204.9 MPa).</figcaption></figure>

(see Figure 2). Whereas the crack path in the reference system clearly follows the tows outer boundary or directly progresses through the tows, the healable system shows a similar crack path location, but with a greater tortuosity which is expected since the system is tougher than pure epoxy. Note that the lower fiber volume fraction in the healable samples probably led to resin pockets known to affect the local stress concentration and reduce the fatigue performance of the material (Silva et al., 2005). However, as the damage mostly propagated through the fiber tows, we assumed that this effect can be neglected when comparing our two types of samples. Nevertheless, the overall fatigue resistance of the healable systems might therefore be underestimated here, keeping a conservative approach in the conclusions drawn.

## Damage Criterion Determination—Weibull Statistics

To evaluate the healing efficiency of the healable systems in fatigue, a damage criterion had to be established, so as to determine when to stop the test for each sample at the same amount of damage. Indeed, simply fixing a stress amplitude and stopping the test at an arbitrary number of cycles would not affect the two systems identically because of their different intrinsic

fatigue response. The same conclusion is reached when considering the reduction in flexural secant modulus because of the statistical nature of the tests. Damage was thus considered as a probability of failure based on a two parameters Weibull distribution characterized by a probability density function f(x) (Khashaba, 2003; Sakin and Ay, 2008):

$$f(x) = \frac{\beta}{\alpha} \left(\frac{x}{\alpha}\right)^{\beta-1} e^{-(x/\alpha)^{\beta}} \quad (1)$$

where α and β are the shape and the scale parameter, respectively. Following the methods described in Sakin and Ay (2008), the cumulative density function is obtained by integrating the probability density function, obtaining a probability of failure Ff and a probability of survival Fs for a number of cycles x:

$$F_f = 1 - F_s = 1 - e^{-(x/\alpha)^\beta}. \quad (2)$$

By taking the natural logarithm and rearranging, the following linear regression model is obtained:

$$Y = mX + c, \quad (3)$$

With X = ln(x), m = β, c = -β(ln(α)).

To determine the Weibull parameters, the unreliability of each fatigue tests must be determined. The median ranks, the true

<figure><img src="image_354.png" /><figcaption>FIGURE 3 | Weibull lines for the healable system.</figcaption></figure>

<figure><img src="image_355.png" /><figcaption>FIGURE 4 | Weibull lines for the reference system.</figcaption></figure>
<figure><img src="image_356.png" /><figcaption>FIGURE 5 | Weibull curves for the healable system.</figcaption></figure>

probability of failure that the ith sample of a group should have in a group of n samples at a 50% confidence level, can be used to estimate this value. In this study, Bernard's median rank approximation was used:

$$MR = \frac{i - 0.3}{n + 0.4} \qquad (4)$$

By grouping the fatigue tests of similar stress amplitude in groups of n samples, the median rank of each sample with a rank i = 1, 2, 3, ..., n could be determined. From the Weibull lines obtained by plotting ln(ln(1/(1-MR))) against ln(Nt) as plotted in **Figures 3 and 4**, the parameters of linear regressions Y(X) could be obtained for each stress level. This approach was carried out for both the healable and reference system allowing to draw survival curves as a function of the test's stress amplitude, **Figures 5 and 6**. The detailed values for each stress conditions and systems as well as the obtained Weibull parameters are shown in **Tables 2 and 3**.

<figure><img src="image_357.png" /><figcaption>FIGURE 6 | Survival curves for the reference system.</figcaption></figure>

<table><caption>TABLE 2 | Weibull's parameter and parameters used to draw the Weibull lines for the healable system.</caption><thead><tr><th>σ<sub>a</sub> (MPa)</th><th># Cycles</th><th>Rank</th><th>α</th><th>β</th></tr></thead><tbody><tr><td rowspan="4">282.13</td><td>466</td><td>1</td><td rowspan="4">1,004</td><td rowspan="4">1.910</td></tr><tr><td>562</td><td>2</td></tr><tr><td>1,069</td><td>3</td></tr><tr><td>1,335</td><td>4</td></tr><tr><td rowspan="4">243.87</td><td>5,198</td><td>1</td><td rowspan="4">17,574</td><td rowspan="4">1.492</td></tr><tr><td>11,840</td><td>2</td></tr><tr><td>19,187</td><td>3</td></tr><tr><td>22,970</td><td>4</td></tr><tr><td rowspan="3">164.38</td><td>1,026,365</td><td>1</td><td rowspan="3">4,900,399</td><td rowspan="3">0.950</td></tr><tr><td>3,500,000</td><td>2</td></tr><tr><td>7,698,653</td><td>3</td></tr><tr><td></td><td>-</td><td>-</td><td></td><td></td></tr></tbody></table>

The curves obtained from the median rank analysis only provide a survival probability for a determined stress amplitude; there is no direct generalization. To extend this analysis to the full range of possible stress values, the dependence of the Weibull parameters with respect to the stress amplitude must be quantified. From the definition of the survival probability and assuming that the number of cycles x generally follows a power law behavior with respect to the applied stress σ:

$$F_S = e^{-(\frac{x}{\alpha})^{\beta}} = e^{-(\frac{\sigma^b}{\alpha})^{\beta}} \qquad (5)$$

The stress can now be directly linked to the first Weibull parameter α as:

$$\alpha(-\ln(F_S))^{1/\beta} = \sigma^b \qquad (6)$$

As β is usually much smaller than α and assuming a constant value for F_S the expression would reduce to:

$$C\alpha \approx \sigma^b. \qquad (7)$$

From this simple analysis, α seems to be linked by a power law to σ. However, fitting the experimental data shows that the relation is closer to:

<table><caption>TABLE 3 | Weibull's parameter and the parameters used to draw the Weibull lines for the reference system.</caption><thead><tr><th>σ<sub>a</sub> (MPa)</th><th># Cycles</th><th>Rank</th><th>α</th><th>β</th></tr></thead><tbody><tr><td rowspan="4">321.33</td><td>320</td><td>1</td><td rowspan="4">959</td><td rowspan="4">1.332</td></tr><tr><td>417</td><td>2</td></tr><tr><td>1,149</td><td>3</td></tr><tr><td>1,314</td><td>4</td></tr><tr><td rowspan="4">274.94</td><td>3,248</td><td>1</td><td rowspan="4">8961</td><td rowspan="4">1.357</td></tr><tr><td>3,763</td><td>2</td></tr><tr><td>9,819</td><td>3</td></tr><tr><td>13,140</td><td>4</td></tr><tr><td rowspan="4">180.91</td><td>332,004</td><td>1</td><td rowspan="4">493,527</td><td rowspan="4">3.256</td></tr><tr><td>358,002</td><td>2</td></tr><tr><td>441,561</td><td>3</td></tr><tr><td>628,491</td><td>4</td></tr></tbody></table>
<figure><img src="image_358.png" /><figcaption>FIGURE 7 | Determination of Weibull parameters as a function of the stress amplitude.</figcaption></figure>

<figure><img src="image_359.png" /></figure>

$$C\alpha \approx \ln(\sigma^b). \qquad (8)$$

From this, it is then possible to find an expression for $\alpha$ as a function of the stress amplitude fitting the experimental results as:

$$\alpha(\sigma) = e^{\left(\frac{\sigma - C_2}{C_1}\right)^{\frac{1}{C_3}}}. \qquad (9)$$

The same approach was performed with the second Weibull parameter $\beta$. Starting from the last expression of the previous analysis:

$$\alpha(-\ln(F_s))^{\frac{1}{\beta}} = \sigma^b \leftrightarrow \frac{1}{\beta}\ln(\alpha \ln(-\ln(F_s))) = b \ln(\sigma) \qquad (10)$$

Assuming again a constant value for $F_s$ and taking the exponential on the expression gives:

$$\sigma = \alpha C e^{\left(\frac{1}{\beta b}\right)} \approx e^{\left(\frac{1}{\beta b}\right)}$$

The stress amplitude is then exponentially linked to $\beta$. It is finally possible to find an expression for $\beta$ as a function of the stress amplitude fitting the experimental results as:

<table><caption>TABLE 4 | Detailed fitting parameters to determine the Weibull parameters as a function of the stress amplitude.</caption><thead><tr><th>Equation</th><th>System</th><th>Parameters</th></tr></thead><tbody><tr><td rowspan="3">α(σ) = e<sup>(σ-C<sub>2</sub>/C<sub>1</sub>)<sup>1/3</sup></sup></td><td rowspan="2">Healable</td><td>C<sub>1</sub> = -13.189</td></tr><tr><td>C<sub>2</sub> = 1.004</td></tr><tr><td rowspan="5">β(σ) = (ln((σ-C<sub>2</sub>/C<sub>1</sub>))<sup>1/3</sup>)</td><td rowspan="5">Reference</td><td>C<sub>3</sub> = 373.918</td></tr><tr><td>C<sub>1</sub> = -6.998</td></tr><tr><td>C<sub>2</sub> = 1.372</td></tr><tr><td>C<sub>3</sub> = 419.663</td></tr><tr><td>C<sub>1</sub> = -611.26</td></tr><tr><td rowspan="5">Reference</td><td rowspan="5">Healable</td><td>C<sub>2</sub> = 1,839.149</td></tr><tr><td>C<sub>3</sub> = -0.104</td></tr><tr><td>C<sub>1</sub> = -165.104</td></tr><tr><td>C<sub>2</sub> = 0</td></tr><tr><td>C<sub>3</sub> = -1.805</td></tr></tbody></table>

$$\beta(\sigma) = \left( \ln \left( \frac{\sigma - C_2}{C_1} \right) \right)^{\frac{1}{C_3}}$$

These fitting approaches were successfully applied to both Weibull parameters on the two studied systems. An example of the obtained fits is shown for the healable system in **Figure 7** as well as the detailed fitting parameters in **Table 4**.

Now that an expression for the Weibull parameters as a function of the applied stress was found, it is possible to determine the survival curve at any stress amplitude. A comparison of the experimental values and the fitted survival curves as well as additional new ones obtained from the fitting approach at intermediate values are shown for the healable and reference system in **Figures 8 and 9**.

The obtained fits closely follow the experimental data giving a good approximation of the survival probability of the

<figure><img src="image_360.png" /><figcaption>FIGURE 8 | Survival curves of the healable system obtained from the fitted Weibull parameters (dashed lines) as well as the initial experimental data (solid dots).</figcaption></figure>
<figure><img src="image_361.png" /><figcaption>**FIGURE 9** | Survival curves of the reference obtained from the fitted Weibull parameters (dashed lines) as well as the initial experimental data (solid dots).</figcaption></figure>

samples as a function of the stress amplitude. It is interesting to note on **Figure 8** that the curve at 164 MPa slightly diverges from the prediction. This can be linked to the reduced number of tested samples at this low stress amplitude because of the time scale of the experiments (sometimes extending up to 17 days of tests).

## Healing Efficiency

The healing efficiencies of each system were tested at 220 MPa ($V_f$ = 50% equivalent) until a probability of survival of 90% was reached. This probability corresponds to 18,401 cycles for the healable system and 32,205 cycles for the reference system (see **Table 5**). This difference is justified by the chosen stress amplitude that corresponds to a region before the crossing

<figure><img src="image_362.png" /><figcaption>**FIGURE 10** | Relative secant modulus as a function of the number of cycles normalized by the number of cycles corresponding to a probability of 90% of survival for the healable and reference system. The vertical dashed lines correspond to each healing cycle.</figcaption></figure>

<table><caption>**TABLE 5** | Summary of the stop criteria for the fatigue/healing tests.</caption><thead><tr><th>System</th><th>σ<sub>a</sub><sup>eq</sup> (MPa)</th><th>σ<sub>a</sub><sup>record</sup> (MPa)</th><th>Stop crit. (# cycles)</th><th>Prob. survival (%)</th></tr></thead><tbody><tr><td>Healable</td><td>220</td><td>172</td><td>18,401</td><td>90</td></tr><tr><td>Reference</td><td>220</td><td>210</td><td>32,205</td><td>90</td></tr></tbody></table>

of the two S-N curves (see **Figure 1**). The change of the relative flexural secant modulus of each system as a function of the number of cycles is shown in **Figure 10**. The efficiency of the healing procedure is clearly observed for the healable system: when stopping the test at a 90% probability of survival, the sample could sustain 10 test/heal repetitions. Note however, that the healable system could not be tested in a row because of the extended testing time. The tests were stopped for 12 h after the 5th healing cycle, explaining the sudden peak of the relative modulus because of the longer time allowed for stress relaxation. The reference system on the other hand shows no improvement in durability from the healing cycle. Only a small regain of properties was observed after the first (and only) healing which was attributed to the relaxation of internal stresses barely affecting its durability. The secant modulus decrease of the reference sample is more marked after healing, indicating that the healing step did not increase the fatigue resistance of the reference system. Interesting to notice as well is the increase in relative modulus for the healable system just after healing, which is especially visible after the 6th healing cycle. This different behavior compared to the reference system can be attributed to the increased toughness of the healable system, artificially increasing the relative modulus by friction. These observations confirm the excellent efficiency of the healable system to recover its mechanical properties in fatigue.

## CONCLUSION

The potential of healable composites as an alternative to conventional composites was confirmed through 3-point bending fatigue tests. The tested healable system showed enhanced fatigue resistance compared to the reference system at high cycle fatigue, opening a door to the use of these materials in highly solicited structural parts. A damage criterion based on Weibull statistics was successfully derived and used to work out a probability of failure. Following this method, the life of the healable system could be extended by at least five times compared to that of the reference system by consecutive testing/healing cycles. Note that this approach could also be used to predict the failure of composite parts during their operation, or to organize the maintenance schedule for parts submitted to cyclic loading during their service. Even if the mechanical properties of this healable system show promising results, its processing under the present form led to a reduced fiber volume fraction compared to conventional systems. This study nonetheless demonstrates the high potential of this material system at equivalent volume fraction.
DATA AVAILABILITY STATEMENT

The raw data supporting the conclusions of this article will be
made available by the authors, without undue reservation.

AUTHOR CONTRIBUTIONS

NH, AC, and VM conceived the experiments, discussed the
methods and results, and prepared the manuscript. NH and
AC manufactured the test coupons and performed the
experimental test and analysis.

REFERENCES

Adam, T. J., and Horst, P. (2017). Fatigue damage and fatigue limits of a GFRP angle-ply laminate tested under very high cycle fatigue loading. *Int. J. Fatigue* 99, 202–214. doi:10.1016/j.ijfatigue.2017.01.045

ASTM International (2015). ASTM D7264. Standard test method for flexural properties of polymer matrix composite materials.

Blaiszik, B. J., Kramer, S. L. B., Olugebefola, S. C., Moore, J. S., Sottos, N. R., and White, S. R. (2010). Self-healing polymers and composites. *Annu. Rev. Mater. Res.* 40, 179–211. doi:10.1146/annurev-matsci-070909-104532

Bond, I. P., Trask, R. S., and Williams, H. R. (2008). Self-healing fiber-reinforced polymer composites. *MRS Bull.* 33, 770–774. doi:10.1557/mrs2008.164

Brown, E., White, S., and Sottos, N. (2005). Retardation and repair of fatigue cracks in a microcapsule toughened epoxy composite-part II: in situ self-healing. *Compos. Sci. Technol.* 65, 2474–2480. doi:10.1016/j.compscitech.2005.04.053

Cohades, A., Branfoot, C., Rae, S., Bond, I., and Michaud, V. (2018). Progress in self-healing fiber-reinforced polymer composites. *Adv. Mater. Interfaces* 5, 17. doi:10.1002/admi.201870084

Cohades, A., and Michaud, V. (2017). Thermal mending in E-glass reinforced poly(e-caprolactone)/epoxy blends. *Compos. Part A Appl. Sci. Manuf.* 99, 129–138. doi:10.1016/j.compositesa.2017.04.013

Deng, W., You, Y., and Zhang, A. (2015). "7-Supramolecular network-based self-healing polymer materials," in *Recent advances in smart self-healing polymers and composites*. Woodhead Publishing Series in Composites Science and Engineering. Sawston, United Kingdom: Wodhead Publishing, 181–210. doi:10.1016/B978-1-78242-280-8.00007-8

Gagel, A., Lange, D., and Schulte, K. (2006). On the relation between crack densities, stiffness degradation, and surface temperature distribution of tensile fatigue loaded glass-fibre non-crimp-fabric reinforced epoxy. *Compos. Part A Appl. Sci. Manuf.* 37, 222–228. doi:10.1016/j.compositesa.2005.03.028

Garcia, S. J. (2014). Effect of polymer architecture on the intrinsic self-healing character of polymers. *Eur. Polym. J.* 53, 118–125. doi:10.1016/j.eurpolymj.2014.01.026

Hamilton, A. R., Sottos, N. R., and White, S. R. (2012). Mitigation of fatigue damage in self-healing vascular materials. *Polymer* 53, 5575–5581. doi:10.1016/j.polymer.2012.09.050

Heo, Y., and Sodano, H. A. (2015). Thermally responsive self-healing composites with continuous carbon fiber reinforcement. *Compos. Sci. Technol.* 118, 244–250. doi:10.1016/j.compscitech.2015.08.015

Herman, W. A., Hertzberg, R. W., and Manson, J. A. (1990). The influence of loading history on fatigue in engineering plastics. *J. Mater. Sci.* 25, 434–440. doi:10.1007/bf00714052

JEC-GROUP (2017). *Overview of the global composites market: at the crossroads*. France: Frédérique Mutel, JEC Group.

Katnam, K. B., Comer, A. J., Roy, D., Da Silva, L. F. M., and Young, T. M. (2015). Composite repair in wind turbine blades: an overview. *J. Adhes.* 91, 113–139. doi:10.1080/00218464.2014.900449

Kessler, M. R., and White, S. R. (2001). Self-activated healing of delamination damage in woven composites. *Compos. Part A Appl. Sci. Manuf.* 32, 638–699. doi:10.1016/s1359-835x(00)00149-4

Khashaba, U. A. (2003). Fatigue and reliability analysis of unidirectional GFRP composites under rotating bending loads. *J. Compos. Mater.* 37 (4), 317–331. doi:10.1177/0021998303037004680

FUNDING

The initial research for this work was supported by the Swiss
National Fund (project no. 20020-15007-1) and later by the EPFL
Enable grants program.

ACKNOWLEDGMENTS

We would like to thank Hexion for providing the reference
epoxy resin.

Kim, S. Y., Sottos, N. R., and White, S. R. (2019). Self-healing of fatigue damage in cross-ply glass/epoxy laminates. *Compos. Sci. Technol.* 175, 122–127. doi:10.1016/j.compscitech.2019.03.016

Kling, S., and Czigány, T. (2014). Damage detection and self-repair in hollow glass fiber fabric-reinforced epoxy composites via fiber filling. *Compos. Sci. Technol.* 99, 82–88. doi:10.1016/j.compscitech.2014.05.020

Ladani, R. B., Wang, C. H., and Mouritz, A. P. (2019). Delamination fatigue resistant three-dimensional textile self-healing composites. *Compos. Part A Appl. Sci. Manuf.* 127, 105626. doi:10.1016/j.compositesa.2019.105626

Luterbacher, R., Trask, R. S., and Bond, I. P. (2016). Static and fatigue tensil properties of crossply laminates containing vasculs for self-healing applications. *Smart Mater. Struct.* 25, 015003. doi:10.1088/0964-1726/25/1/015003

Manfredi, E., Cohades, A., Richard, I., and Michaud, V. (2015). Assessment of solvent capsule-based self-healing for woven E-glass fibre-reinforced polymers. *Smart Mater. Struct.* 24, 1–11. doi:10.1088/0964-1726/24/1/015019

MarketsAndMarkets (2017). Report code CH 4906. Composite repairs market by type (structural, semi-structural, cosmetic), process (hand lay-up, vacuum infusion, autoclave), end-use industry (aerospace and defense, wind energy, automotive and transportation), and region—global forecast to 2026. doi:10.1016/j.ijfatigue.2017.12.015

Movahedi-Rad, A. V., Keller, T., and Vassilopoulos, A. P. (2018). Fatigue damage in angle-ply GFRP laminates under tension-tension fatigue. *Int. J. Fatigue* 109, 60–69. doi:10.1016/j.ijfatigue.2017.12.015

Norris, C. J., Bond, I. P., and Trask, R. S. (2013). Healing of low-velocity impact damage in vascularised composites. *Compos. Part A Appl. Sci. Manuf.* 44, 78–85. doi:10.1016/j.compositesa.2012.08.022

Park, J. S., Kim, H. S., and Thomas Hahn, H. (2009). Healing behavior of a matrix crack on a carbon fiber/mendomer composite. *Compos. Sci. Technol.* 69, 1082–1087. doi:10.1016/j.compscitech.2009.01.031

Patel, A. J., Sottos, N. R., Wetzel, E. D., and White, S. R. (2010). Autonomic healing of low-velocity impact damage in fiber-reinforced composites. *Compos. Part A Appl. Sci. Manuf.* 41, 360–368. doi:10.1016/j.compositesa.2009.11.002

Patrick, J. F., Hart, K. R., Krull, B. P., Diesendruck, C. E., Moore, J. S., and White, S. R. (2014). Continuous self-healing life cycle in vascularized structural composites. *Adv. Mater.* 26, 4302–4308. doi:10.1002/adma.201400248

Pingkarawat, K., Bhat, T., Craze, D. A., Wang, C. H., Varley, R. J., and Mouritz, A. P. (2013). Healing of carbon fibre-epoxy composites using thermoplastic additives. *Polym. Chem.* 4, 5007. doi:10.1039/c3py00459g

Pingkarawat, K., Wang, C. H., Varley, R. J., and Mouritz, A. P. (2012). Self-healing of delamination fatigue cracks in carbon fibre-epoxy laminate using mendable thermoplastic. *J. Mater. Sci.* 47, 4449–4456. doi:10.1007/s10853-012-6303-8

Pingkarawat, K., Wang, C. H., Varley, R. J., and Mouritz, A. P. (2014). Healing of fatigue delamination cracks in carbon-epoxy composite using mendable polymer stitching. *J. Intell. Mater. Syst. Struct.* 25, 75–86. doi:10.1177/1045389X13505005

Post, W., Cohades, A., Michaud, V., van der Zwaag, S., and Garcia, S. J. (2017). Healing of a glass fibre reinforced composite with a disulphide containing organic-inorganic epoxy matrix. *Compos. Sci. Technol.* 152, 85–93. doi:10.1016/j.compscitech.2017.09.017

Préau, M., and Hubert, P. (2016). Processing of co-bonded scarf repairs: void reduction strategies and influence on strength recovery. *Compos. Part A Appl. Sci. Manuf.* 84, 236–245. doi:10.1016/j.compositesa.2016.01.016
Sakin, R., and Ay, İ. (2008). Statistical analysis of bending fatigue life data using Weibull distribution in glass-fiber reinforced polyester composites. *Mater. Design* 29, 1170–1181. doi:10.1016/j.matdes.2007.05.005

Shabani, P., Shokrieh, M. M., and Saeedi, A. (2020). A novel model to simulate the formation and healing of cracks in self-healing cross-ply composites under flexural loading. *Compos. Struct.* 235, 111750. doi:10.1016/j.compstruct.2019.111750

Silva, J. M. A., Devezas, T. C., Silva, A. P., and Ferreira, J. A. M. (2005). Mechanical characterization of composites with embedded optical fibers. *J. Compos. Mater.* 39 (14), 1261–1281. doi:10.1177/0021998305050423

Sordo, F., and Michaud, V. (2016). Processing and damage recovery of intrinsic self-healing glass fiber reinforced composites. *Smart Mater. Struct.* 25, 084012. doi:10.1088/0964-1726/25/8/084012

Tsangouri, E., Aggelis, D., and Van Hemelrijck, D. (2015). Quantifying thermoset polymers healing efficiency: a systematic review of mechanical testing. *Prog. Polym. Sci.* 49–50, 154–174. doi:10.1016/j.progpolymsci.2015.06.002

Ye, X. J., Zhu, Y., Yuan, Y. C., Song, Y. X., Yang, G. C., and Rong, M. Z. (2017). Improvement of fatigue resistance of epoxy composite with microencapsulated

epoxy-SbF5 self-healing system. *Express Polym. Lett.* 11, 853–862. doi:10.3144/expresspolymlett.2017.82

Zhu, Y., Ye, X. J., Rong, M. Z., and Zhang, M. Q. (2016). Self-healing glass fiber/ epoxy composites with polypropylene tubes containing self-pressurized epoxy and mercaptan healing agents. *Compos. Sci. Technol.* 135, 146–152. doi:10.1016/j.compscitech.2016.09.020

**Conflict of Interest:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2020 Hostettler, Cohades and Michaud. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
<figure><img src="image_363.png" /></figure>

Check for
updates

# Supramolecular Structure and Mechanical Properties of Wet-Spun Polyacrylonitrile/Carbon Nanotube Composite Fibers Influenced by Stretching Forces

Hamideh Mirbaha<fnref n="1" />,²*, Paolo Scardi³, Mirco D'Incau³, Shahram Arbab⁴, Parviz Nourpanah² and Nicola M. Pugno<fnref n="1" />,⁵*

**OPEN ACCESS**

**Edited by:**
Patricia Krawczak,
IMT Lille Douai, France

**Reviewed by:**
Yaodong Liu,
Institute of Coal Chemistry (CAS),
China
Wei Xu,
IMDEA Nanociencia, Spain

**\*Correspondence:**
Hamideh Mirbaha
hamideh.mirbaha@unitn.it
Nicola M. Pugno
nicola.pugno@unitn.it

**Specialty section:**
This article was submitted to
Polymeric and Composite Materials,
a section of the journal
Frontiers in Materials

**Received:** 29 April 2020
**Accepted:** 22 June 2020
**Published:** 24 July 2020

**Citation:**
Mirbaha H, Scardi P, D'Incau M, Arbab S, Nourpanah P and Pugno NM (2020) Supramolecular Structure and Mechanical Properties of Wet-Spun Polyacrylonitrile/Carbon Nanotube Composite Fibers Influenced by Stretching Forces. *Front. Mater. 7:226*. doi: 10.3389/fmats.2020.00226

The effect of different elongation conditions on the crystalline structure and physical and mechanical properties of polyacrylonitrile/carbon nanotube (PAN/CNT) microfibers during the wet spinning process was studied. It turns out that the response of polymer chains in PAN/CNT and in PAN fibers to the stretching forces from jet stretching and steam drawing is different. The X-ray diffraction (XRD) results showed that the crystalline domain size in steam-drawn PAN/CNT fibers is 1.5 times larger than in PAN fibers. CNTs alter the optimum stretching conditions, as they improve the crystalline structure of the PAN/CNT fibers at lower steam drawing ratios than PAN fibers, through nucleation of crystals on their surface. Synchrotron-radiation XRD studies revealed that the presence of CNTs improves the crystal orientation of PAN/CNT fibers significantly. In addition, steam drawing is more effective in improving the crystal orientation than jet stretching. The mechanical properties of PAN/CNT fibers have also been affected by steam drawing more than jet stretching. Multiwalled CNTs have the biggest impact on Young's modulus. The Young's modulus of PAN/CNT fibers could increase up to 19% higher than PAN fibers at specific stretching conditions, i.e., steam drawing ratio of 2.5. Better orientation of polymers and crystals in the fiber direction is the reason for the enhancement of Young's modulus. To our knowledge, the differences between the response of PAN/CNT and PAN fibers to stretching forces inside coagulation bath and after fiber coagulation as well as the difference in evolution of crystalline structure at different stretching stages has not been reported elsewhere.

**Keywords:** PAN/CNT composite fibers, wet spinning, jet stretch, steam drawing, crystalline structure, mechanical properties

<footnote n="1">¹ Laboratory of Bio-Inspired, Bionic, Nano, Meta Materials & Mechanics, Department of Civil, Environmental and Mechanical Engineering, University of Trento, Trento, Italy, ² Department of Textile Engineering, Amirkabir University of Technology, Tehran, Iran, ³ Department of Civil, Environmental and Mechanical Engineering, University of Trento, Trento, Italy, ⁴ Advanced Textile Materials and Technology Research Institute, Department of Textile Engineering, Amirkabir University of Technology, Tehran, Iran, ⁵ School of Engineering and Materials Science, Queen Mary University of London, London, United Kingdom</footnote>
INTRODUCTION

Incorporation of CNTs in PAN microfibers has the potential to improve the physical, mechanical, electrical, and thermal properties of PAN fibers. This can be achieved without any significant weight change and or need to use different production equipment (Brown et al., 2014; Sahin et al., 2014; Li et al., 2018; Mirbaha et al., 2019). These fibers have the potential to be used as a carbon fiber precursor (Mikolajczyk et al., 2010; Sahin et al., 2014). By addition of CNTs, Li et al. (2015) obtained 35 and 17% improvement in tensile strength and tensile modulus of PAN fibers, respectively. To have PAN/CNT fibers with better mechanical properties, the internal structure of fibers needs to be controlled and designed (Ji et al., 2007). The formation of a fiber structure in the wet spinning process is a very important step in determining the properties of produced fibers (Wilms et al., 2013; Wei et al., 2020). Since there are many effective parameters on the evolution of a fiber structure in the wet spinning stage, the favored production of PAN/CNT composite fibers with promising properties involves thorough research on the spinning process, which is also the main aim of this work.

PAN/CNT composite fibers can be produced by the wet spinning method (Mirbaha et al., 2019). During the wet spinning process, shear and elongational forces are applied on filaments, affecting the alignment of polymer chains (Mai et al., 2011; Dintcheva et al., 2011). Shear stress is applied on the spinning dope while it flows inside the spinneret (Mikolajczyk et al., 2010). According to Cao et al. (2004), a strong enough shear force can align polymer chains in the flow direction. The melt spinning process of polypropylene/CNT composite fibers was studied by Kang and Ryu (2013). They observed that shear stress applied to the spinning dope can also be transferred to CNTs, resulting in their orientation in the flow direction. Minus et al. (2012) observed that CNTs can act as nucleating agents for the crystallization of polyethylene. A small amount of CNT dispersed within a polymer can influence its orientation and crystallinity.

Stretching is one of the most necessary steps in fiber production, leading to improvement in molecular orientation along the fiber direction as well as fineness, tensile strength, and modulus (Zeng et al., 2007; Arbab et al., 2011a). The study on the effect of elongational forces on the initial structure of the PAN fibers in the presence of CNTs can help control the alignment of polymers and CNTs, and the formation of an interphase between them, subsequently playing an important role on the final properties of the fiber (Lin et al., 2019). Fiber stretching can be done at different production stages including jet stretching inside the coagulation bath and hot drawing of coagulated fibers in different environments (Arbab et al., 2011a; Lian et al., 2012). Jet stretching reduces the residence time of filaments in the coagulation bath, decreasing the available time for formation and growth of voids (Arbab et al., 2011b). Arbab et al. (2011b) obtained a denser structure with smaller voids, when they increased the jet stretch inside the coagulation bath of the wet spinning process of PAN fibers. Li et al. (2015) observed that drawing has two opposite effects on the crystallization

of PAN fibers. Applying stretch results in orientation assisted crystallization and crystallinity improvement. On the other hand, increasing the stretch ratio shortens the crystallization step and reduces the crystallinity. According to Cao et al. (2004), the stretching force inside the coagulation bath is able to unwind the random coils inside fiber filaments, stretch the chains, improve the chain orientation, and increase the compactness of the structure. While shear forces inside the spinneret can cause the alignment and orientation of polymer chains, the coexistence of shear and elongational forces results in stretching and orientation of polymer chains and opening of random coils (Cao et al., 2004). According to Mai et al. (2011), polymer chains and CNTs are likely to change from an entangled shape to a stretched form under elongational force. Polymer chains and CNTs have a greater chance of interaction in their stretched form. Polymer chains can crystallize around CNTs, thus producing a better interface interaction. Mai et al. (2011) reported better dispersion of CNTs at higher stretch ratios during melt spinning of PA66 as well as better orientation of fibers in the presence of CNTs.

After formation of a fiber structure in the coagulation bath, it is possible to improve the orientation and fiber properties by hot drawing the as-spun fibers (Zeng et al., 2007). The time scale of the reaction of polymer chains to hot drawing of coagulated fibers is different from the jet stretching of the polymer fluid (Zeng et al., 2007). During the hot drawing stage, the orientation of a fibrillar network and polymer chains in amorphous and crystalline regions increases, improving the mechanical strength (Arbab et al., 2011b).

Most of the researches in the field of PAN/CNT composite fiber production have focused on the structure and properties of the final products, while studies on the evolution of the structure during the production phases are still limited. The coexistence of shear and elongation in the wet spinning stage alters the orientation and placement of the polymer chains and CNTs and determines the initial fiber structure that can be developed during the next production stages. In a previous work, we reported the different effects of shear rates and jet stretching ratios inside the coagulation bath – where the filaments are still not fully solidified – on the morphology, crystalline structure, and physical and mechanical properties of as-spun PAN/CNT fibers (Mirbaha et al., 2019). In this work, we kept the shear rate constant, while applying a new stretching step in water vapor after full fiber coagulation. The objective of this research is the exploration of the effect of elongational forces at different stretching conditions on the formation and evolution of the PAN/CNT fiber structure and properties affected by the presence of CNTs. Understanding the effect of the mentioned variables can be regarded as a scheme part to control the spinning process and improve the physical and mechanical properties of the final PAN/CNT fibers. For this reason, the effect of jet stretching and hot drawing on the structure and properties of PAN/CNT composite fibers is studied using physical and mechanical measurements, X-ray diffraction, and FESEM microscopy. The impact of the existence of CNTs inside fibers on structural development is also explored in comparison with PAN reference fibers.
# EXPERIMENTAL

## Materials

PAN powder with methyl acrylate and sodium methallyl sulfonate comonomers was obtained from Polyacryl Corp., Iran. The viscosity molecular weight was 112,000 g/mole. Analysis-grade dimethyl formamide (DMF) solvent was obtained from Merck Co., Germany, and was used as received. Multiwalled CNTs were obtained from the Research Institute of Petroleum Industry, Iran. The average length and diameter of nanotubes are 10 micron and 10–30 nanometers, respectively. The purity was over 95%.

## Methods

### Preparation of the Spinning Dope

The spinning dope for reference PAN fibers was prepared from dissolving PAN powder in DMF solvent at a concentration of 20% w/v. The spinning dope for PAN/CNT composite fibers was prepared by dissolving PAN powder in DMF solvent at a concentration of 20% w/v, as in the spinning dope for PAN fibers. Based upon the literature review and optimization tests, the CNT concentration of 0.75% wt. proportional to the weight of the PAN polymer was chosen. For production of 35 cc of polymer solution, 0.0525 g CNT was added to 400 mL of DMF solvent. CNTs were dispersed for 30 min in DMF solvent using a UP200S ultrasonic homogenizer (Hielscher Ultrasonics GmbH, Germany). Then, a mechanical stirrer was placed in the CNT/DMF dispersion and polymer solution was continuously added to the dispersion under shear force until a homogenous solution was attained. The composite solution was further homogenized for 30 min using a ultrasonic homogenizer. An optical image of the final state of composite solution after ultrasonication is shown in **Figure 1**. Using mild heating and shear force, the extra solvent was evaporated and the final PAN/CNT/DMF solution with PAN and CNT concentrations of 20 and 0.75%, respectively, was attained. For determination of the end point of solvent evaporation, a finely scaled narrow glass container was used and the volume reduction during evaporation was closely observed until the final volume of the spinning solution was achieved.

### Wet Spinning

A laboratory wet spinning machine was used to produce composite PAN/CNT and reference PAN fibers. A schematic diagram of this wet spinning machine is reported in our previous paper (Mirbaha et al., 2019). A metering pump was used to pump the spinning dope at the constant shear rate of 8,000 s⁻¹ to the coagulation bath through a one-hole spinneret with a diameter of 200 microns. The coagulation bath contained water as non-solvent and the temperature was kept at 20°C to reduce the phase separation speed. The freely spun fibers were produced without any stretching force applied on the filaments. These fibers were collected from the coagulation bath and then were thoroughly washed and dried at room temperature. Jet-stretched fibers were picked up at take-up rollers. The take-up velocity to the extrusion velocity of the spinning dope is defined as the jet stretch ratio. Jet stretch ratios of 2 and 2.5 were applied

on fibers by changing the take-up speed in relation to the extrusion velocity. Jet-stretched fibers were washed thoroughly with water, and they were dried naturally in air. The hot drawing stage was performed in water vapor at 100°C in a laboratory-designed drawing machine (Figure 2). Fibers were fixed between two steel grips inside the hot steam chamber and were drawn up to different drawing ratios with constant speed of 1.3 mm/sec. The drawn fibers were then dried at room temperature.

The sample names and production conditions of PAN/CNT and PAN fibers are listed in Table 1. The letters of P and PC in sample names represent PAN and PAN/CNT, respectively. The numbers coming after the letters of S and D are jet-stretch ratios inside the coagulation bath and steam-draw ratios, respectively. The fiber codes without these numbers are freely spun fibers. Steam-draw ratios are defined as the final length of fibers after drawing in hot water vapor at 100°C to the initial length of fibers before drawing. The freely spun and jet-stretched fibers are similar to those reported in our previous work (Mirbaha et al., 2019). They are reported in this work for comparison to steam-drawn fibers.

<figure><img src="image_364.png" /><figcaption>**FIGURE 1** | An optical image of the composite solution after sonication and mechanical stirring.</figcaption></figure>
<figure><img src="image_365.png" /><figcaption>**FIGURE 2** | Schematic diagram of the steam drawing chamber used in this research.</figcaption></figure>

<table><caption>**TABLE 1** | Denotation and production conditions of PAN/CNT and PAN fibers.</caption><thead><tr><th rowspan="2">Condition</th><th colspan="2">Sample names</th><th rowspan="2">Jet stretch ratio</th><th rowspan="2">Jet stretch ratio</th></tr><tr><th>PAN/CNT</th><th>PAN</th></tr></thead><tbody><tr><td>Freely spun</td><td>PC</td><td>P</td><td>-</td><td>-</td></tr><tr><td rowspan="2">Jet stretched</td><td>PC-S2</td><td>P-S2</td><td>2</td><td>-</td></tr><tr><td>PC-S2.5</td><td>P-S2.5</td><td>2.5</td><td>-</td></tr><tr><td rowspan="3">Steam srawn</td><td>PC-D2.5</td><td>P-D2.5</td><td>-</td><td>2.5</td></tr><tr><td>PC-D5</td><td>P-D5</td><td>-</td><td>5</td></tr><tr><td>PC-S2-D5</td><td>P-S2-D5</td><td>2</td><td>5</td></tr></tbody></table>

## Characterization

In order to evaluate the dispersion of CNTs and determine the optimum dispersion conditions, UV-visible analysis was utilized. For this purpose, the PAN/DMF solution with a concentration of 20% and the CNT/DMF suspension with a CNT concentration of 0.75% proportional to the weight of the PAN polymer were prepared. Similar to preparation of the spinning dope, the CNT/DMF solution was first ultrasonicated for 30 min. Then, by the help of a mechanical stirrer, the polymer solution was added to the CNT/DMF dispersion under shear force. The dispersion process was again continued for 15, 30, 60, and 75 min. At each of these times, a sample was collected, diluted, and tested in a Cary 100 UV-visible spectrophotometer (Agilent, United States) in a wave number ranging from 200 to 700 nm.

A Microphot FXA optical microscope (Nikon, Japan) at a magnification of 50× was used to measure the fiber diameter at several points along the fiber. The average values of different replicates are reported here. A GR-200 balance (A and D Balance, Japan) was used to measure the mass of different fibers at various lengths, and the average value of linear density was calculated and reported in unit of tex as the mass in grams of 1,000 meters of fibers. The overall porosity of fibers was approximately calculated

as follows (Arbab et al., 2008):

$$ \text{Overall porosity (\%)} = \frac{\pi R^2 L - m \rho^{-1}}{\pi R^2 L} \times 100 \quad (1) $$

where R, L, and m are the radius, length, and mass of fibers, respectively, and ρ is the polymer density.

The microscopic images of fiber cross sections at magnifications of 500×, 7,000×, and 35,000× were obtained using a MIRA3 FESEM (Tescan, Czech Republic). The operation voltage was 15 kV. The samples were coated by gold, and the intact cross sections were cut in liquid nitrogen.

A Panalytical X'Pert MRD X-ray diffractometer (United Kingdom) was utilized to study the crystalline parameters of PAN/CNT and PAN fibers. The diffractometer worked at an operation voltage of 40 kV and current of 40 mA. Co Kα radiation at wavelength of 1.790 Å was used. Fibers were placed as bundles on the sample holder. The scanning range was 10°–48° (2θ), the scanning step was 0.15°, and the counting time was 25 s/step. To extract data from X-ray patterns, the diffraction peaks were resolved using Origin 2017 software (OriginLab, United States) based on three Lorentzian functions. The crystalline parameters were calculated from the intense peak at 2θ~17°. Bragg law was used to calculate the plane spacing (d), and lateral crystallite size (Lc) was estimated using the Scherrer equation (Ji et al., 2007; Gong et al., 2014):

$$ L_c = \frac{K\lambda}{\beta\cos\theta} \quad (2) $$

where K is a constant with a value of 0.89 and β is the full width at half maximum (FWHM) intensity of the peak.

The degree of crystallinity was assessed using Gupta–Singhal’s method (Ouyang et al., 2015):

$$ \text{Crystallinity} = \frac{A_{\text{crystal}}}{A_{\text{total}}} \times 100 \quad (3) $$

where Acrystal is the area of crystalline peaks and Atotal is the total area of diffraction after subtracting the background.

Synchrotron X-ray patterns were collected at the MCX beamline of Elettra synchrotron in Trieste, Italy (Rebuffi et al., 2014). A beamline energy of 17.5 keV was used. The single fiber was placed horizontally with its axis perpendicular to X-ray radiation beam. Therefore, the vertical and horizontal directions in two-dimensional (2D) patterns are related to equatorial and meridional directions of fibers, respectively. 2D X-ray diffraction patterns were integrated to two one-dimensional (1D) equatorial and meridional scans. The intensity and area under the peak were calculated by modeling with a Python script. The modeling is made using a decaying exponential background and a pseudo-Voigt peak profile.

A universal testing machine MIDI 10 (Messphysik Materials Testing GmbH, Austria) was used to measure the mechanical properties of single threads of PAN and PAN/CNT fibers. A 10-N load cell was utilized. The initial distance between two steel grips was 1 cm. An elongation rate of 0.1 mm/s based on Standard D3822 was chosen. The mechanical properties reported in this paper are the average values of at least 10 test replicates. The
coefficient of variation (% CV) was calculated for each set of tests. The breaking tenacity was measured from the load at break divided by the linear density of fibers in tex. Young's modulus was measured from the slope of the initial linear section of a specific stress-strain curve. The mechanical properties are reported in cN/tex (force divided by mass per unit length) instead of MPa, as the cross section of some fibers cannot be easily defined. Considering 1 tex = 0.001 g/m, tensile data in cN/tex (force per unit linear density) can be converted to Pa (force per unit area of cross section) by multiplying by the density of the fiber in proper units (Ward, 1997).

# RESULTS AND DISCUSSION

## Determination of Optimum Dispersion Time

CNTs tend to form aggregates in solution due to Van der Waals forces. Therefore, one of the challenges for producing PAN/CNT composite fibers is dispersion of CNTs in the form of individual tubes or in small bundles inside the viscous polymer matrix (Basu-Dutt et al., 2012). At these conditions, the interface area between CNT and polymer increases, leading to improvement in fiber mechanical properties.

The addition of polymer to dispersion solution after 30 min of sonication is due to the fact that polymer chains can be adsorbed on the surface of CNTs and penetrate to their bundles with the help of ultrasonic force (Bakhtiar Davijani and Kumar, 2015). Therefore, CNT bundles can become smaller. As the dispersion time increases, the PAN/CNT/DMF dilute solution becomes darker, while the aggregations that can be recognized with the naked eye decrease. These aggregations completely disappear at dispersion times of 60 min or longer (Figure 3A).

UV-visible spectra of PAN/CNT/DMF solutions contain one absorption peak at a wavelength of about 265 nm (Figure 3B). Since only individual CNTs can absorb UV and visible lights, the maximum absorbance during dispersion can be used as a method to evaluate the dispersion state of CNTs (Sapalidis et al., 2018). The maximum UV-visible absorbance increases with an increase in dispersion time (Figure 3B). At dispersion time of 60 min or longer, the maximum absorbance stays constant, revealing that maximum possible dispersion of CNTs is achieved. Increasing the dispersion time to longer than 60 min not only does not improve the dispersion state of CNTs in the solution but also increases the energy consumption as well as the possibility of undesired damages to the CNTs. As a result, dispersion period of 60 min was used in this research.

## Physical Properties

The existence of CNTs among PAN polymer chains causes a decrease in the fiber diameter in all spinning and stretching conditions, i.e., freely spun, jet-stretched, and steam-drawn fibers (Table 2). The diameter decrease is attributed to the smaller die swell as a result of the effect of CNTs on lowering the elastic recovery of polymer chains (Mirbaha et al., 2019). CNTs decrease the elastic recovery of polymer chains by reducing the chain mobility and formation of the interphase between polymers

and CNTs. Consequently, polymer chains can preserve their orientation better in the presence of CNTs.

At identical stretch ratios, the reduction in diameter and linear density after steam drawing is less than that of jet stretch inside the coagulation bath (Table 2). For example, the diameter of steam-drawn PAN/CNT fibers at the ratio of 2.5 is about 1.5 times larger than that of jet-stretched PAN/CNT fibers at the ratio of 2.5.

The porosity of PAN fibers increases when CNT exists in the structure. The high porosity of fiber samples shows that the spinning conditions need further optimization. The overall fiber porosity is attributed to macrovoids (finger-like voids between fibrillar parts) as well as nanovoids (among fibrils) (Arbab et al., 2008). The more voids are present in the structure, the more random and disordered will be the arrangement of molecular chains, leading to less fiber orientation (Ji et al., 2007). The biggest difference between porosity of PAN/CNT and PAN fibers is observed when fibers are jet stretched (Table 2). For PAN reference fibers, jet stretch has brought structural elements close to each other reducing the voids, whereas for PAN/CNT fibers the void volume has increased by application of jet stretch. In addition, the presence of CNTs among polymer chains and formation of interphase with polymers has probably prevented from free movement of chains, intensifying the formation of voids under stretching forces inside coagulation bath. The higher porosity of PAN/CNT fibers compared to PAN fibers could also possibly be due to increased solution elasticity. The smallest difference between porosity of PAN/CNT and PAN fibers is observed when fibers are spun under steam drawing. Steam drawing has increased the porosity of PAN fibers sharply compared to jet-stretched fibers, whereas the difference between porosity of jet-stretched and steam-drawn PAN/CNT fibers is negligible. These results show the different response of PAN/CNT and PAN fibers to jet stretch and steam drawing. The lowest porosity of PAN/CNT fibers is obtained when fibers are jet stretched at the ratio of 2 and then steam drawn at the ratio of 5. The combination of jet stretch and steam drawing results in diameter reduction, formation of smaller voids during coagulation, and narrowing and healing of the voids in water steam, leading to a more compact structure.

The porosity of P-S2-D5 has increased compared to P-S2 fibers, whereas the opposite trend is observed in PC fibers. Jet stretching has brought the structural parameters in PAN fibers close to each other, reducing the voids and porosity of structure. However, jet stretching at high ratios used in this research, which turned out to be higher than optimum values for PAN/CNT fibers, has increased the porosity of PAN/CNT fiber structure. Steam drawing at temperatures higher than glass transition temperatures has been more successful in healing the pores and voids in the structure of PAN/CNT fibers. This may be due to lower crystallinity of the structure of jet-stretched PAN/CNT fibers (as it will be discussed in section "Laboratory X-Ray Diffraction"), which makes the structural arrangement and orientation easier than PAN fibers. In PAN fibers, however, steam drawing at the temperature and rate used in this research has not reduced the porosity of structure. This might result from chain breakage during steam drawing and/or slippage of
<figure><img src="image_366.png" /><figcaption>**FIGURE 3** | (A) Dilute PAN/CNT/DMF solutions after ultrasonication at different times. (B) UV-visible spectrum of PAN/CNT/DMF solutions at different ultrasonication times.</figcaption></figure>

some fibrils and crystalline domains to rearrange and reorganize new domains, introducing new voids in the structure. Liu et al. (2015) also observed an increase in porosity at some hot drawing ratios and lower drawing temperatures during the production of ultra-high molecular weight polyethylene fibers.

## Fiber Morphology

During the coagulation process of PAN fibers, a structure consisting of loose fibrils and finger-like voids along the radius of the fiber is formed. This structure is formed because of counter diffusion of solvent/non-solvent and rapid phase separation (Ouyang et al., 2011). Finger-like voids reduce the mechanical properties of the fibers (Arbab et al., 2008). Application of jet-stretch ratios of 2 and 2.5 to PAN/CNT fibers makes the cross

section irregular with deep surface grooves (Figure 4 in Mirbaha et al., 2019). On the other hand, application of steam drawing with the ratio of 5 to jet-stretched PAN/CNT fibers heats the surface grooves to a large extent, Figure 4A. The cross section of these fibers is also more regular compared to jet-stretched PAN/CNT fibers (Mirbaha et al., 2019). The discrepancy of the shape of the cross section from a circular form affects the mechanical properties of the fibers (Chen et al., 2010). PAN reference fibers at jet stretch ratios of 2 and 2.5 have a circular cross section (Figure 4 in Mirbaha et al., 2019). Application of a steam drawing ratio of 5 to these fibers also does not change the cross section shape significantly (Figure 4B). This difference suggests that the tolerability of PAN/CNT fibers to jet stretch is less than that of PAN fibers. The voids in the cross section of steam-drawn fibers are shrunk as a result of elongational forces. In addition, the fibrils are more oriented and compact after steam drawing. Fibrillar compactness of PAN/CNT fibers is more than that of PAN fibers. This is attributed to the smaller die swell in PAN/CNT fibers (Mirbaha et al., 2019). As it was reported in our previous paper (Mirbaha et al., 2019), we could not identify aggregations of CNTs in fiber cross sections even at magnifications of 150,000×. This can possibly show the good state of dispersion of CNT in the matrix and wrapping individual CNTs with polymer chains (Mirbaha et al., 2019).

<table><caption>**TABLE 2** | Physical properties of PAN/CNT and PAN fibers at different spinning conditions.</caption><thead><tr><th>Condition</th><th>Sample name</th><th>Diameter (μm)</th><th>% CV</th><th>Linear density (tex)</th><th>% CV</th><th>Porosity (%)</th></tr></thead><tbody><tr><td colspan="7"><strong>PAN/CNT</strong></td></tr><tr><td>Freely spun</td><td>PC</td><td>275.0</td><td>0.04</td><td>15.1</td><td>0.05</td><td>78.5</td></tr><tr><td rowspan="2">Jet stretched</td><td>PC-S2</td><td>147.9</td><td>0.07</td><td>4.0</td><td>0.09</td><td>80.3</td></tr><tr><td>PC-S2.5</td><td>138</td><td>0.13</td><td>2.8</td><td>0.16</td><td>84.1</td></tr><tr><td rowspan="3">Steam Drawn</td><td>PC-D2.5</td><td>197.6</td><td>0.04</td><td>6.3</td><td>0.07</td><td>82.6</td></tr><tr><td>PC-D5</td><td>157</td><td>0.06</td><td>3.3</td><td>0.08</td><td>85.5</td></tr><tr><td>PC-S2-D5</td><td>77.4</td><td>0.15</td><td>1.1</td><td>0.19</td><td>79.1</td></tr><tr><td colspan="7"><strong>PAN</strong></td></tr><tr><td>Freely spun</td><td>P</td><td>335.4</td><td>0.05</td><td>28.2</td><td>0.06</td><td>73.1</td></tr><tr><td rowspan="2">Jet stretched</td><td>P-S2</td><td>185.5</td><td>0.06</td><td>10.0</td><td>0.05</td><td>68.7</td></tr><tr><td>P-S2.5</td><td>164.2</td><td>0.11</td><td>7.3</td><td>0.05</td><td>70.7</td></tr><tr><td rowspan="3">Steam Drawn</td><td>P-D2.5</td><td>241.4</td><td>0.05</td><td>12.4</td><td>0.06</td><td>77.2</td></tr><tr><td>P-D5</td><td>196.2</td><td>0.05</td><td>6.8</td><td>0.07</td><td>81.1</td></tr><tr><td>P-S2-D5</td><td>102</td><td>0.04</td><td>2.1</td><td>0.06</td><td>76.1</td></tr></tbody></table>

## Crystalline Structure

### Laboratory X-Ray Diffraction

A sharp peak at 2θ~17° exists in the X-ray pattern of PAN/CNT fibers, which is attributed to the diffraction peak from the (100) planes (Ouyang et al., 2015). A broad diffuse scattering appearing around 2θ~20~30° is attributed to the contribution of amorphous regions of fibers (Figure 5). These peaks in freely spun fibers are wide and weak. Application
<figure><img src="image_367.png" /><figcaption>**FIGURE 4** | FESEM images of the cross section and morphology of PAN/CNT and PAN fibers at magnifications of 500×, 7,000×, and 35,000×. (A) PC-S2-D5, and (B) P-S2-D5.</figcaption></figure>
of jet stretch increases the peak intensity to some extent, whereas steam drawing sharpens the diffraction peak at 2θ~17° and increases its intensity significantly. In addition, a small crystalline peak appears at 2θ~30°, showing an improvement in the crystalline structure of PAN/CNT fibers with steam drawing.

During steam drawing, polymer chains are oriented and packed in a way that orientation-induced crystallization can take place. As a result, the crystallinity and crystallite size increase sharply (Table 3). Application of steam drawing at the ratio of 2.5 has better efficiency in improving the crystalline structure of the freely spun PAN/CNT fibers compared to a total draw ratio of 10 (Table 3), whereas the improvement in the crystalline structure of PAN reference fibers at a steam drawing ratio of 2.5 is negligible. For improving the crystalline structure of PAN fibers and increasing their crystallite size in the absence of CNTs, application of both jet stretch and steam drawing at higher drawing ratios is necessary. To our

knowledge, the difference between the response of polymer chains in the composite PAN/CNT and reference PAN fibers to the stretching force inside the coagulation bath and after fiber coagulation as well as the difference in evolution of the crystalline structure at different stretching steps has not been reported elsewhere. This information can help optimize the stretching stages of PAN/CNT fibers to obtain better fiber properties.

Comparison of crystalline parameters of PC-S2.5 and PC-D2.5 shows the larger influence of steam drawing on the crystalline structure of PAN/CNT fibers compared to jet stretching (Table 3). Steam drawing has resulted in the growth of crystals and significant improvement in fiber crystallinity. As reported by Gissinger et al. (2018), this can improve the interfacial shear strength of CNTs in the matrix and result in stronger cohesion.

In the presence of the water vapor molecules at the temperature of 100°C, polymer chains have more freedom for mobility and alignment. In addition, there is more intense nucleation and growth of crystals in the presence of CNTs, whereas PAN reference fibers show a different behavior.

Crystallite size and average crystallite layers of PAN/CNT fibers are larger than those of corresponding PAN fibers at most spinning conditions (Figure 6A). The presence of CNTs among the polymer chains facilitates the growth of crystallites by joining smaller crystallites together and joining more chains from the amorphous and transition phases to the crystal phase. We propose that nucleation and growth of PAN crystals on the surface of CNTs has resulted in improvement in the crystalline domain size. Other researchers also reported the growth of crystals in the vicinity of CNTs (Mai et al., 2011; Zhang et al., 2013).

Despite larger crystallites, the crystallinity of PAN/CNT fibers is less than that of PAN reference fibers at most spinning conditions (Figure 6B). It shows the presence of larger crystals at fewer numbers in PAN/CNT composite fibers. Jet-stretched PAN/CNT fibers have significantly lower crystallinity than corresponding PAN fibers, whereas the crystallinity of PAN/CNT fibers increases sharply with steam drawing, bringing it close to or even higher than that of PAN fibers. This shows the larger

<figure><img src="image_368.png" /><figcaption>TABLE 3 | Parameters of crystalline structure of PAN/CNT and PAN fibers at different stretching conditions.</figcaption></figure>

<table><thead><tr><th>Condition</th><th>Sample name</th><th>Angle (2θ)</th><th>FWHM (2θ)</th><th>d (Å)</th><th>L<sub>c</sub> (nm)</th><th>L<sub>c</sub>/d</th><th>C (%)</th></tr></thead><tbody><tr><td colspan="8"><strong>PAN/CNT fibers</strong></td></tr><tr><td>Freely spun</td><td>PC</td><td>17.09</td><td>3.07</td><td>5.185</td><td>2.6</td><td>5.0</td><td>13.0</td></tr><tr><td rowspan="2">Jet stretched</td><td>PC-S2</td><td>16.93</td><td>2.19</td><td>5.233</td><td>3.6</td><td>6.9</td><td>34.1</td></tr><tr><td>PC-S2.5</td><td>16.76</td><td>2.10</td><td>5.286</td><td>3.8</td><td>7.1</td><td>36.5</td></tr><tr><td rowspan="2">Steam drawn</td><td>PC-D2.5</td><td>16.98</td><td>1.29</td><td>5.218</td><td>6.1</td><td>11.8</td><td>54.5</td></tr><tr><td>PC-S2-D5</td><td>16.98</td><td>1.15</td><td>5.216</td><td>6.9</td><td>13.3</td><td>58.4</td></tr><tr><td colspan="8"><strong>PAN fibers</strong></td></tr><tr><td>Freely spun</td><td>P</td><td>17.02</td><td>3.10</td><td>5.206</td><td>2.6</td><td>4.9</td><td>13.5</td></tr><tr><td rowspan="2">Jet stretched</td><td>P-S2</td><td>16.91</td><td>2.32</td><td>5.240</td><td>3.4</td><td>6.5</td><td>53.9</td></tr><tr><td>P-S2.5</td><td>16.82</td><td>2.29</td><td>5.266</td><td>3.5</td><td>6.6</td><td>54.6</td></tr><tr><td rowspan="2">Steam drawn</td><td>P-D2.5</td><td>16.86</td><td>2.16</td><td>5.255</td><td>3.7</td><td>7.0</td><td>53.2</td></tr><tr><td>P-S2-D5</td><td>16.97</td><td>1.03</td><td>5.222</td><td>7.7</td><td>14.7</td><td>60.5</td></tr></tbody></table>

FWHM, Full width at half maximum; d, interplanar spacing; L<sub>c</sub>, crystalline domain size; L<sub>c</sub>/d, average crystallite layers; C, crystallinity.
<figure><img src="image_369.png" /><figcaption>**FIGURE 6** | Comparison of crystalline parameters of PAN/CNT and PAN fibers at different stretching conditions: (A) Crystallite size, (B) Crystallinity.</figcaption></figure>

influence of steam drawing on crystalline evolution of PAN/CNT fibers compared to PAN fibers.

## Synchrotron X-Ray Diffraction

Synchrotron radiation 2D X-ray diffraction patterns of PAN/CNT fibers and their 1D equatorial and meridional cross section profiles are shown in **Figure 7**. Freely spun fibers exhibited broad reflections with similar equatorial and meridional profiles, resulting from low crystallinity and irregular placement of crystals inside fibers. Application of jet stretch sharpens the peak slightly with small differences in diffraction area in equatorial and meridional directions. After steam drawing, the reflections become sharper and their intensity on the equator increases due to increased crystal perfection. Since single fibers were used for synchrotron X-ray diffraction, the intensity of diffraction ring decreases with increasing the stretching ratio and diameter reduction.

The difference between intensity (I) and area (A) under the peak in equatorial (eq.) and meridional (mer.) profiles can be used as a measure of crystal orientation along the fiber direction (Table 4). Ieq/Imer, and Aeq/Amer. ratios of jet-stretched PAN/CNT fibers have been increased slightly compared to freely

spun fibers. Application of steam drawing with the ratio of 5 to the jet-stretched fibers increases these ratios significantly, revealing the remarkable improvement in the crystal orientation.

When PAN/CNT fibers are stretched, CNTs can also orient along the fiber direction (Dintcheva et al., 2011; Mai et al., 2011). Since crystals can nucleate in the vicinity of CNTs, increasing the orientation of CNTs improves the crystal

<figure><img src="image_370.png" /><figcaption>**FIGURE 7** | Synchrotron radiation 2D X-ray diffraction patterns of PAN/CNT fibers at different stretching conditions and their 1D equatorial and meridional cross section profiles. (A) PC, (B) PC-S2, (C) PC-S2.5, (D) PC-D5, (E) PC-S2-D5.</figcaption></figure>
<table><caption>**TABLE 4** | Parameters extracted from equatorial ad meridional scans of synchrotron radiation X-ray diffraction patterns of PAN/CNT and PAN fibers spun at different stretching conditions.</caption><thead><tr><th>Sample name</th><th>Spinning condition</th><th><i>l</i><sub>eq./lmer.</sub></th><th><i>A</i><sub>eq./Amer.</sub></th></tr></thead><tbody><tr><td>PC</td><td>Freely spun</td><td>1.4</td><td>1.49</td></tr><tr><td>PC-S2</td><td>Jet stretched</td><td>1.51</td><td>1.58</td></tr><tr><td>PC-S2.5</td><td>Jet stretched</td><td>1.78</td><td>1.81</td></tr><tr><td>PC-D5</td><td>Steam drawn</td><td>4.44</td><td>4.76</td></tr><tr><td>PC-S2-D5</td><td>Jet stretched and steam drawn</td><td>18.44</td><td>15.65</td></tr><tr><td>P</td><td>Freely spun</td><td>0.81</td><td>0.53</td></tr><tr><td>P-D5</td><td>Steam drawn</td><td>4.05</td><td>4.09</td></tr><tr><td>P-S2-D5</td><td>Jet stretched and steam drawn</td><td>7.91</td><td>10.38</td></tr></tbody></table>

I, diffraction intensity; A, area under the peak; eq, equatorial; mer, meridional.

orientation at polymer/CNT interface, which leads to higher crystal orientation compared to the structures without CNTs (Table 4). It confirms the possibility of crystal growth in the vicinity of CNTs.

## Mechanical Properties

### Young's Modulus

The Young's modulus of PAN and PAN/CNT fibers increases with jet stretch (Table 5). Even with the lower crystallinity of jet-stretched PAN/CNT fibers in comparison with the corresponding PAN reference fibers (Table 3), Young's modulus of PC-S2 is still 9% higher than that of P-S2, and Young's modulus of PC-S2.5 and P-S2.5 are almost identical. It seems that better orientation of the polymer chains and CNTs by application of jet stretch can result in the growth of crystals on CNTs and in increasing the crystal orientation along the fiber axis. Hence, the jet-stretched PAN/CNT fibers have higher Young's modulus despite high porosity and low crystallinity compared to the corresponding PAN fibers.

When steam drawing is applied to coagulated fibers, Young's modulus has increased significantly due to improved polymer and crystal orientation as well as larger crystalline domain size

and crystallinity even at high porosity (Table 5 and Figure 8A). Comparison of Young's modulus of the steam-drawn fibers at the ratio of 2.5 with the jet-stretched fibers at the ratio of 2.5 shows the larger impact of steam drawing on Young's modulus. It should be noted that steam drawing has larger impact on Young's modulus of PAN/CNT fibers than that of PAN reference fibers. For example, Young's modulus of PC-D2.5 fibers is 100% higher than that of PC-S2.5, whereas Young's modulus of P-D2.5 is 67% higher than that of P-S2.5. In addition, Young's modulus of PC-D2.5 is 20% higher than that of P-D2.5. It is evident that the presence of multiwalled CNTs inside the fiber structure has improved Young's modulus. The improvement in orientation of polymer chains and crystals in the fiber direction is the reason for the improvement of Young's modulus in PAN/CNT fibers.

## Tenacity

The mechanical properties of fibers increase with application of elongational forces inside the coagulation bath through alignment of bonds of the chain backbone in the fiber direction and crystallinity enhancement. Jet stretch has increased the tenacity of PAN/CNT and PAN reference fibers up to 21 and 24%, respectively (Table 5).

Steam drawing increases the tenacity of PAN/CNT and PAN fibers significantly (Table 5 and Figure 8B). For example, application of steam drawing at the ratio of 5 results in 320-350% increase in tenacity of PAN/CNT and PAN fibers. Steam drawing is more effective in improving the structure and mechanical properties than jet stretching. The tenacity of steam-drawn PAN/CNT composite fibers at the ratio of 2.5 is 65% higher than that of jet-stretched PAN/CNT fibers at the ratio of 2.5. Steam drawing at temperatures above the glass transition temperature of PAN has increased the freedom of movement of polymer chains and CNTs in PAN/CNT fibers. Crystalline domain size, crystallinity, and crystal orientation have also been improved as a result of better orientation of polymer chains in the fiber direction (Tables 3, 4). However, steam drawing has mostly increased the

<table><caption>**TABLE 5** | Mechanical properties of PAN/CNT and PAN reference fibers at different stretching conditions.</caption><thead><tr><th>Condition</th><th>Sample name</th><th>Tenacity (cN/tex)</th><th>% CV</th><th>Strain at break (%)</th><th>% CV</th><th>Young's modulus (cN/tex)</th><th>% CV</th></tr></thead><tbody><tr><td colspan="8"><strong>PAN/CNT</strong></td></tr><tr><td>Freely spun</td><td>PC</td><td>2.4</td><td>11.6</td><td>40</td><td>16.2</td><td>50.1</td><td>16.4</td></tr><tr><td rowspan="2">Jet stretched</td><td>PC-S2</td><td>2.6</td><td>4.9</td><td>19</td><td>18.1</td><td>73.0</td><td>8.8</td></tr><tr><td>PC-S2.5</td><td>2.9</td><td>5.7</td><td>18</td><td>16.2</td><td>77.6</td><td>17.0</td></tr><tr><td rowspan="3">Steam Drawn</td><td>PC-D2.5</td><td>4.7</td><td>15.5</td><td>15</td><td>21.3</td><td>154.2</td><td>4.6</td></tr><tr><td>PC-D5</td><td>9.9</td><td>6.9</td><td>28</td><td>6.3</td><td>156.2</td><td>16.6</td></tr><tr><td>PC-S2-D5</td><td>6.8</td><td>7.6</td><td>13</td><td>8.9</td><td>166.1</td><td>11.4</td></tr><tr><td colspan="8"><strong>PAN</strong></td></tr><tr><td>Freely spun</td><td>P</td><td>2.4</td><td>8.9</td><td>33</td><td>14.0</td><td>48.1</td><td>11.6</td></tr><tr><td rowspan="2">Jet stretched</td><td>P-S2</td><td>2.7</td><td>7.5</td><td>36</td><td>7.0</td><td>67.2</td><td>7.0</td></tr><tr><td>P-S2.5</td><td>3.2</td><td>2.8</td><td>34</td><td>7.2</td><td>78.1</td><td>8.2</td></tr><tr><td rowspan="3">Steam drawn</td><td>P-D2.5</td><td>6.8</td><td>5.9</td><td>25</td><td>13.2</td><td>130.2</td><td>15.0</td></tr><tr><td>P-D5</td><td>10.8</td><td>7.0</td><td>35</td><td>6.9</td><td>153.8</td><td>16.4</td></tr><tr><td>P-S2-D5</td><td>13.1</td><td>2.3</td><td>24</td><td>5.4</td><td>263.7</td><td>6.1</td></tr></tbody></table>
<figure><img src="image_371.png" /><figcaption>**FIGURE 8** | Mechanical properties of PAN/CNT and PAN fibers at different stretching conditions, (A) Young's modulus, (B) Tenacity.</figcaption></figure>

chain orientation of PAN reference fibers without increasing its crystallinity.

The highest tenacity of PAN/CNT fibers was achieved at steam drawing ratio of 5 (9.928 cN/tex), whereas PAN fibers had the highest tenacity when they were jet stretched at the ratio of 2 and then steam drawn at the ratio of 5 (13.112 cN/tex). Carbon nanotubes alter the optimum stretching conditions by affecting the orientation and crystallinity of structure. At similar steam drawing conditions, the tenacity of PAN fibers is higher than that of PAN/CNT fibers. Supposedly, CNTs have reduced the structure integrity, increasing the porosity. Not only does the polymer orientation affect the fiber strength, but also defects and voids in the molecular structure have an important role in mechanical properties of fibers (Ouyang et al., 2015). Macrovoids act as breaking points during stretching. To increase the tenacity of PAN/CNT fibers, it is necessary to control the coagulation and drying stage during the wet-spinning process to lower the porosity of the structure. Comparing the influence of various types of CNTs on the properties of PAN/CNT fibers, Chae et al. (2005) reported that multiwalled CNTs have a larger effect on tenacity and strain at break. However, the results of our work show that Young's modulus

of fibers is more affected by the presence of multiwalled CNTs than tenacity.

When a combination of jet stretch at the ratio of 2 and steam drawing at the ratio of 5 is applied on fibers, the mechanical properties of PAN fibers exceeds that of PAN/CNT fibers. This can possibly be due to the negative effect of the high jet stretch ratio on porosity, morphology, crystalline structure, and mechanical properties of PAN/CNT fibers. Taking into account the outcomes of our previous paper (Mirbaha et al., 2019) as well as the observations in the current work, it is clear that for production of PAN/CNT fibers, lower jet-stretch ratios should be applied on fibers. In addition, controlling the porosity of structure by optimization of spinning conditions (i.e., coagulation, stretching and drying conditions) can lead to improvement in mechanical properties of PAN/CNT fibers compared to PAN fibers.

## CONCLUSION

The present research on the impact of different stretching conditions on the crystalline structure and properties of wet-spun PAN/CNT microfibers led to the following outcomes:

(1) Steam drawing has a larger effect on crystalline structure and Young's modulus of PAN/CNT fibers than jet stretching, whereas jet stretching and steam drawing have a comparable effect on the crystalline structure of PAN fibers. Application of the steam draw ratio of 2.5 has improved Young's modulus of PAN/CNT fibers by a maximum of 100% in comparison to corresponding PAN/CNT fibers jet-stretched at the ratio of 2.5.

(2) The effect of steam drawing on the crystalline structure and Young's modulus of PAN/CNT fibers is larger than its effect on PAN fibers. Young's modulus of PAN/CNT fibers could increase up to 19% higher than that of PAN fibers at a steam drawing ratio of 2.5. The existence of multiwalled CNTs inside the fibers has provided better crystalline structure of PAN/CNT fibers at lower steam drawing ratios compared to PAN fibers via nucleation of crystals around the CNTs. Improvement in the orientation of polymer chains and crystals in the fiber direction is the reason for the enhancement Young's modulus.

The results of this research can help to understand how different stretching stages in wet spinning (i.e., jet stretching inside coagulation bath and steam drawing after fiber coagulation) affect the structure and properties of composite PAN/CNT fibers in comparison to PAN fibers.

## DATA AVAILABILITY STATEMENT

The raw data supporting the conclusions of this article will be made available by the authors, without undue reservation, to any qualified researcher.
## AUTHOR CONTRIBUTIONS

HM contributed to the conceptualization, doing the experiments, original draft preparation, writing, and editing the manuscript. PS contributed to editing the manuscript. MD'I contributed to the data acquiring. PS and MD'I contributed to the methodology. PS, MD'I, and HM contributed to the data analysis. PS, NP, and SA contributed to the providing resources. PN and NP contributed to the supervision and project administration. NP contributed to the original draft preparation and editing. All authors contributed to the article and approved the submitted version.

## REFERENCES

Arbab, S., Noorpanah, P., Mohammadi, N., and Soleimani, M. (2008). Designing index of void structure and tensile properties in wet-spun polyacrylonitrile (PAN) Fiber. I. Effect of dope polymer or nonsolvent concentration. *J. Appl. Polym. Sci.* 109, 3461–3469. doi: 10.1002/app.28458

Arbab, S., Noorpanah, P., Mohammadi, N., and Zeinolebadi, A. (2011a). Exploring the effects of non-solvent concentration, jet-stretching and hot-drawing on microstructure formation of poly(acrylonitrile) fibers during wet-spinning. *J. Polym. Res.* 18, 1343–1351. doi: 10.1007/s10965-010-9537-7

Arbab, S., Noorpanah, P., Mohammadi, N., and Zeinolebadi, A. (2011b). Simultaneous effects of polymer concentration, jet-stretching, and hot-drawing on microstructural development of wet-spun poly(acrylonitrile) fibers. *Polym. Bull.* 66, 1267–1280. doi: 10.1007/s00289-010-0422-7

Bakhtiar Davijani, A. A., and Kumar, S. (2015). Ordered wrapping of poly(methyl methacrylate) on single wall carbon nanotubes. *Polymer* 70, 278–281. doi: 10.1016/j.polymer.2015.06.018

Basu-Dutt, S., Minus, M. L., Jain, R., Nepal, D., and Kumar, S. (2012). Chemistry of carbon nanotubes for everyone. *J. Chem. Educ.* 89, 221–229. doi: 10.1021/ed1005163

Brown, T. R., Choudhury, R. P., Meree, C. E., Beckham, H. W., Kumar, S., and Shofner, M. L. (2014). Viscoelastic properties and structure of poly(acrylonitrile-co-methacrylic acid) polymer solutions for gel spinning at long aging times. *J. Appl. Polym. Sci.* 131:39821.

Cao, C., Chung, T., Chen, S., Dong, Z., and Jiang, S. (2004). The study of elongation and shear rates in spinning process and its effect on gas separation performance of poly(ether sulfone) (PES) hollow fiber membrane. *Chem. Eng. Sci.* 59:1053. doi: 10.1016/j.ces.2003.10.023

Chae, H. G., Sreekumar, T. V., Uchida, T., and Kumar, S. (2005). A comparison of reinforcement efficiency of various types of carbon nanotubes in polyacrylonitrile fiber. *Polymer* 46, 10925–10935. doi: 10.1016/j.polymer.2005.08.092

Chen, J., Ge, H., Liu, H., Li, G., and Wang, C. (2010). The coagulation process of nascent fibers in PAN Wet spinning. *J. Wuhan Uni. Technol.* 25, 200–205. doi: 10.1007/s11595-010-2200-7

Dintcheva, N. T., Arrigo, R., Nasillo, G., Caponetti, E., and Mantia, F. P. L. (2011). On the role of extensional flow in morphology and property modifications of MWCNT/polyamide-based fibers. *Macromol. Mater. Eng.* 296, 645–657. doi: 10.1002/mame.201000355

Gissinger, J. R., Pramanik, C., Newcomb, B., Kumar, S., and Heinz, H. (2018). Nanoscale structure-property relationships of polyacrylonitrile/CNT composites as a function of polymer crystallinity and CNT diameter. *ACS Appl. Mater. Interfaces* 10, 1017–1027. doi: 10.1021/acsami.7b09739

Gong, Y., Du, R., Mo, G., Xing, X., Lu, C. X., and Wu, Z. (2014). In-situ Microstructural changes of polyacrylonitrile based fibers with stretching deformation. *Polymer* 55, 4270–4280. doi: 10.1016/j.polymer.2014.06.073

Ji, B., Wang, C., and Wang, Y. (2007). Effect of jet stretch on polyacrylonitrile as-spun fiber formation. *J. Appl. Polym. Sci.* 103, 3348–3352. doi: 10.1002/app.25490

## FUNDING

NP was supported by the European Commission under the FET Proactive (“Neurofibres”) Grant No. 732344. NP and HM were also supported by the PRIN-2017 7TTP3S.

## ACKNOWLEDGMENTS

We thank MCX beamline and the technical staff of the Italian synchrotron Elettra for collaboration.

Kang, D. W., and Ryu, S. H. (2013). Orientation of carbon nanotubes in polypropylene melt. *Polym. Int.* 62, 152–157. doi: 10.1002/pi.4270

Li, X., Qin, A., Zhao, X., Liu, D., Wang, H., and He, C. (2015). Drawing dependent structures, mechanical properties and cyclization behaviors of polyacrylonitrile and polyacrylonitrile/carbon nanotube composite fibers prepared by plasticized spinning. *Phys. Chem. Chem. Phys.* 17, 21856–21865. doi: 10.1039/c5cp02498f

Li, Y., Yu, Y., Liu, Y., and Lu, C. (2018). Interphase Development in Polyacrylonitrile/SWNT Nano-composite and its Effect on Cyclization and Carbonization for Tuning Carbon Structures. *ACS Appl. Nano Mater.* 1, 3105–3113. doi: 10.1021/acsnanm.8b00125

Lian, F., Liu, J., Ma, Z., and Liang, J. (2012). Stretching-induced deformation of polyacrylonitrile chains both in quasicrystals and in amorphous regions during the in situ thermal modification of fibers prior to oxidative stabilization. *Carbon* 50, 488–499. doi: 10.1016/j.carbon.2011.09.003

Lin, Y., Che, J., Pu, S., and Wang, K. (2019). Structural origins of mechanical strengthening in poly(phenylene sulfide)multiwalled carbon nanotube nanocomposites obtained via hot-stretching. *Polym. Comp.* 40, E589–E598.

Liu, H., Lv, F., Li, J., Cao, T., Wan, C., Zhang, W., et al. (2015). Two-stage drawing process to prepare high-strength and porous ultrahigh-molecular-weight polyethylene fibers: cold drawing and hot drawing. *J. Appl. Polym. Sci.* 132:42823. doi: 10.1002/app.42823

Mai, F., Pan, D., Gao, X., Yao, M., Deng, H., Wang, K., et al. (2011). Extension-induced mechanical reinforcement in melt-spun fibers of polyamide 66/multiwalled carbon nanotube composites. *Polym. Int.* 60, 1646–1654. doi: 10.1002/pi.3144

Mikolajczyk, T., Szparaga, G., Bogun, M., Fraczek-Szczypta, A., and Blazewicz, S. (2010). Effect of spinning conditions on the mechanical properties of polyacrylonitrile fibers modified with carbon nanotubes. *J. Appl. Polym. Sci.* 115, 3628–3635. doi: 10.1002/app.31414

Minus, M. L., Chae, H. G., and Kumar, S. (2012). Polyethylene crystallization nucleated by carbon nanotubes under shear. *ACS Appl. Mater. Interface* 4, 326–330. doi: 10.1021/am2013757

Mirbaha, H., Noorpanah, P., Scardi, P., D'incau, M., Greco, G., Valentini, L., et al. (2019). The impact of shear and elongational forces on structural formation of polyacrylonitrile/carbon nanotubes composite fibers during wet spinning process. *Materials* 12:2797. doi: 10.3390/ma12172797

Ouyang, Q., Chen, Y., Wang, X., Ma, H., Li, D., and Yang, J. (2015). Supramolecular structure of highly oriented wet-spun polyacrylonitrile fibers used in the preparation of high-performance carbon fibers. *J. Polym. Res.* 22:229.

Ouyang, Q., Chen, Y. S., Zhang, N., Mo, G. M., Li, D. H., and Yan, Q. (2011). Effect of jet swell and jet stretch on the structure of wet-spun polyacrylonitrile fiber. *J. Macromol. Sci. B Phys.* 50, 2417–2427. doi: 10.1080/00222348.2011.564104

Rebuffi, L., Plaisier, J. R., Abdellatief, M., Lausi, A., and Scardi, P. (2014). MCX: a synchrotron radiation beamline form X-ray diffraction line profile analysis. *Z. Anorg. Allg. Chem.* 640, 3100–3106. doi: 10.1002/zaac.201400163

Sahin, K., Fasanella, N. A., Chasiotis, I., Lyons, K. M., Newcomb, B. A., Kamath, M. G., et al. (2014). High strength micron size carbon fibers from polyacrylonitrile-carbon nanotube precursors. *Carbon* 77, 442–453. doi: 10.1016/j.carbon.2014.05.049
Sapalidis, A., Sideratou, Z., Panagiotaki, K. N., Sakellis, E., Kouvelos, E. P., Papageorgiou, S., et al. (2018). Fabrication of antibacterial poly(Vinyl Alcohol) nanocomposite films containing dendritic polymer functionalized multi-walled carbon nanotubes. *Front. Mater.* 5:11. doi: 10.3389/fmats.2018.00011

Ward, I. M. (1997). *Structure and Properties of Oriented Polymers*. Netherlands: Springer.

Wei, H., Suo, X., Lu, C., and Liu, Y. (2020). A comparison of coagulation and gelation on the structures and stabilization behaviors of polyacrylonitrile fibers. *J. Appl. Polym. Sci.* 137:48671. doi: 10.1002/app.48671

Wilms, C., Seide, G., and Gries, T. (2013). The relationship between process technology, structure development and fibre properties in modern carbon fibre production. *Chem. Eng. Trans.* 32, 1609–1614.

Zeng, X., Hu, J., Zhao, J., Zhang, Y., and Pan, D. (2007). Investigating the jet stretch in the wet spinning of PAN fiber. *J. Appl. Polym. Sci.* 106, 2267–2273. doi: 10.1002/app.26929

Zhang, Y., Song, K., Meng, J., and Minus, M. L. (2013). Tailoring polyacrylonitrile interfacial morphological structure by crystallization in the presence of single-wall carbon nanotubes. *ACS Appl. Mater. Interface* 5, 807–814. doi: 10.1021/am302382m

**Conflict of Interest:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2020 Mirbaha, Scardi, D'Incau, Arbab, Nourpanah and Pugno. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
<figure><img src="image_372.png" /></figure>

Check for
updates

# Role of Compatibilizer in Improving the Properties of PLA/PA12 Blends

Amulya Raj<fnref n="1" />,²*, Cedric Samuel<fnref n="1" />,² and Kalappa Prashantha<fnref n="1" />,²,<fnref n="3" />

Engineering properties of partially-biobased poly(L-Lactide)/poly(amide-12) (PLA/PA12) blends are here addressed using a functional compatibilizer, namely poly(L-Lactide) grafted maleic anhydride (PLA-g-MA). PLA/PA12 blends with compatibilizers are processed using twin-screw extrusion, and the amount of PLA-g-PMA is here optimized with respect to final blend properties such as ductility, impact toughness, and thermal resistance. Blend morphologies depict the enhanced PA12 dispersion into PLA-g-MA with strong interfacial adhesion, in accordance with a compatibilization effect of PLA-g-MA in PLA/PA12 blends. Ductility of the compatibilized blends shows remarkable improvements with the highest ductility close to 290% for an optimized PLA/PA12/PLA-g-MA weight composition (respectively, 69/30/1). The impact strength of compatibilized blends showed a similar two-fold improvement for compatibilized blends. Morphology of injection molded samples revealed finer dispersion and stronger interfaces. The morphology of the transversal section displayed fibrillation of PA12 for PLA-g-MA content up to 2 wt.%. Fibrillation and lower droplet size seemed to play a key role in the improvement of ductility in compatibilized blends. Considerable synergy was detected in the dynamic mechanic analysis of the compatibilized blends. The thermal resistance of the PLA/PA12 blends showed improvement, and this behavior could be due to the formation of copolymer leading to additional interfacial crystallization. This study consequently indicates that PLA-g-MA is an efficient compatibilizer for partially biobased PLA/PA12 blends with potential applications in the automotive and electronic industry.

**Keywords:** biobased polymers, poly(lactic acid), polymer blends, compatibilization, morphologies, engineering properties

OPEN ACCESS

**Edited by:**

Jacqueline Anne Johnson,
University of Tennessee Space
Institute (UTSI), United States

**Reviewed by:**

M. E. Ali Mohsin,
University of Technology
Malaysia, Malaysia
Debora Puglia,
University of Perugia, Italy

*Correspondence:

Amulya Raj
amulya.raj@imt-lille-douai.fr

**Specialty section:**

This article was submitted to
Polymeric and Composite Materials,
a section of the journal
Frontiers in Materials

**Received:** 17 March 2020

**Accepted:** 25 May 2020

**Published:** 17 July 2020

**Citation:**

Raj A, Samuel C and Prashantha K (2020) Role of Compatibilizer in Improving the Properties of PLA/PA12 Blends. Front. Mater. 7:193. doi: 10.3389/fmats.2020.00193

## INTRODUCTION

The replacement of petroleum-based polymers to biobased polymers is an immediate need due to the detrimental effects of the former on the environment (Mecking, 2004; Gandini et al., 2011; Mülhaupt, 2013; Harmsen et al., 2014). Among the various biobased polymers, Poly(Lactide) (PLA), also known as Poly(Lactic acid), has gained special attention in the past years owing to its remarkable properties such as renewability (Drumright et al., 2000), biocompatibility (Madhavan Nampoothiri et al., 2010; Farah and Anderson, 2016), transparency (Farah and Anderson, 2016), high modulus (Pang et al., 2010), and biodegradability (Pang et al., 2010; Saeidlou et al., 2012; Hamad et al., 2018; Nofar et al., 2019). These properties led to the commercialization of PLA, thereby making it suitable for various applications such as resorbable sutures

<footnote n="1">¹ Ecole Nationale Supérieure Mines Telecom Lille Douai, Institut Mines Telecom Lille Douai (IMT Lille Douai), Département Technologie des Polymères et Composites & Ingénierie Mécanique (TPCIM), Douai, France, ² Université de Lille, Lille, France,</footnote>

<footnote n="3">³ ACU-Centre for Research and Innovation, Faculty of Natural Science, Adichunchanagiri University, Mandya, India</footnote>
(Pawar et al., 2014), food and beverage packaging (Auras et al., 2004; Ahmed and Varshney, 2011), surgical implants (Pawar et al., 2014), scaffolds (Lasprilla et al., 2012; Lopes et al., 2012), and textiles (Gupta et al., 2007). Despite its propitious properties and applications, the intricacy of implementing PLA in high value or durable applications is restricted due to limitations arising from low melt strength (Nofar et al., 2019), inherent brittleness (Madhavan Nampoothiri et al., 2010; Farah and Anderson, 2016; Hamad et al., 2018), poor impact strength (Farah and Anderson, 2016; Nagarajan et al., 2016), and low heat resistance (Madhavan Nampoothiri et al., 2010; Hamad et al., 2018). Various strategies have been investigated to overcome the drawbacks of PLA, such as toughening mechanisms (Nagarajan et al., 2016; Anderson et al., 2018), copolymerization (Liu and Zhang, 2011; Nagarajan et al., 2016), stereocomplexation of PLA (Li et al., 2016; Tan et al., 2016), addition of nanoparticles and nucleating agents (Shi et al., 2015; Krishnan et al., 2016), and polymer blending (Hamad et al., 2018; Nofar et al., 2019). Among all of the strategies, polymer blending stands out owing to its feasible way to develop new material with versatile and tailor-made properties with the additional benefit of being cost-effective (Utracki and Wilkie, 2014; Mishra et al., 2017; Hamad et al., 2018; Nofar et al., 2019). Blending PLA with engineering polymers is an efficient way to overcome a majority of the limitations of PLA. Over recent years, intensive research is being carried on PLA based blends with engineering polymers such as polycarbonates (Bouzouita et al., 2016), poly(methylmethacrylate) (Samuel and Jean-Marie Raquez, 2013), polyesters (Jiang et al., 2014), and polyamides (Stoclet et al., 2011; Raj et al., 2019). Among the engineering polymers, polyamides stand out, as the blending would enhance the targeted properties of PLA such as low ductility, poor impact strength, and low heat resistance. The improvement in these properties would help in the market penetration of PLA to high-end or durable applications.

Polyamides are engineering polymers with propitious properties such as high ductility, impact strength, stiffness, and higher thermal resistance. In general, PLA/PA blends are immiscible but compatible in nature. Different types of polyamides have been blended with PLA including PA6 (Wang et al., 2010), PA11 (Stoclet et al., 2011), PA6-10 (Pai et al., 2013), PA10-10 (Cailloux et al., 2018), PA10-12 (Raj et al., 2019), and PA12 (Raj et al., 2019). In order to achieve better final properties, it is necessary to improve the compatibility further. One of the classic methods to improve the compatibility is by incorporation of compatibilizers (Utracki and Wilkie, 2014; Zeng et al., 2015). Multifarious compatibilization strategies have been employed in PLA/PA blends by addition of compatibilizers such as polyalkenyl-poly-maleic-anhydride-imide/amide agents, maleic anhydride fraction grafted on polyethylene-octene elastomer, epoxy based compounds, and poly(ethylene-coglycidyl methacrylate-graft-styrene-co-acrylonitrile) (Wang et al., 2010; Kucharczyk et al., 2012; Dong et al., 2014; Walha et al., 2016; Cailloux et al., 2018; Pai et al., 2019; Rasselet et al., 2019). Walha F et al. studied the effect of multifunctional epoxide (Joncryl ADR®-4368) on PLA and PA11 blends, and two routes of compatibilization were investigated: (i) one step

mixing of adding all the components at once to the extruder, (ii) two-step process where initially PLA and Joncryl are premixed and PA11 is introduced later (Walha et al., 2016). Rheological studies revealed good compatibility between the phases and formation of PLA-joncryl-PA11 copolymer was also confirmed. It was observed that both the processes improved the interfacial adhesion, thus reducing droplet size (<1 µm). Significant improvement in ductility was also observed with the highest being for PLA/PA11(80/20) and 0.7 wt.% of Joncryl ADR®-4368 (~355%) (Rasselet et al., 2019). In another interesting study, maleic anhydride grafted on polyethylene octene elastomer was utilized as a compatibilizer in PLA/PA6 (Wang et al., 2010). The blends showed good dispersion with lower domain size, good ductility, and impact strength. This is due to the formation of POE-co-PA6 and POE-co-PLA copolymers, thereby improving the compatibility between the blend components (Wang et al., 2010). There have been some recent studies about reactive compatibilization of PLA/PA11; catalysts such as p-toluenesulfonic acid and titanium isopropoxide have been incorporated to induce ester amide interchange reactions between the blend components (Patel et al., 2014; Gug and Sobkowicz, 2016). Patel et al. observed that the mechanical properties such as ductility, tensile modulus, and toughness of the PLA/PA11 blends were in between the pure polymer's values, suggesting that although the catalyst titanium isopropoxide induced interchange reactions between the polymers, this did not improve the mechanical properties significantly (Patel et al., 2014).

PLA grafted maleic anhydride (PLA-g-MA) has shown to be a potential compatibilizer for PLA blends. It has been incorporated in PLA based starch, poly(ethylene terephthalate glycol) (PETG), and Poly(butylene adipate terephthalate) (PBAT) blends (Zeng et al., 2015). In PLA/starch blends the addition of PLA-g-MA has brought about an improvement in the interface by decreasing the droplet size of the dispersed phase. Significant improvement in tensile strength and ductility were also observed (Huneault and Li, 2007; Wang et al., 2007; Zeng et al., 2015). Additional improvement in morphology, tensile properties, and ductility were found in PLA/PBAT and PLA/PETG blends in the presence of PLA-g-MA as the compatibilizer (Zeng et al., 2015). Maleic anhydride based coupling agents have been previously utilized in PLA/PA blends as mentioned above (Wang et al., 2010; Kucharczyk et al., 2012). However, there are no significant studies of PLA-g-MA as a compatibilizer in the PLA/PA blend system. It would be interesting to observe the effect of PLA-g-MA as a compatibilizer on PLA/PA12 blend properties and investigate if similar improvement in morphology and ductility can be obtained.

In our previous work, we have reported various PLA blends with long-chain aliphatic polyamides. It was found that PA12 showed the best compatibility with PLA with low interfacial tensions, good dispersion, and exceptional mechanical properties (ductility > 170% for 40 wt.% PA12), good impact properties, and thermal resistance (Raj et al., 2019). Hence, in an attempt to obtain further improvement thermal resistance and mechanical properties we present the effect of PLA-g-MA as a compatibilizer on the properties of PLA/PA12 blends.
<figure><img src="image_373.png" /><figcaption>**FIGURE 1** | Stress-strain curves of PLA/PA12 blends with varying concentrations of compatibilizer (PLA-g-MA).</figcaption></figure>

<table><caption>**TABLE 1** | Stress at yield, strain at break, and impact strength for as-produced PLA/PA12/PLA-g-MA blends.</caption><thead><tr><th>PLA/PA12/<br>PLA-g-MA<br>compositions</th><th>Stress at<br>yield<br>(MPa)</th><th>Strain at<br>break<br>(%)</th><th>Young's Modulus<br>modulus<br>(GPa)</th><th>Impact<br>strength<br>(un-notched)<br>(kJ/m²)</th></tr></thead><tbody><tr><td>100/0/0</td><td>66.6 (1.8)</td><td>3.4 (1.5)</td><td>2.9 (0.06)</td><td>18.1 (0.6)</td></tr><tr><td>0/100/0</td><td>39.5 (1.2)</td><td>138.3 (2.3)</td><td>1.4 (0.05)</td><td>No break</td></tr><tr><td>70/30/0</td><td>56.5 (1.6)</td><td>153.0 (3.3)</td><td>2.2 (0.06)</td><td>28.3 (0.7)</td></tr><tr><td>69/30/1</td><td>55.3 (1)</td><td>291.2 (16.6)</td><td>2.1 (0.05)</td><td>45.4 (0.8)</td></tr><tr><td>68/30/2</td><td>53.9 (2.4)</td><td>282.2 (16.2)</td><td>2.1 (0.04)</td><td>46.8 (0.8)</td></tr><tr><td>67/30/3</td><td>52.8 (3.1)</td><td>268.7 (22.9)</td><td>2.1 (0.09)</td><td>40.5 (0.2)</td></tr><tr><td>66/30/4</td><td>52.4 (4.5)</td><td>268.2 (15)</td><td>2.1 (0.1)</td><td>40.4 (0.2)</td></tr><tr><td>65/30/5</td><td>53.7 (0.9)</td><td>260.7 (5.04)</td><td>2.0 (0.3)</td><td>32.1 (2.4)</td></tr></tbody></table>

<figure><img src="image_374.png" /><figcaption>**FIGURE 2** | Ductility of PLA/PA12 blends with varying PLA-g-MA content.</figcaption></figure>

<figure><img src="image_375.png" /><figcaption>**FIGURE 3** | Impact strength of the PLA/PA12 blends with varying concentration of PLA-g-MA.</figcaption></figure>

# MATERIALS AND METHODS

## Materials

Poly(L-lactide) (PLA, grade 4032D, biobased content 100%, density 1.24 g/cm³, melting temperature 171°C, melt volume index 6.5 cm³/10 min at 210°C, according to manufacturer data) was procured from NatureWorks®, USA. Poly(amide-12) (PA12, grade Rilsamid AESNO, density 1.01 g/cm³, melting temperature 179°C, melt volume index 2.1 cm³/10 min at 235°C, according to manufacturer data) was supplied by Arkema, France. PLA and PA12 were used as received in the pellet form. PLA grafted maleic anhydride (PLA-g-MA) was kindly provided to us by the University of Mons and the preparation method can be referred in the research work published by Quintana et al. (2016). It has a maleic anhydride content of 0.45%.

## Processing of PLA/PA12 Blends

Before compounding, the polymers were dried thoroughly under vacuum at 80°C for 8h to remove any traces of moisture in the materials. PLA/PA12/PLA-g-MA (wt.% 70/30/0, 69/30/1, 68/30/2, 67/30/3, 66/30/4, 65/30/5) were processed by twin-screw extrusion using a Haake Rheomex PTW 16 OS twin-screw extruder (Thermoscientific, Germany, screw diameter 16 mm, L/D 40). The average speed maintained during extrusion was 80 rpm, with a total mass flow rate close to 500 g/h and a mean residence time of ~2 min. The barrel temperature regulation is carried out on 10 zones and the temperature profile was maintained constant for all experiments (200-205-210-210-210-210-210-210-200-185°C), from the hopper to the die. The extrudates were cooled in a water bath and pelletized. The
A

<figure><img src="image_376.png" /></figure>

B

<figure><img src="image_377.png" /></figure>

C

<figure><img src="image_378.png" /></figure>

D

<figure><img src="image_379.png" /></figure>

E

<figure><img src="image_380.png" /></figure>

F

<figure><img src="image_381.png" /><figcaption>**FIGURE 4** | Morphology of PLA/PA/PLA-g-MA of various compositions (A) 70/30/0, (B) 69/30/1, (C) 68/30/2, (D) 67/30/3, (E) 66/30/4, and (F) 65/30/5.</figcaption></figure>

pellets were dried at 80°C under vacuum overnight. After thorough drying pellets were used for the following molding processes. Pristine polymers were also extruded through a similar procedure.

Tensile test specimens were prepared by following the ISO-527-1 standards. The dimensions of the specimens are as follows: gauge length 50 mm, width of 5 mm, and thickness of 1.5 mm. They were processed by injection-molding using a Haake Minijet II molding machine. The following processing conditions were used for the preparation of injection-molded specimens: mold temperature 45°C, chamber temperature 215°C, injection pressure 550 bar, cooling time 5 s. Impact specimens were prepared by injection molding (Babyplast 6/10P, Spain). The processing parameters maintained during the preparation of samples were as follows: the temperature ranged from 200 to 210°C and the mold temperature was around 35–50°C. The injection and the consecutive pressure

were 60 and 80 bars, respectively, with a cooling time of 10 s. Specimens for morphological characterization were processed by compression-molding, and Dolouet molding press was utilized for this process. The parameters and conditions were used for all compression-molded specimens (melt temperature: 205°C, melting time and low-pressure cycle: 8 min, high pressure cycle: 5 min).

## Methods

### Scanning Electron Microscopy (SEM)

The morphology of the compression-molded PLA/PA blends was examined using a scanning electron microscope (SEM, Neoscope II, JEOL, magnification ×60,000). Prior to the observation of morphologies, the specimens were cryo-fractured in liquid nitrogen. These fractured samples were then metallized by gold sputtering under vacuum in order to make them conductive. Droplet sizes of the dispersed phase were measured
<table><caption>**TABLE 2** | Number-average and volume-average PA12 droplet diameters (Dn and Dv), PA droplet size polydispersity (Ip) with standard deviation in brackets of PLA/PA12/PLA-g-MA blends (compression molded samples).</caption><thead><tr><th>Blend</th><th>D<sub>n</sub> (μm)</th><th>D<sub>v</sub> (μm)</th><th>I<sub>p</sub></th></tr></thead><tbody><tr><td>70/30/0</td><td>1.1 (0.1)</td><td>1.4 (0.03)</td><td>1.2</td></tr><tr><td>69/30/1</td><td>1 (0.06)</td><td>1.2 (0.05)</td><td>1.2</td></tr><tr><td>68/30/2</td><td>1 (0.04)</td><td>1.2 (0.06)</td><td>1.2</td></tr><tr><td>67/30/3</td><td>1.3 (0.1)</td><td>1.5 (0.09)</td><td>1.1</td></tr><tr><td>66/30/4</td><td>1.3 (0.1)</td><td>1.6 (0.06)</td><td>1.2</td></tr><tr><td>65/30/5</td><td>1.5 (0.07)</td><td>1.8 (0.05)</td><td>1.2</td></tr></tbody></table>

by image analysis using Image J software. The analysis was done on 50 droplets at least followed by size classification into 7/8 classes. The volume and number average droplet radius diameters (Dv, Dn) with the droplet diameter polydispersity (ID) were calculated using Equations 1–3. The standard deviations of the size of the droplets were presented by measuring two samples.

$$D_n = 2 \times \frac{\sum n_i \times R_i}{\sum n_i} \qquad (1)$$

$$D_v = 2 \times \frac{\sum n_i \times R_i^4}{\sum n_i \times R_i^3} \qquad (2)$$

$$I_D = \frac{D_v}{D_n} \qquad (3)$$

with Dn the droplet number-average diameter, Dv the droplet volume-average diameter, Ri the average droplet radius of the i-class, and ni the number of droplets into the i-class.

## Tensile and Impact Tests

Tensile tests were carried out on PLA/PA blends and neat polymers in a tensile test machine (Instron 3110, U.K) and a 1 kN force cell to obtain stress-strain curves. Tensile properties were determined in accordance with ISO 527-1 (pre-load of 0.5 N, strain rate 1 mm/min for Young's modulus, and 10 mm/min for general properties). Five dumbbell specimens were tested at the least. Charpy un-notched impact tests were carried out as per ISO 179-1 standard by using a pendulum impact machine (Model 5101, Zwick, Germany) at 25°C and 50% RH and impact energy of 7.1 J. The reported values were evaluated as averages of seven specimens.

## Dynamic Mechanical Analysis (DMA)

The dynamic mechanical analysis of the specimens was carried out on dual cantilever beam in the flexion mode (DMA+150, MetraviB, France). The samples were heated over a temperature range (25–130°C) at a heating rate of 2°C/min and frequency of 10 Hz. Two rectangular specimens were tested at the least. The cold crystallization peak of PLA was avoided by annealing the specimens at 110°C for 6 h. The approximate thermal resistance of blends which could correspond to heat deflection temperature (HDTa and HDTb) at two different storage moduli

<figure><img src="image_382.png" /><figcaption>**FIGURE 5** | Emulsification curve of PLA/PA12 blends with varying PLA-g-MA content.</figcaption></figure>

(400 and 250 MPa), respectively, were evaluated by the DMA graphs. A method was partially implemented in several studies (Takemori, 1979; Spinella et al., 2016). These storage modulus values were first calibrated using known values on polyamides and temperatures where storage modulus drops down to 400 and 250 MPa were named T<sub>400MPa-HDTa</sub> and T<sub>250MPa-HDTb</sub>, respectively (Battegazzore et al., 2016).

## Differential Scanning Calorimetry (DSC)

Differential scanning calorimetry (DSC Mettler Toledo) was utilized to measure the PLA crystallinity of pristine PLA and PLA/PA12/PLA-g-MA blends. The evaluation was performed on injection-molded specimens following an annealing step. The DSC evaluation was carried out under nitrogen atmosphere and a single heating scan at 10°C/min was used for this purpose. PLA melting temperatures and associated enthalpy of fusion (Tm and ΔHm) were obtained and the PLA crystallinity index (Xc-PLA) was assessed according to Equation 4. It should be noted that PLA crystallinities were normalized to the amount of PLA in PLA/PA blends

$$X_{c-PLA} = \frac{\Delta H_{m-PLA} - \Delta H_{cc-PLA}}{\Delta H_{m}^{0}} \times 100 \qquad (4)$$

with Xc is the degree of crystallinity, ΔHm-PLA the melting enthalpy of PLA, ΔHcc-PLA the cold crystallization enthalpy of PLA during heating, and ΔHm° the melting enthalpy of fusion for 100% crystalline PLA (93 J/g) (Jia et al., 2017).

## Extraction/Separation Experiments by Centrifugation

The PLA matrix and the PA12 dispersed phase of PLA/PA12/PLA-g-MA blends were selectively extracted/separated by PLA dissolution followed by centrifugation in centrifugal rotary machine (Hettich D7200, Germany). Predetermined weights (~1 g) of extruded pellets of PLA/PA12 blends were dissolved in dichloromethane (1 mL)
<figure><img src="image_383.png" /><figcaption>**FIGURE 6** | Cross-sectional SEM analysis of injection-molded PLA/PA12/PLA-g-MA blends (A) 70/30/0, (B) 69/30/1, (C) 68/30/2, (D) 67/30/3, (E) 66/30/4, and (F) 65/30/5.</figcaption></figure>

and acetone (9 mL). The solution was then submitted to centrifugation. Centrifugation time was set 1 h at a rotation speed of 4,000 rpm. After every centrifugation cycle, the dissolved phase was carefully transferred in a beaker and a fresh solvent system (dichloromethane/acetone 1:9) was added to the centrifugation tube containing the insoluble PA12 fraction. The dissolution-centrifugation was repeated 5–6 times. The PLA solution and settled PA12 particle were dried for 48 h prior to further characterizations.

## Fourier Transform Infrared Studies

The chemical structure of the neat PLA, PA12, and extracted PA12 phase was evaluated by Fourier transform infrared (FTIR) spectroscopy (Nicolet 380 FTIR, ThermoScientific, France). The analysis was carried out on the extracted PA12 phase in reflection mode. The absorption spectra were logged utlising a microscope in reflection or transmission mode. The range of spectra employed is 4,000–800 cm⁻¹ by the resolution of 16 cm⁻¹ and utilizing 62 scans.

## RESULTS AND DISCUSSIONS

### Mechanical Properties

The major emphasis of this study was to investigate the effect of compatibilizer on engineering properties such as mechanical and the same under dynamic load. The blends were processed by injection molding and subjected to tensile testing; the related stress-strain curves are plotted in **Figure 1**. The values of stress at yield, elongation at break, and Young's modulus evaluated by tensile tests are tabulated in **Table 1**. The poor ductility of PLA (3.4%) drastically improved by the addition of 30 wt.% PA12 (153%), indicating good compatibility as confirmed in our previous studies (Raj et al., 2019). PLA-g-MA was incorporated in the blends to observe the potential evolution of compatibility and its effect on mechanical properties.

There was a significant improvement in the ductility with the addition of PLA-g-MA as a compatibilizer in PLA/PA12 blends as can be seen in **Figure 1**. Even at low concentrations of PLA-g-MA (1 wt.%), the improvement in ductility (≈291%)
<table><caption>**TABLE 3** | Number-average and volume-average PA12 droplet diameters ($D_n$ and $D_v$), PA droplet size polydispersity ($I_p$) with standard deviation in brackets of PLA/PA12/PLA-g-MA blends.</caption><thead><tr><th>Blend</th><th>D<sub>n</sub> (μm)</th><th>D<sub>v</sub> (μm)</th><th>I<sub>p</sub></th></tr></thead><tbody><tr><td>70/30/0</td><td>1.1 (0.09)</td><td>1.3 (0.09)</td><td>1.2</td></tr><tr><td>69/30/1</td><td>0.94 (0.07)</td><td>1.06 (0.04)</td><td>1.1</td></tr><tr><td>68/30/2</td><td>0.83 (0.04)</td><td>0.9 (0.03)</td><td>1.1</td></tr><tr><td>67/30/3</td><td>0.9 (0.05)</td><td>1 (0.03)</td><td>1.1</td></tr><tr><td>66/30/4</td><td>0.9 (0.03)</td><td>1.02 (0.04)</td><td>1.1</td></tr><tr><td>65/30/5</td><td>0.96 (0.06)</td><td>1.04 (0.09)</td><td>1.1</td></tr></tbody></table>

is noteworthy. Hence, it could be said that very little amount of PLA-g-MA is adequate to further improve the compatibility, leading to higher ductility.

However, with further increase in content of PLA-g-MA, a minor decrease in the ductility is observed, the lowest ductility obtained for 5 wt.% of PLA-g-MA (260%). The decrease in ductility is linear with PLA-g-MA content, and it should be noted that the decrease in ductility is not drastic, suggesting that even at higher concentrations there is no interfacial debonding leading to poor ductility.

The values of Young's modulus and stress at yield are in the range of 2.0–2.1 GPa and 54–55 MPa, respectively for all the compositions. These values are intermediate of PLA (2.9 GPa, 67 MPa) and PA12 (1.4 GPa, 39.5 MPa) in accordance with previous studies of PLA and PA (Stoclet et al., 2011; Nuzzo et al., 2014; Rashmi et al., 2015).

The evolution of ductility with varying content of PLA-g-MA in PLA/PA12 blends is plotted in **Figure 2**. The content of PLA-g-MA required to improve the ductility remarkably (291%) is as low as 1 wt.%, and a further increase in the content does not have an incremental effect on ductility.

The impact strength of the injection molded blend samples was evaluated. The Charpy un-notched impact strength of the PLA/PA12/PLA-g-MA blends are tabulated in **Table 1**. The variation of impact strength with different concentrations of PLA-g-MA is plotted in **Figure 3** for a clear depiction of the influence of compatibilizer on PLA/PA blends. The toughening effect of PA12 can be clearly seen in the PLA/PA12 (70/30, 28.3 kJ/m²) blends without compatibilizer. It is noteworthy to observe that with the incorporation of PLA-g-MA there is further increase in impact strength, suggesting the positive effect of compatibility in blends. The maximum impact strength (45 kJ/m²) achieved was for the blends containing 2 wt.% of PLA-g-MA. Further addition of PLA-g-MA in blends showed a slight linear decrease in impact strength and the least value (32.1 kJ/m²) being for 5 wt.% of PLA-g-MA. Ductility and toughness are considered important properties for durable applications. The addition of compatibilizer (PLA-g-MA) in PLA/PA12 blends has a positive effect on ductility and toughness. Hence, it can be said that these blends could be suitable for high value applications with high ductility requirements.

The SEM images of the blends with various amounts of compatibilizer could give us a better understanding of the

<figure><img src="image_384.png" /><figcaption>**FIGURE 7** | Droplet diameter of PLA/PA12 blends with varying concentration of PLA-g-MA.</figcaption></figure>

improvement in properties. The morphologies of compressed and injection molded samples of blends and their correlation with mechanical properties are discussed in the following section.

## Morphology

PLA/PA12/PLA-g-MA blend samples were compressed and cryogenically fractured; the SEM images of the same are represented in **Figure 4**. The continuous phase is represented by PLA, and PA12 droplets are dispersed in the PLA matrix (**Figure 4**). The blends displayed good dispersion and the dispersion seems to improve with the addition of PLA-g-MA. The $D_n$, $D_v$, and $I_p$ are tabulated in **Table 2**, where it can be observed that the polydispersity index remains constant (≈1.2) for all compositions.

The domain size of PA12 decreased with the addition of PLA-g-MA (≈1 µm) until 2 wt.%; further addition increased the droplet size of PA12. The droplet size of PA12 in blends with the highest concentration of PLA-g-MA (5 wt.%) is ≈1.5 µm, and it could be said that the blends reached saturation at 2 wt.% of PLA-g-MA as further increase in the content results in the bigger droplet size of dispersed phase (PA12) (**Figure 5**). The reduction in the droplet size could be related to compatibilization effect of PLA-g-MA on PLA/PA12 blends at the interface, resulting in decrease of coalescence of droplets. However, the interfacial adhesion is strong between the blend components and remains unaltered with the increase in concentration of PLA-g-MA, suggesting no interfacial debonding occurs even at higher concentrations of PLA-g-MA.

## Morphology of Injection Molded Samples-Correlation With Mechanical Properties

Morphology of injection-molded samples was observed to get an in-depth knowledge of the effect of compatibilizer on the
<figure><img src="image_385.png" /><figcaption>**FIGURE 8** | Transversal section SEM analysis of injection-molded PLA/PA12/PLA-g-MA blends (A) 70/30/0, (B) 69/30/1, (C) 68/30/2, (D) 67/30/3, (E) 66/30/4, and (F) 65/30/5.</figcaption></figure>

mechanical properties of the blends. The SEM images of the transversal and cross-sectional part of the injection-molded samples are depicted in **Figures 6, 8**. The blends display a fine dispersion of PA12 droplets in PLA matrix. Dn, Dv, and Ip of the PLA/PA12/PLA-g-MA blends are tabulated in **Table 3**. The droplet diameter of PA12 ranges from micronic to submicronic with lowest being for 2 wt.% of PLA-g-MA (0.8 µm), and the polydispersity index is ≈1.1 for all the blend compositions.

The evolution of droplet diameter as a function of PLA-g-MA content has been plotted in **Figure 7**. It can be clearly seen that the droplet diameter decreases linearly up to 2 wt.% and an increment in the droplet size was noticed with a further increase in PLA-g-MA concentration. The lower droplet diameters, better interfacial adhesion, and stronger interface between the blend components of injection molded samples (**Figure 6**) when compared to compression molded samples (**Figure 4**) could be the reason for higher ductility in blends (Stoclet et al., 2011;

Walha et al., 2016; Raj et al., 2019). It is interesting to observe that for injection molded samples, the droplet size of 5 wt.% PLA-g-MA (0.96 µm) is lower than PLA/PA12 (70/30) (1.1 µm); this fine dispersion of PA12 droplets even at higher concentration of PLA-g-MA could be the reason for the ductility behavior of the blends at 5 wt.% of PLA-g-MA (260%). The fibrillation of PA12 can be observed in the **Figures 8A–D**.

At higher concentrations of PLA-g-MA the fibrillation of PA12 appears to be hindered as can be observed in the highlighted area in **Figures 8E,F**. This could be due to higher concentration of compatibilizer at the interface limiting the formation of elongated PA12 in the injection molded samples (Mantia et al., 2017; Kuzmanović et al., 2018). This behavior of hindered elongation of PA12 could be related to the slightly lower ductility and impact strength values (260%, 32 kJ/m²) obtained for the blends containing a higher concentration of PLA-g-MA (5 wt.%). To further study the interactions between the compatibilizers and the polymers, FTIR was performed.
<figure><img src="image_386.png" /><figcaption>**FIGURE 9** | FTIR studies of PLA/PA12 blends with varying content of compatibilizers.</figcaption></figure>

<figure><img src="image_387.png" /><figcaption>**FIGURE 10** | Storage modulus as a function of the temperature of PLA/PA12 blends with varying content of PLA-g-MA(wt.%).</figcaption></figure>

## FTIR Studies

FTIR was performed to study the possible grafting or copolymer formation between PLA and PA due to the presence of compatibilizer. FTIR was done on the extracted PA12 phase from centrifugation. The relative absorptions at 3,286 and 1,751 cm⁻¹ correspond to characteristic stretching vibration absorptions of NH amide groups and C=O ester groups of PA12 and PLA, respectively, (Figure 9) (Feng and Ye, 2010; Ogunniran et al., 2012). The existence of the absorption peak corresponding to the carbonyl group of PLA in the extracted phase of PA12 could be considered as evidence to grafting/copolymer formation in low amounts.

Additionally, the absorption peak of the carbonyl group appears to have shifted from 1,751 to 1,758 cm⁻¹ for the blend (Figure 9), suggesting hydrogen bonding interaction

<table><caption>**TABLE 4** | Thermal properties of PLA/PA12 blends with varying content of PLA-g-MA.</caption><thead><tr><th rowspan="2">PLA/PA12/<br>PLA-g-MA<br>compositions</th><th rowspan="2">T<sub>400Mpa</sub><br>HDT<sub>A</sub> (°C)</th><th rowspan="2">T<sub>250Mpa</sub><br>HDT<sub>B</sub> (°C)</th><th colspan="2">Melting point (°C)</th><th rowspan="2">T<sub>g</sub> (°C)</th></tr><tr><th>(PLA)</th><th>(PA12)</th></tr></thead><tbody><tr><td>100/0/0</td><td>78 (0.5)</td><td>84 (1)</td><td>168.0</td><td>-</td><td>61.0</td></tr><tr><td>70/30/0</td><td>81 (2.5)</td><td>95 (3.7)</td><td>169.4</td><td>177.8</td><td>64.7</td></tr><tr><td>69/30/1</td><td>80 (0.5)</td><td>92 (0.0)</td><td>168.6</td><td>176.7</td><td>64.7</td></tr><tr><td>68/30/2</td><td>84 (1)</td><td>99 (0.5)</td><td>168.3</td><td>176.4</td><td>64.2</td></tr><tr><td>67/30/3</td><td>80.5 (1)</td><td>86 (1)</td><td>168.1</td><td>176.4</td><td>64.2</td></tr><tr><td>66/30/4</td><td>81 (1.5)</td><td>89 (1.5)</td><td>169.3</td><td>176.7</td><td>64.2</td></tr><tr><td>65/30/5</td><td>79 (1)</td><td>86 (0.8)</td><td>169.8</td><td>177.2</td><td>64.0</td></tr></tbody></table>

between the molecules of PLA and PA12 resulting from the grafting/copolymer formation of PLA and PA12 (Feng and Ye, 2010). Overall the enhancement of mechanical properties could be due to the combined effects of grafting, PA12 droplet size, and PA12 fibrillation. However, PA12 fibrillation and PA12 droplet size could be a predominant factor for the improvement in ductility and toughness. The amount of PLA grafting could represent a key parameter to control in order to avoid losing PA12 fibrillated structures.

## (Thermo)mechanical Properties

Thermomechanical properties including thermal resistance or heat deflection temperature (HDT) are prominent for high value applications. Thermomechanical properties were evaluated using DMA to obtain the evolution of the storage modulus/loss factor with temperature.

The HDT (T<sub>400MPa-HDTa</sub> and T<sub>250MPa-HDTb</sub>) were subsequently extracted from storage modulus with a specific experimental procedure previously reported (Raj et al., 2019). The HDT of PLA/PA12 blends under load was evaluated by dynamic mechanical analysis. The influence of varying concentrations of compatibilizer on the HDT of blends was investigated through these tests. The variation of storage modulus with temperature has been plotted in Figure 10. It is important to note that the blends were annealed at 110°C to eliminate the cold crystallization of PLA. The storage modulus (E') gradually decreases with temperature with a classical drop for temperatures higher than 70°C associated with the glass transition/α-relaxation of PLA (Figure 10). The α-relaxation temperature of PLA/PA12(70/30) blends is noticed close to 70°C. Higher storage modulus is observed for 2 wt.% of PLA-g-MA blends in comparison with the other compositions. This behavior is still indeterminate but an additional PLA crystallization at PLA/PA12 interface could induce a significant increase in PLA crystallinity (Zhang and Thomas, 2011; Srithep et al., 2013; Raj et al., 2019). The impact of the concentration of compatibilizer on storage modulus is not significant in the glassy state. However, in rubbery state (T > 70-80°C), a prominent influence of 2 wt.% of PLA-g-MA is observed. The other concentrations did not show this synergism.
<figure><img src="image_388.png" /><figcaption>FIGURE 11 | DSC cooling (left) and heating (right) thermogram of PLA/PA12 blends with PLA-g-MA.</figcaption></figure>

Such thermomechanical modifications in the glassy/rubbery state can be correlated with HDT/thermal resistance improvements. Heat deflection temperature (T<sub>400MPa-HDTa</sub> and T<sub>250MPa-HDTb</sub>) of blends were subsequently evaluated from DMA graphs and tabulated in **Table 4**. PLA/PA12/PLA-g-MA blends exhibited similar HDT values as uncompatibilized blends with an exception of PLA-g-MA 2wt.%. Significant improvement of HDT was observed for 2wt.% of PLA-g-MA (84°C-T<sub>400MPa-HDTA</sub> and 99°C-T<sub>250MPa-HDTB</sub>). The HDT values do not display similar trends that were observed in mechanical properties. However, above 2 wt.% of PLA-g-MA, decrease in HDT values was observed with the least being for 5 wt.% (79°C-T<sub>400MPa-HDTA</sub>, 86°C-T<sub>250MPa-HDTB</sub>).

In conclusion, the addition of compatibilizer to the PLA/PA12 blends is beneficial to improve thermal resistance with a considerable increase in the thermal resistance for 2 wt.% of PLA-g-MA. Similar to tensile and impact properties, the optimum content of PLA-g-MA to enhance thermal resistance of PLA/PA12 blends is found to be at 2 wt.%.

Differential scanning calorimetry was performed on PLA/PA12 blends to investigate several thermal properties such as glass transition temperature, melting temperature, and cold crystallization behavior. The T<sub>g</sub> of all the blends is around 64°C, which is an intermediate value between PLA and PA12. This is a typical additive rule behavior and it could be said that the compatibilizer does not affect T<sub>g</sub> of the blend. It can be observed in heating thermograms (Figure 11) that the melting point of the blends with compatibilizer is slightly lower than that of uncompatibilized blends, the reaction of anhydride group of PLA-g-MA and amino group PA12 could be the reason, thereby suggesting a formation of copolymer at the interface. This copolymer could be a hindrance to the motion of segments causing difficulty in polymer chain arrangement (Wang et al., 2010). However, there is a minor increase in the melting point at 4 and 5 wt.% of PLA-g-MA. The cooling thermograms of the blends show a typical droplet and bulk crystallization behavior

of PLA/PA blends (Stoclet et al., 2011; Rashmi et al., 2015; Raj et al., 2019). The temperature at which droplet crystallization occurs remains almost unaltered (97°C) for uncompatibilized and compatibilized blends. However, the bulk crystallization behavior shows an effect of compatibilizer (PLA-g-MA), with an increase in temperature of bulk crystallization as the content of PLA-g-MA increases. This might be due to larger droplets or the copolymer at interface delaying the bulk crystallization of the blends.

CONCLUSIONS

The major emphasis of this study is to examine the compatibilization effects of PLA-g-MA on PLA/PA12(70/30) blends in a twin-screw extruder. Various concentrations of PLA-g-MA were incorporated in the blends to assess the optimum content of compatibilizer required to achieve properties. The positive effect of PLA-g-MA is evident in morphology, mechanical properties, and thermal resistance. Concerning the mechanical properties, remarkable improvement in ductility was observed in the blends with varying content of PLA-g-MA. The impact properties of the compatibilized blends showed similar trends, and the toughening behavior of PA12 is observed. Based on this study the content of PLA-g-MA required to improve the mechanical properties (ductility and impact strength) is 1–2 wt.%. Morphology of the compression molded blend samples with different concentrations of PLA-g-MA revealed fine dispersion of PA12 in the PLA matrix with a strong interface and good interfacial adhesion. To get a better in-depth understanding of the improvement in mechanical properties, the morphology of injection molded specimens was examined. The cross-sectional SEM images revealed finer dispersion, stronger interface, and better interfacial adhesion when compared to compression molded samples. This could be the reason for the improvement in ductility and toughness of the blends.
Transversal sections of injection molded samples were also examined to assess the relation between fibrillation of PA12 and mechanical properties of the blends; the fibrillation of PA12 is evident in the SEM images. However, for higher content of PLA-g-MA hindrance in the fibrillation might be due to the higher concentration of PLA-g-MA at the interface, impeding the elongation of PA12. FTIR studies were also done to study the probable grafting and its correlation with mechanical properties. The absorption peak corresponding to carbonyl group of PLA in the extracted phase of PA12 could be considered an indication of grafting/copolymer formation. This could be considered as a contributing factor for the improvement in ductility. The influence of compatibilizer on PLA/PA12 blends on the thermal resistance of the blends (HDT) was not as significant as mechanical properties. In conclusion, PLA-g-MA is a suitable compatibilizer for PLA/PA12 blends. Remarkable improvement in ductility was observed and considerable improvement in toughness and heat resistance was attained. These blends could be well suited for durable applications which require high ductility and toughness. Further studies related to particular applications are required to discover the apt application for these blends.

## REFERENCES

Ahmed, J., and Varshney, S. K. (2011). Polylactides—chemistry, properties and green packaging technology: a review. *Int. J. Food Prop.* 14, 37–58. doi: 10.1080/10942910903125284

Anderson, K. S., Schreck, K. M., Hillmyer, M. A., and Schreck, K. M. (2018). Toughening polylactide. *Polym. Rev.* 48, 85–108. doi: 10.1080/15583720701834216

Auras, R., Harte, B., and Selke, S. (2004). An overview of polylactides as packaging materials. *Macromol. Biosci.* 4, 835–864. doi: 10.1002/mabi.200400043

Battegazzore, D., Salvetti, O., Frache, A., Peduto, N., De Sio, A., and Marino, F. (2016). Thermo-mechanical properties enhancement of bio-polyamides (PA10.10 and PA6.10) by using rice husk ash and nanoclay. *Compos. A Appl. Sci. Manuf.* 81, 193–201. doi: 10.1016/j.compositesa.2015.11.022

Bouzouita, A., Samuel, C., Notta-Cuvier, D., Odent, J., Lauro, F., Dubois, P., et al. (2016). Design of highly tough poly (l-lactide)-based ternary blends for automotive applications. *J. Appl. Polym. Sci.* 133:n-a-n/a. doi: 10.1002/app.43402

Cailloux, J., Abt, T., Santana, O., and Carrasco, F. (2018). Effect of the viscosity ratio on the PLA / PA10. 10 bioblends morphology and mechanical properties. *exPRESS Polym. Lett.* 12, 569–582. doi: 10.3144/expresspolymlett. 2018.47

Dong, W., Cao, X., and Li, Y. (2014). High-performance biosourced poly(lactic acid)/polyamide 11 blends with controlled salami structure. *Polym. Int.* 63, 1094–1100. doi: 10.1002/pi.4618

Drumright, R. E., Gruber, P. R., and Henton, D. E. (2000). Polylactic acid technology. *Adv. Mater.* 12, 1841–1846. doi: 10.1002/1521- 4095(200012)12:23<1841::AID-ADMA1841>3.0.CO;2-E

Farah, S., and Anderson. D. G. (2016). Physical and mechanical properties of PLA, and their functions in widespread applications — a comprehensive review. *Adv. Drug Deliv. Rev.* 107, 367–392. doi: 10.1016/jaddr.2016.06.012

Feng, F., and Ye, L. (2010). Structure and property of polylactide/polyamide blends. *J. Macromol. Sci. B* 49, 1117–1127. doi: 10.1080/0022234100 3609179

Gandini, A., Rodríguez, H., Rahman, M., Rogers, R. D., Labidi, J., Hedenqvist, M. S., et al. (2011). The irruption of polymers from renewable resources on the scene of macromolecular science and technology. *Green Chem.* 13, 1061–1083. doi: 10.1039/c0gc00789g

## DATA AVAILABILITY STATEMENT

All datasets generated for this study are included in the article/supplementary material.

## AUTHOR CONTRIBUTIONS

AR, CS, and KP planned the conceptualization and design of the study. AR carried out the experiments and characterizations of the materials and wrote the first draft of the manuscript. All authors contributed to manuscript revision and read and approved the submitted version.

## ACKNOWLEDGMENTS

We are gratefully acknowledge the International Campus on Safety and Intermodality in Transportation (CISIT, France), the European Community (FEDER funds), as well as the Hauts-de-France Region (France) for their contributions to funding extrusion equipment and characterization tools (dynamic rheometers, microscopes, goniometers, calorimeters, and tensile benches).

Gug, J., and Sobkowicz, M. J. (2016). Improvement of the mechanical behavior of bioplastic poly(lactic acid)/polyamide blends by reactive compatibilization. *J. Appl. Polym. Sci.* 133:n-a-n/a. doi: 10.1002/app.43350

Gupta, B., Revagade, N., and Hilborn, J. (2007). Poly(lactic acid) fiber: an overview. *Prog. Polym. Sci.* 32, 455–482. doi: 10.1016/j.progpolymsci.2007.01.005

Hamad, K., Kaseem, M., Ayyoob, M., Joo, J., and Deri, F. (2018). Polylactic acid blends: the future of green, light and tough. *Prog. Polym. Sci.* 85, 83–127. doi: 10.1016/j.progpolymsci.2018.07.001

Harmsen, P. F. H., Hackmann, M. M., and Bos, H. L. (2014). Green building blocks for bio-based plastics. *Biofuel. Bioprod Bior.* 8, 306–324. doi: 10.1002/ bbb.1468

Huneault, M. A., and Li, H. (2007). Morphology and properties of compatibilized polylactide/thermoplastic starch blends. *Polymer (Guildf)*. 48, 270–280. doi: 10.1016/j.polymer.2006.11.023

Jia, S., Yu, D., Zhu, Y., Wang, Z., Chen, L., Fu, L., et al. (2017). Morphology, crystallization and thermal behaviors of PLA-based composites: wonderful effects of hybrid GO/PEG via dynamic impregnating. *Polymers (Basel)*. 9:528. doi: 10.3390/polym9100528

Jiang, W.-R., Bao, R.-Y., Yang, W., Liu, Z.-Y., Xie, B.-H., and Yang, M.-B. (2014). Morphology, interfacial and mechanical properties of polylactide/poly(ethylene terephthalate glycol) blends compatibilized by polylactide-g-maleic anhydride. *Mater. Des.* 59, 524–531. doi: 10.1016/j.matdes.2014.03.016

Krishnan, S., Pandey, P., Mohanty, S., and Nayak, S. K. (2016). Toughening of polylactic acid: an overview of research progress. *Polym. Plast. Technol. Eng.* 55, 1623–1652. doi: 10.1080/03602559.2015.1098698

Kucharczyk, P., Sedlarik, V., Miskolczi, N., Szakacs, H., and Kitano, T. (2012). Properties enhancement of partially biodegradable polyamide/polylactide blends through compatibilization with novel polyalkenyl-poly-maleic-anhydride-amide/imide-based additives. *J. Reinf. Plast. Compos.* 31, 189–202. doi: 10.1177/0731684411434150

Kuzmanović, M., Delva, L., Mi, D., Martins, C., Cardon, L., and Ragaert, K. (2018). Development of crystalline morphology and its relationship with mechanical properties of PP/PET microfibrillar composites containing POE and POE-g-MA. *Polymers (Basel)*. 10:291. doi: 10.3390/polym10030291

Lasprilla, A. J. R., Martinez, G. A. R., Lunelli, B. H., Jardini, A. L., and Filho, R. M. (2012). Poly-lactic acid synthesis for application in biomedical devices — a review. *Biotechnol. Adv.* 30, 321–8. doi: 10.1016/j.biotechadv.2011. 06.019
Li, Z., Tan, B. H., Lin, T., and He, C. (2016). Recent advances in stereocomplexation of enantiomeric PLA-based copolymers and applications. *Prog. Polym. Sci.* 62, 22–72. doi: 10.1016/j.progpolymsci.2016.05.003

Liu, H., and Zhang, J. (2011). Research progress in toughening modification of poly(lactic acid). *J. Polym. Sci. B Polym. Phys.* 49, 1051–1083. doi: 10.1002/polb.22283

Lopes, M. S., Jardini, A. L., and Filho, R. M. (2012). Poly (lactic acid) production for tissue engineering applications. *Procedia Eng.* 42, 1402–1413. doi: 10.1016/j.proeng.2012.07.534

Madhavan Nampoothiri, K., Nair, N. R., and John, R. P. (2010). An overview of the recent developments in polylactide (PLA) research. *Bioresour. Technol. 101*, 8493–8501. doi: 10.1016/j.biortech.2010.05.092

Mantia, F., Ceraulo, M., Giacchi, G., Mistretta, M., and Botta, L. (2017). Effect of a compatibilizer on the morphology and properties of polypropylene/polyethylentherephthalate spun fibers. *Polymers (Basel)*. 9:47. doi: 10.3390/polym9020047

Mecking, S. (2004). Nature or petrochemistry?—biologically degradable materials. *Angew. Chemie Int. Ed.* 43, 1078–1085. doi: 10.1002/anie.200301655

Mishra, J., Tiwari, S. K., Abolhasani, M. M., Azimi, S., and Nayak, G. C. (2017). “Fundamental of polymer blends and its thermodynamics,” in *Micro and Nano Fibrillar Composites (MFCs and NFCs) From Polymer Blends*, eds S. Thomas, R. Mishra, and N. Kalarikkal (Dhanbad: Elsevier), 27–55. doi: 10.1016/B978-0-08-101991-7.00002-9

Múlhaupt, R. (2013). Green polymer chemistry and bio-based plastics: dreams and reality. *Macromol. Chem. Phys.* 214, 159–174. doi: 10.1002/macp.201200439

Nagarajan, V., Mohanty, A. K., and Misra, M. (2016). Perspective on polylactic acid (PLA) based sustainable materials for durable applications: focus on toughness and heat resistance. *ACS Sustain. Chem. Eng.* 4, 2899–2916. doi: 10.1021/acssuschemeng.6b00321

Nofar, M., Sacligil, D., Carreau, P. J., Kamal, M. R., and Heuzey, M.-C. (2019). Poly (lactic acid) blends: processing, properties and applications. *Int. J. Biol. Macromol.* 125, 307–360. doi: 10.1016/j.ijbiomac.2018.12.002

Nuzzo, A., Coiai, S., Carrocchio, S. C., Dintcheva, N. T., Gambarotti, C., and Filippone, G. (2014). Heat-resistant fully bio-based nanocomposite blends based on poly(lactic acid). *Macromol. Mater. Eng.* 299, 31–40. doi: 10.1002/mame.201300051

Ogunniran, E. S., Sadiku, R., Sinha Ray, S., and Luruli, N. (2012). Morphology and thermal properties of compatibilized PA12/PP blends with boehmite alumina nanofiller inclusions. *Macromol. Mater. Eng.* 297, 627–638. doi: 10.1002/mame.201100254

Pai, F.-C., Chu, H.-H., and Lai, S.-M. (2019). Preparation and properties of thermally conductive PLA/PA 610 biomass composites. *J. Elastomers Plast.* 52:0095244318825286. doi: 10.1177/0095244318825286

Pai, F.-C., Lai, S.-M., and Chu, H.-H. (2013). Characterization and properties of reactive poly(lactic acid)/polyamide 610 biomass blends. *J. Appl. Polym. Sci.* 130, 2563–2571. doi: 10.1002/app.39473

Pang, X., Zhuang, X., Tang, Z., and Chen, X. (2010). Polylactic acid (PLA): research, development and industrialization. *Biotechnol. J.* 5, 1125–1136. doi: 10.1002/biot.201000135

Patel, R., Ruehle, D. A., Dorgan, J. R., Halley, P., and Martin, D. (2014). Biorenewable blends of polyamide-11 and polylactide. *Polym. Eng. Sci.* 54, 1523–1532. doi: 10.1002/pen.23692

Pawar, R. P., Tekale, S. U., Shisodia, S. U., Totre, J. T., and Domb, A. J. (2014). Biomedical applications of poly(lactic acid). *Recent Pat. Regen. Med.* 4, 40–51. doi: 10.2174/2210296504666140402235024

Quintana, R., Persenaire, O., Lemmouchi, Y., Bonnaud, L., and Dubois, P. (2016). Compatibilization of co-plasticized cellulose acetate/water soluble polymers blends by reactive extrusion. *Polym. Degrad. Stab.* 126, 31–38. doi: 10.1016/j.polymdegradstab.2015.12.023

Raj, A., Prashantha, K., and Samuel, C. (2019). Compatibility in biobased poly(L-lactide)/polyamide binary blends: from melt-state interfacial tensions to (thermo)mechanical properties. *J. Appl. Polym. Sci.* 137:48440. doi: 10.1002/app.48440

Rashmi, B. J., Prashantha, K., Lacrampe, M.-F., and Krawczak, P. (2015). Toughening of poly(lactic acid) without sacrificing stiffness and strength by melt-blending with polyamide 11 and selective

localization of halloysite nanotubes. *Express Polym. Lett.* 9, 721–735. doi: 10.3144/expresspolymlett.2015.67

Rasselet, D., Caro-Bretelle, A.-S., Taguet, A., and Lopez-Cuesta, J.-M. (2019). Reactive compatibilization of PLA/PA11 blends and their application in additive manufacturing. *Materials (Basel)*. 12:485. doi: 10.3390/ma12030485

Saeidlou, S., Huneault, M. A., Li, H., and Park, C. B. (2012). Poly(lactic acid) crystallization. *Prog. Polym. Sci.* 37, 1657–1677. doi: 10.1016/j.progpolymsci.2012.07.005

Samuel, C., Jean-Marie Raquez, P. D. (2013). PLLA/PMMA blends: a shear-induced miscibility with tunable morphologies and properties? *Polymer (Guildf)*. 54, 3931–3939. doi: 10.1016/j.polymer.2013.05.021

Shi, X., Zhang, G., Phuong, T., and Lazzeri, A. (2015). Synergistic effects of nucleating agents and plasticizers on the crystallization behavior of poly(lactic acid). *Molecules* 20, 1579–1593. doi: 10.3390/molecules 2011579

Spinella, S., Samuel, C., Raquez, J.-M., McCallum, S. A., Gross, R., and Dubois, P. (2016). Green and efficient synthesis of dispersible cellulose nanocrystals in biobased polyesters for engineering applications. *ACS Sustain. Chem. Eng.* 4, 2517–2527. doi: 10.1021/acssuschemeng.5b01611

Srithep, Y., Nealey, P., and Turng, L.-S. (2013). Effects of annealing time and temperature on the crystallinity and heat resistance behavior of injection-molded poly(lactic acid). *Polym. Eng. Sci.* 53, 580–588. doi: 10.1002/pen.23304

Stoclet, G., Seguela, R., and Lefebvre, J.-M. (2011). Morphology, thermal behavior and mechanical properties of binary blends of compatible biosourced polymers: polylactide/polyamide11. *Polymer (Guildf)*. 52, 1417–1425. doi: 10.1016/j.polymer.2011.02.002

Takemori, M. T. (1979). Towards an understanding of the heat distortion temperature of thermoplastics. *Polym. Eng. Sci.* 19, 1104–1109. doi: 10.1002/pen.760191507

Tan, B. H., Muiruri, J. K., Li, Z., and He, C. (2016). Recent progress in using stereocomplexation for enhancement of thermal and mechanical property of polylactide. *ACS Sustain. Chem. Eng.* 4, 5370–5391. doi: 10.1021/acssuschemeng.6b01713

Utracki, L. A., and Wilkie, C. A. (2014). *Polymer Blends Handbook*. Dordrecht: Springer. doi: 10.1007/978-94-007-6064-6

Walha, F., Lamnawar, K., Maazouz, A., Jaziri, M., Walha, F., Lamnawar, K., et al. (2016). Rheological, morphological and mechanical studies of sustainably sourced polymer blends based on poly(lactic acid) and polyamide 11. *Polymers (Basel)*. 8:61. doi: 10.3390/polym8030061

Wang, N., Yu, J., and Ma, X. (2007). Preparation and characterization of thermoplastic starch/PLA blends by one-step reactive extrusion. *Polym. Int.* 56, 1440–1447. doi: 10.1002/pi.2302

Wang, Y.-L., Hu, X., Li, H., Ji, X., and Li, Z.-M. (2010). Polyamide-6/poly(lactic acid) blends compatibilized by the maleic anhydride grafted polyethylene-octene elastomer. *Polym. Plast. Technol. Eng.* 49, 1241–1246. doi: 10.1080/036025590.2010.496418

Zeng, J.-B., Li, K.-A., and Du, A.-K. (2015). Compatibilization strategies in poly(lactic acid)-based blends. *RSC Adv.* 5, 32546–32565. doi: 10.1039/C5RA01655J

Zhang, M., and Thomas, N. L. (2011). Blending polylactic acid with polyhydroxybutyrate: the effect on thermal, mechanical, and biodegradation properties. *Adv. Polym. Technol.* 30, 67–79. doi: 10.1002/adv.20235

**Conflict of Interest:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2020 Raj, Samuel and Prashantha. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.
<figure><img src="image_389.png" /></figure>

Check for
updates

# *In-situ* Monitoring of the Out-Of-Autoclave Consolidation of Carbon/Poly-Ether-Ketone-Ketone Prepreg Laminate

Florence Saffar¹,²,³*, Camille Sonnenfeld¹, Pierre Beauchêne¹ and Chung Hae Park²,³

¹DMAS, ONERA, Université Paris Saclay, Châtillon, France, ²Technology of Polymers and Composites & Mechanical Engineering Department, IMT Lille Douai, Institut Mines-Télécom, Douai, France, ³Université de Lille, Lille, France

The main objectives of this article were to examine the feasibility of high quality laminate (carbon PEKK—Poly-Ether-Ketone-Ketone) manufacturing under low pressure and to analyze the principal phenomena governing the consolidation quality. The quality of laminate was evaluated in terms of the interlaminar shear strength measured by short-beam shear test and the residual voids observed by micrographic images. This work underlined the possibility to obtain a good interlaminar consolidation quality, i.e., interlaminar shear strength of 100 MPa under a low pressure of 7.0 × 10⁴ Pa even for prepregs which were not designed for out-of-autoclave processes. To better understand the interlaminar consolidation phenomena, we developed an experimental set-up for the *in-situ* monitoring of the laminate consolidation to measure the change of laminate thickness and the temperature gradient in the thickness direction during the manufacturing process. In particular, we identified two major phenomena, the establishment of intimate contact between the adjacent layers at the glass transition temperature and the molten matrix flow at the melting temperature. The assumption on the intimate contact establishment at the glass transition temperature was confirmed by the heat transfer simulation considering the change of thermal contact resistance at the interlaminar interface.

**Keywords:** polymer-matrix-composites (PMCs), interfacial strength, thermal properties, out-of-autoclave (OOA) consolidation, *in-situ* monitoring

## INTRODUCTION

In the aircraft design and manufacturing, thermoplastic composites have many technical advantages over carbon/epoxy composites, such as high impact resistance, a short process cycle time, favorable storage conditions, and ease of patch repair. Among a wide spectrum of thermoplastic matrices for high performance composites are Poly-Aryl-Ether-Ketone (PAEK) matrices which have attracted great attention owing to high environmental resistance at elevated temperature. For example, unidirectional carbon fiber reinforced Poly-Ether-Ether-Ketone (PEEK) prepreg is a representative thermoplastic composite material used in the aircraft design (Bishop, 1985; Carlile et al., 1989; Denault and Dumouchel, 1998; El Kadi and Denault, 2001). Recently, more attention is being paid to Poly-Ether-Ketone-Ketone (PEKK) as a promising matrix material for

## OPEN ACCESS

**Edited by:**
Maria Chiara Bignozzi,
University of Bologna, Italy

**Reviewed by:**
Liqing Wei,
United States Forest Service (USDA),
United States
Veronique Michaud,
École Polytechnique Fédérale de
Lausanne, Switzerland

*Correspondence:
Florence Saffar
florence.saffar@onera.fr

**Specialty section:**
This article was submitted to Polymeric and Composite Materials, a section of the journal Frontiers in Materials

**Received:** 29 March 2020
**Accepted:** 25 May 2020
**Published:** 25 June 2020

**Citation:**
Saffar F, Sonnenfeld C, Beauchêne P and Park CH (2020) *In-situ* Monitoring of the Out-Of-Autoclave Consolidation of Carbon/Poly-Ether-Ketone-Ketone Prepreg Laminate. Front. Mater. 7:195. doi: 10.3389/fmats.2020.00195
carbon fiber reinforced composites (Chang, 1988). PEKK has a higher glass transition temperature (155°C) than PEEK (145°C) and offers a better high temperature resistance, whereas it has a similar melting temperature (335–340°C) which leads to almost the same manufacturing condition as PEEK. In spite of its advantages, the data on the carbon/PEKK consolidation is seldom found in the literature, however, whereas the characterization of carbon/PEEK prepreg systems has been extensively reported (Blundell et al., 1989; Yang and Colton, 1995; Phillips et al., 1997; Ageorges et al., 1998; Gao and Kim, 2000, 2001; Wang et al., 2001; Stokes-Griffin and Compston, 2016a; Talbott et al., 2016). Only a small number of articles about the degradation and isothermal crystallization of PEKK have been published by Choupin et al. and other authors (Hsiao et al., 1991; Krishnaswamy and Kalika, 1996; Choupin, 2017; Tadini et al., 2017; Choupin et al., 2018; Li and Strachan, 2019).

Moreover, a critical drawback in the use of such materials is the high manufacturing cost. In general, the price of those materials is very high. To reduce the labor cost due to the manual lay-up, automated tape placement (ATP) is widely considered (Khan et al., 2010; Comer et al., 2015; Stokes-Griffin and Compston, 2016a,b). Nevertheless, the high consolidation quality cannot be achieved only by ATP and a post-consolidation process is required. For this post-consolidation step, an expensive manufacturing equipment such as autoclave is generally used because high temperature and pressure are needed to consolidate laminates. Indeed, all the commercial grades of carbon/PEKK prepreg available in the current market are designed for the autoclave manufacturing.

Nowadays, there is a growing interest in out-of-autoclave consolidation processes where low compaction pressure is employed to reduce the manufacturing cost (Witik et al., 2012; Centea et al., 2015). For example, laminates can be consolidated under vacuum condition in a heated oven.

About the high pressure consolidation of thermoplastic composites, different phenomena have already been identified such as intimate contact, autohesion, and crystallization.

The phenomenon of intimate contact has been described by different models. In the literature, the principal mechanism of the intimate contact establishment has been assumed to be the flattening of surface roughness which has been modeled by the squeeze flow of the material under compaction pressure (Lee and Springer, 1987; Loos and Dara, 1987; Mantell and Springer, 1992; Butler et al., 1998; Yang and Pitchumani, 2001). In the early work, the main focus was on the description of the geometric pattern on the rough surface of prepreg. For example, Dara and Loos proposed a model where the surface roughness of the prepreg was represented by a series of rectangles with different dimensions (Loos and Dara, 1987). Lee and Springer simplified this model by representing the prepreg surface roughness by a series of rectangles with the same dimensions (Lee and Springer, 1987). Nevertheless, the limits of this model have also been underlined. Butler et al., and Yang and Pitchumani explained that the identification of the geometric parameters in the models was very complex (Butler et al., 1998; Yang and Pitchumani, 2001). The results of profilometric and micrographic studies showed that the prepreg surface roughness could not necessarily

be described by rectangles. Butler et al. highlighted that according to the geometric parameters, the time for the intimate contact establishment could vary between 1 and $10^5$ s (Butler et al., 1998). Hence, Yang and Pitchumani introduced Cantor fractals to represent the surface roughness of prepreg ply (Yang and Pitchumani, 2001). By this method, the estimation of intimate contact establishment time could be close to a real value. Yet, the complexity to determine the geometric parameters still remains as a critical issue to address in this approach.

In another approach, the evolution of the intimate contact during the consolidation was assessed by the change of the thermal contact resistance between the adjacent plies of a laminate. To monitor the evolution of this thermal contact resistance, Schaefer et al. installed thermocouples between the adjacent layers in the middle of composite laminate (Schaefer et al., 2016, 2017). With the temperature values measured by the thermocouples, the thermal diffusivities of initial prepreg and of final consolidated laminate were estimated. Subsequently, they determined the thermal contact resistance at the interface between the adjacent layers and found that it decreased during the consolidation process. The relation between the thermal contact resistance and the intimate contact has also been studied by Levy et al. (2014) and Cassidy and Monaghan (1994). It should be noted that in all the aforementioned works, the composite material was considered to be viscous fluid and the intimate contact was assumed to take place at high temperature (i.e., around the melting point of the matrix) and under high compaction pressure (i.e., over $3.0 \times 10^5$ Pa). Moreover, all the works have focused on the PEEK, Polyimide, and PIMA-M matrix composites, whereas PEKK matrix composites have not been investigated (Lee and Springer, 1987; Butler et al., 1998; Yang and Pitchumani, 2001). Finally, the intimate contact phenomenon of carbon fibers/PEKK laminates began to be investigated in a recent work by Çelik et al. (2020). This last study focuses on the intimate contact phenomenon during laser assisted automated fiber placement manufacturing, and so under high pressure conditions.

Therefore, the main subject of this paper is the low pressure consolidation of carbon fibers/PEKK laminates. In particular, the influence of pressure level on the consolidation is experimentally investigated. Moreover, an in-situ monitoring set-up to study the consolidation phenomena in terms of the temperature differential in the thickness direction and the change of the total thickness of laminates is presented. Finally, thermal modeling of the intimate contact phenomenon under low pressure is proposed.

# EXPERIMENTAL METHODS AND NUMERICAL SIMULATION

## Material

We used unidirectional carbon fiber reinforced PEKK prepreg. The weight fraction of fibers was 66 %, the areal weight was $145 \text{ g/m}^2$ and the thickness of a single ply before the consolidation was 0.15 mm. A microscopic image of the prepreg ply is presented in **Figure 1**. The prepreg was placed in an oven
at 120°C for 80 min before the consolidation process to dry the samples.

## Thermal Expansion Characterization

The thermal expansion of the prepreg in the thickness direction was evaluated by dilatometer (TGA-TDA Setsys Evolution 16/18). The chosen temperature cycle was a 5 K/min heating rate from 40 to 300°C and an inert gas (argon) was used in order to obtain the conditions close to the vacuum manufacturing conditions. Tested samples were laminates with a thickness of 4.3mm consolidated using a press. The prepregs were placed in an oven at 120°C for 80 min before the measurement of thermal expansion coefficient.

## Two Set-Ups for Out-Of-Autoclave Manufacturing

Two different heating methods were employed for out-of-autoclave consolidation, namely, oven and heating plate.

<figure><img src="image_390.png" /><figcaption>**FIGURE 1** | Microscopic image of the prepreg ply.</figcaption></figure>

## Oven Heating

In order to study the feasibility of out-of-autoclave manufacturing, laminate consolidation in a heated oven was used. Using this experimental set-up, only the temperature was measured whereas the thickness of laminate was not monitored. A laminate lay-up of 16 UD plies was chosen. The planar dimensions of the laminates were 115 × 115 mm². Three thermocouples of K type were installed on the planar surfaces of laminate to monitor the temperature evolution as shown in **Figure 2A**. Because the laminate stack was heated in an oven, relatively uniform heating of laminates could be obtained. The temperature cycle of the oven was controlled by Eurotherm system. The vacuum pressure applied to the laminate stack was controlled and recorded by a pressure gauge (Pfeiffer APR250).

## One-Side Heating Plate

The second method which has been developed is a heating plate set-up in order to monitor both the temperature difference and the laminate thickness change. A stack of 16 UD prepreg plies with planar dimensions of 60 × 60 mm² was placed on a stainless steel plate which was heated. The prepreg stack was subsequently covered by a speckled vacuum bag as shown in **Figure 2B**. Two thermocouples of K type with a diameter of 0.25 mm were installed beneath and on the laminate. It should be noted that the laminate was heated only from the bottom side which induced a non-uniform temperature distribution through the thickness direction. As previously done for the oven heating method, the temperature cycle of the stainless steel plate was controlled by Eurotherm system and the vacuum pressure was controlled and recorded by a pressure gauge.

To monitor the change of the laminate thickness during the consolidation, speckles were applied on the upper face of the vacuum bag and two optical cameras (AVT Pike-421, 4/80 lens and 41 mm extension ring) were placed 1,000 mm above the laminate stack. The angle between the two cameras was around 33° to measure the change of the out-of-plane displacement. The data acquisition interval was 30 s. Vis-Snap 8 was used as a data acquisition system and the data processing was performed by Vic-3D. For each time step, the mean value of laminate thickness

A

<figure><img src="image_391.png" /></figure>

B

<figure><img src="image_392.png" /></figure>

C

<figure><img src="image_393.png" /><figcaption>**FIGURE 2** | *In-situ* monitoring system for out-of-autoclave consolidation of thermoplastic composites. (A) Scheme of thermocouples monitoring for oven manufacturing. (B) Schematic of prepreg lay-up. (C) Schematic diagram of *in-situ* monitoring system.</figcaption></figure>
was calculated from around 35 data on the sampling surface of 5 × 5 cm² which was located in the center of the laminate surface of 6 × 6 cm² in order to avoid the edges effects. During the consolidation, the pressure, the temperatures at the top and at the bottom of the laminate and the laminate thickness were recorded. The whole in-situ monitoring system is represented by a schematic drawing in **Figure 2C**.

# Manufacturing Conditions

## Temperature Cycle

Two different temperature cycles were applied in the heating plate method, to investigate the influence of temperature on the interlaminar consolidation.

In the first case, a classic temperature cycle was used for manufacturing with the heating plate. A prepreg stack was heated from room temperature at a heating rate of 5 K/min. Once the temperature reached 360°C (=Melting temperature of PEKK + 25°C), this temperature was maintained for 15 min. Then, the prepreg stack was cooled down at a free cooling rate (see **Figure 3A**).

To better understand the consolidation phenomena, another identical successive consolidation cycle was applied again on the laminate which had already been consolidated.

In the second case, a stack of prepreg which had not been consolidated was submitted to the peak temperature of 230°C which was maintained for 15 min, while employing the same heating and cooling rates (see **Figure 3B**), to observe the interlaminar microstructure of laminate consolidated at the intermediate temperature (i.e., 230°C) between the glass transition temperature ($T_g$) and the melting temperature ($T_m$).

Finally, for the manufacturing in an oven, the same temperature cycle as shown in **Figure 3A** was applied, while employing a different cooling rate of 5 K/min.

## Vacuum Pressure

The compaction pressure applied on the prepreg stack was determined by the difference between the atmospheric pressure and the vacuum pressure. During the manufacturing cycles, a steady vacuum pressure was applied. The used reference value was –995 × 10² Pa. In order to determine the minimum

compaction pressure to obtain a good consolidation quality, different compaction pressure values (i.e., –995 × 10², –973 × 10², –915 × 10², –825 × 10², –740 × 10², –511 × 10², –253 × 10², –209 × 10², –185 × 10², –93 × 10², –24 × 10²) were tested in the oven manufacturing method. Moreover, two different vacuum pressures (–995 × 10² Pa, –325 × 10² Pa) were used for heating plate method to better understand the influence of vacuum pressure on interlaminar consolidation phenomena.

# Assessment of Consolidation Quality: Interlaminar Shear Strength and Void Content

Despite the drawbacks of interlaminar shear strength (ILSS) test related to the multiple failure modes which can occur for thermoplastic composites (Kadlec et al., 2014), ILSS test is widely accepted because the interlaminar shear strength of laminate is a representative parameter which is sensitive to interlaminar properties in shear, and in addition it is easy and fast to perform. Furthermore, their results are reproducible from a test to another and enough discriminating. Three rectangular specimens were cut from each laminate according to the same orientation in the laminate to conduct short-beam shear three-point bending. The dimensions of the specimens were 25 × 12 × 2.35 mm³. For the three-point short-beam shear test, the radius of the loading pin was 5 mm and the radius of the two supporting pins was 2 mm. The distance between the two supporting pins was 12 mm and the displacement speed of the loading pin was 0.5 mm/min. The tests were performed on Zwick/Roell Z010 with a load cell of 10 kN.

Reference laminate samples were fabricated by the compression molding with a high compaction pressure (i.e., 10 × 10⁵ Pa), while applying the same temperature cycle as shown in **Figure 3A**. The corresponding ILSS value was regarded as the reference ILSS of the present prepreg samples.

The void content was evaluated by image analysis. The transverse cross-section of laminates was observed using an optical microscope and the void content was determined by thresholding method on (Image J software, 2020). The observations were made at the zone around the interlaminar interfaces and the observation window size was 1,740 × 60 µm².

<figure><img src="image_394.png" /><figcaption>**FIGURE 3** | Manufacturing cycles. (A) First cycle. (B) Third cycle.</figcaption></figure>
<table><caption>**TABLE 1** | Thermal properties of the prepreg.</caption><thead><tr><th></th><th>T = 50°C</th><th>T = 100°C</th><th>T = 150°C</th><th>T = 200°C</th><th>T = 250°C</th><th>T = 300°C</th></tr></thead><tbody><tr><td>Cp (J/g·°C)</td><td>0.6681</td><td>0.8481</td><td>1.025</td><td>1.269</td><td>1.454</td><td>1.639</td></tr><tr><td>δ (m²/s)</td><td>4.28 × 10<sup>-7</sup></td><td>3.99 × 10<sup>-7</sup></td><td>3.71 × 10<sup>-7</sup></td><td>3.24 × 10<sup>-7</sup></td><td>3.06 × 10<sup>-7</sup></td><td>2.80 × 10<sup>-7</sup></td></tr></tbody></table>

## Characterization of the Thermal Properties and Numerical Simulation of Heat Transfer

One dimensional finite element model was used for the numerical simulation of heat transfer during the heating plate manufacturing method. The computational domain for heat conduction was composed of 16 layers of composite prepreg with a thickness of 0.15 mm and 15 interlaminar interface layers with a thickness of 8 µm which corresponded to the mean height of the prepreg surface roughness. For the heat transfer simulation, the heat transfer module of COMSOL was used. As the boundary conditions, a convective heat transfer coefficient of 15 W·m⁻²·K⁻¹ was applied on the top surface and a preassigned temperature profile was given at the bottom surface. The density of prepreg was 1,400 kg/m³. The thermal diffusivity (δ) has been measured for temperatures between room temperature and 300°C at every 50°C by flash method on the consolidated laminate with a thickness of 2.3 mm. Finally, the specific heat capacity (Cp) of prepreg was evaluated by DSC tests. All the material properties which were used in this model are summarized in Table 1. For the interface layers, the thermal properties were assumed to change, as the intimate contact was established according to the increase of temperature. The thermal conductivity of the interface layers was represented by a sigmoid temperature dependent function. This function was defined by four parameters, viz. the initial conductivity (kᵢₘₐt), the final conductivity (kᵢₚₐt), the inflection temperature (Tᵣ), and the temperature range (dT) (see Figure 4). These four parameters were determined by an inverse identification using Levenberg-Marquardt method (Marquardt, 1963).

## RESULTS AND DISCUSSION

### In-situ Monitoring Data: Temperature Difference and Laminate Thickness

From the in-situ monitoring of thermoplastic laminates manufacturing on the heating plate, the data for the thickness of the laminate and the average temperature gradient through the thickness during the whole consolidation cycle were obtained. Figure 5 shows the time-dependent change of the laminate thickness obtained by the two optical cameras and the change of the average temperature gradient.

The temperature gradient was greatly increased from “Tg-30°C” and reached its maximum value when the temperature of the bottom layer was Tg(see Figure 5). Afterwards, the temperature gradient dropped fast as the temperature increased further from Tg. At the same time, an increase of the laminate thickness was observed as the temperature increased above Tg. Both the laminate thickness and the temperature differential were decreased again around

the melting temperature of the matrix, even if the temperature of the bottom layer continued to increase. We could verify the reproducibility of these observations from the repeated experiments under the same condition. This speculation can be verified by the successive consolidation experiment.

In Figure 6, the evolution of the temperature difference between top and bottom layers is shown during the two identical successive consolidation steps. We can see two sudden decreases of temperature difference at the glass transition temperature and the melting temperature, respectively, during the first consolidation. During the second consolidation of the same laminate, however, such sudden decrease of temperature is not observed at the glass transition temperature. Thus, we can assume that the intimate contact was totally established during the first consolidation and there was no more change of thermal contact resistance at the interlaminar interfaces during the second consolidation.

The change of the laminate thickness can be explained by the competition between the thermal expansion effect and the intimate contact establishment which tends to decrease the thickness due to the flattening of prepreg surface roughness. Until the melting temperature of the matrix, the increase of the laminate thickness according to the increase of temperature can be correlated with the coefficient of thermal expansion (CTE) in the thickness direction of laminate which had been experimentally measured (see Figure 7).

Conversely, it is assumed that the molten matrix flow at the interlaminar interfaces accounts for the decreases of the laminate

<figure><img src="image_395.png" /><figcaption>**FIGURE 4** | Evolution of thermal conductivity of interlaminar interfaces with temperature.</figcaption></figure>
<figure><img src="image_396.png" /><figcaption>**FIGURE 5** | Results of average temperature gradient and thickness against time.</figcaption></figure>

<figure><img src="image_397.png" /><figcaption>**FIGURE 6** | Result of temperature difference for two successive consolidation cycles.</figcaption></figure>

thickness at the melting temperature. Thus, it can be assumed that the intimate contact is not totally complete between the glass transition temperature and the melting temperature, however and there are still some voids at the interlaminar interfaces. The perfect contact between adjacent layers is only obtained around the melting temperature (see **Figure 8**). Microscopic observations of laminates highlight a relatively low void content

which shows the establishment of intimate contact at the interlaminar interfaces (see **Figure 9**). The most of voids are concentrated at the three or four top interlaminar interfaces close to the vacuum bag and the approximate void content at these zones is 5%. At the other zones, it is not easy to identify the interfaces between two adjacent layers and the void content is lower than 2%. This difference between the upper
<figure><img src="image_398.png" /><figcaption>**FIGURE 7** | Comparison between thermal expansion and laminate thickness results.</figcaption></figure>

<figure><img src="image_399.png" /></figure>

<figure><img src="image_400.png" /></figure>

<figure><img src="image_401.png" /></figure>

<figure><img src="image_402.png" /><figcaption>**FIGURE 8** | Schematic of the interlaminar consolidation phenomena. (A) Below T<sub>g</sub>. (B) At T<sub>g</sub> (Establishment of intimate contact). (C) At T<sub>m</sub> (Molten matrix flow). (D) Above T<sub>m</sub> (Interdiffusion of polymer molecules).</figcaption></figure>
and bottom zones in the laminate can be explained by the delay in the temperature rise due to one side heating from the bottom heating plate. The maximum temperature in the upper part near the vacuum bag just reached the melting temperature of the PEKK when that in the bottom part near the heating plate was sufficiently above the melting temperature. Thus, we can suppose that some portion of matrix in the upper layers was not sufficiently melted and therefore the matrix could not flow to fill the interlaminar gaps in these zones. Otherwise, voids could migrate from the bottom zone to the upper zone of the laminate because they might move along the negative temperature gradient according to the phenomenon described by Lu et al. (2004).

## Laminate Consolidated at the Low Temperature

Figure 10 shows a micrographic image of the laminate obtained by low temperature consolidation (i.e., 230°C). Because the laminate was consolidated above the glass transition temperature, the intimate contact began to be established even if it was not totally complete. The surface roughness was reduced and the adjacent layers became in contact with each other. Nevertheless,

there was still a high average content (e.g. ~11%) of interlaminar voids. It was more difficult to assess the gradient of void content in the thickness direction in a partially consolidated laminate (consolidated at the low temperature, i.e., 230°C) than in a fully consolidated laminate (consolidate at the high temperature, i.e., 360°C). It could be noticed only that the void content at the interlaminar interfaces was higher at the outer layers (>10%) than at the middle layers (<6%) of the laminate.

The final laminate quality was also evaluated by ILSS tests. The interlaminar shear strength of the laminate consolidated at 230°C was around 1 MPa. We can conclude that the intimate contact establishment at the glass transition temperature is not enough to build the interlaminar bonding strength. Once the matrix temperature reaches at the melting temperature, the matrix is melted and flows to fill the small gaps between the adjacent layers (see Figure 8C). As the temperature rises further above the melting point, there is an interdiffusion of polymer molecules across the interlaminar interfaces to erase the interface between the adjacent layers which establishes the interlaminar bonding strength (see Figure 8D). Thus, the phenomena at the melting temperature, viz. molten matrix flow and molecule interdiffusion play the principal role in the laminate consolidation process.

<figure><img src="image_403.png" /><figcaption>FIGURE 9 | Micrographic image of laminate consolidated by one side heating plate.</figcaption></figure>

<figure><img src="image_404.png" /><figcaption>FIGURE 10 | Micrographic image of the laminate consolidated at 230°C.</figcaption></figure>
It should be emphasized that the intimate contact begins to be established at a low temperature, viz. around the glass transition temperature, whereas in the literature it has been assumed to occur at a high temperature, i.e., around the melting temperature of the matrix. Hence, the initial surface roughness of prepreg ply is changed above the glass transition temperature, which improves the intimate contact between adjacent layers. This phenomenon can be correlated with the reduction of thermal contact resistance at the interlaminar interfaces. We can see in Figure 5 that the decrease of the average temperature gradient took place around the glass transition temperature. As the intimate contact is established, the heat conduction path is enlarged and the thermal resistance is decreased at the interlaminar interfaces.

## Influence of Compaction Pressure

The influence of the compaction pressure on the interlaminar consolidation of laminate was investigated for both heating methods, i.e., heating plate and oven.

For the heating plate method, two levels of compaction pressure were tested, viz. 995 × 10² Pa and 325 × 10² Pa. Figure 11 presents the in-situ monitoring results of the evolution of the average temperature gradient, which is the temperature difference divided by the laminate thickness (Figure 11A) and the evolution of the laminate thickness (Figure 11B). It should be noticed that the laminate thickness was not uniform on the surface of the laminate. The final thickness of laminate was greater in the zones far from the pressure pump than near the pressure pump. This difference is assumed to be come from the non-uniform distribution of compaction pressure. Moreover, the standard deviation of the laminate thickness was increased with the decreasing of compaction pressure applied during the consolidation cycle. Thus, the maximum value for standard deviation was obtained for the compaction pressure of 325 × 10² Pa which is presented in Figure 11C.

The average temperature gradient and the laminate thickness were highly dependent on the compaction pressure. In the case of low compaction pressure (i.e., 325 × 10² Pa), there was no remarkable change of the average temperature gradient and the laminate thickness at the glass transition temperature. Moreover,

it should be noted that the average temperature gradient at the plateau of peak temperature was the highest for the lowest value of compaction pressure (i.e., 325 × 10² Pa). This high average temperature gradient implies a weak interlaminar consolidation because of the too low compaction pressure.

In the same manner, the result of the laminate thickness evolution shows that the lowest compaction pressure led to a higher laminate thickness during the whole consolidation cycle. This implies the presence of air layers at the interlaminar interfaces which are likely to swell during the heating phase. Finally, the sudden change of the average temperature gradient at the glass transition temperature is also related with the change of the thermal conductivity of the air layers at the interlaminar interfaces, which can be correlated with the compaction pressure according to Equation 1 adopted by Wu et al. (Cassidy and Monaghan, 1994).

$$k_{air} = \frac{k_0}{1 + \frac{7.6 \times 10^{-5} T}{PD}} \quad (1)$$

where $k_{air}$ is the thermal conductivity of air, $k_0$ is the thermal conductivity of air at the room temperature and pressure, T is the temperature (in K), P is the air pressure and D is the air gap height.

When the compaction pressure was important, the pressure in vacuum bag was below 5.0 × 10² Pa and the corresponding thermal conductivity of air was around 1 mW/(m·K). This low value increased the magnitude of the average temperature gradient. On the contrary, when the compaction pressure was around 325 × 10² Pa, the thermal conductivity of air in the vacuum bag and in the gaps at the interlaminar interfaces was around 25 mW/(m·K) at room temperature. The difference of the thermal conductivity between the prepreg plies and the interlaminar zones was not big enough to highlight the intimate contact phenomenon.

As a result, it has been found that thermoplastic laminates can be consolidated under low pressure (< 1.0 × 10⁵ Pa) and that under these conditions the in-situ monitoring is helpful to improve the understanding of the interlaminar consolidation.

In the oven heating method, the temperature distribution is more uniform than in the heating plate method. Hence, the

<figure><img src="image_405.png" /></figure>
influence from the non-uniform heating on the interlaminar consolidation can be minimized and eventually ruled out. The interlaminar consolidation quality was assessed by ILSS tests. **Figure 12** shows the results of the interlaminar shear strength of the laminates against the compaction pressure which was applied during the consolidation. The compaction pressure has a huge influence on the final consolidation quality of the material. The ILSS changes from 21 MPa to 107 MPa when the compaction pressure increases from $25 \times 10^2$ Pa to 980 $\times 10^2$ Pa. Moreover, above $700 \times 10^2$ Pa the ILSS values are superior to 100 MPa which is so close to the reference value, viz. 110 MPa (from the compression molded samples). The failure modes of the laminates were more complex to identify. Indeed, the delamination mostly occurred at very low compaction pressure ($-10 \times 10^2$ Pa) whereas it could also take place for the laminates consolidated at $-825 \times 10^2$ Pa (see

**Figure 13A**). However, this failure mode was relatively rarely observed. Most of samples showed large plastic deformation with or without a small drop of stress after the elastic region (see **Figure 13B**) which can reveal the presence of crack propagation (van Rijswijk et al., 2009; Vieille et al., 2011; See et al., 2015; Obande et al., 2019). According to these results the failure mode cannot be directly related with the pressure level of the consolidation.

These results highlight that high interlaminar consolidation quality can be achieved by Out-Of-Autoclave (OOA) or Vacuum-Bag-Only (VBO) manufacturing techniques.

## Heat Transfer Simulation Result

The four parameters of the sigmoid function for the thermal conductivity of interface layers (see **Figure 4**) have been identified. The temperature range ($dT$) was 22 K. The inflection temperature ($T_r$ = 440 K) corresponded to the glass transition temperature and the improvement of conductivity is associated with the matrix phase change. The initial conductivity ($k_{init}$) was $0.003 \text{ W} \cdot \text{m}^{-1} \cdot \text{K}^{-1}$ which corresponded to the air conductivity under very low pressure ($1-5 \times 10^2$ Pa). The final conductivity ($k_{fin}$) was $0.032 \text{ W} \cdot \text{m}^{-1} \cdot \text{K}^{-1}$ which was still lower than the matrix conductivity, because the intimate contact was not complete under the melting point. **Figure 14** presents the experimental result and the simulation result for the temperature evolution at the top layer during the consolidation process. We can see that the fast increase of the temperature at the top layer above the glass transition temperature was well-reproduced by the numerical simulation considering the change of thermal contact resistance at the interlaminar interfaces. As the amorphous zones of the matrix softened above the glass transition temperature, the surface roughness was deformed and flattened out to increase the contact area between the adjacent layers. Subsequently, the heat conduction path was enlarged and the thermal contact resistance was decreased. As a result, the heat transfer was enhanced and the temperature at the top layer was fast increased.

<figure><img src="image_406.png" /></figure>

<figure><img src="image_407.png" /></figure>