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# Smoothing and Cicatrization of Isotactic Polypropylene/Fe 3 O 4 Nanocomposites via Magnetic Hyperthermia

Pablo Griffiths, Simon Fritz, Laura Ea, Florent Dalmas, Jérôme Adrien, Vivien Truchot, Françoise Méchin, Cyril Benoit, Fabien Dutertre, Jean-Charles Majesté, et al.

► To cite this version:

Pablo Griffiths, Simon Fritz, Laura Ea, Florent Dalmas, Jérôme Adrien, et al. "Smoothing and Cicatrization of Isotactic Polypropylene/Fe 3 O 4 Nanocomposites via Magnetic Hyperthermia." ACS Applied Polymer Materials, 2024, 6 (23), pp.14084-14094. 10.1021/acsapm.3c03205. http://onlinelibrary.wiley.com/doi/abs/10.1021/acsapm.3c03205

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Smoothing & Cicatrization of iPP/Fe₃O₄
Nanocomposites via Magnetic Hyperthermia

Pablo Griffiths,† Simon Fritz,† Laura Ea,† Florent Dalmas,† Jérôme Adrien,†

Vivien Truchot,‡ Françoise Méchin,‡ Cyril Benoit,¶ Fabien Dutertre,¶

Jean-Charles Majesté,¶ Niels Pernoux,§ Jerome Grando,‖ Jerome Chevalier,† and

Guilhem P. Baeza*,†

†INSA Lyon, Universite Claude Bernard Lyon 1, CNRS, MATEIS, UMR5510, 69621,
Villeurbanne, France

‡INSA Lyon, Universite Claude Bernard Lyon 1, Universite Jean Monnet, CNRS, UMR
5223, Ingénierie des Matériaux Polymères, F-69621, Villeurbanne, France

¶Université de Lyon, CNRS, Université Claude Bernard Lyon 1, INSA Lyon, Université Jean Monnet, UMR 5223, Ingénierie des Matériaux Polymères F-42023, Saint-Etienne
Cédex, France

§Lines Manufacturing, 67810, Holtzheim, France

‖Plastic Omnium Intelligent Exterior System, 01150, Sainte-Julie, France

E-mail: guilhem.baeza@insa-lyon.fr
Abstract

We prepare industrially relevant magneto-responsive thermoplastic nanocomposites capable of being cicatrized and smoothed after additive manufacturing through the application of an oscillatory magnetic field (OMF). The materials are made of an iso-tactic polypropylene (iPP) matrix filled with magnetite nanoparticles (2-22 wt.%, 75 nm in diameter) synthesized from steel waste, providing them with a limited ecological impact on top of their ability to be repaired. Nanoparticles are found to have no significant impact on the thermal properties of the iPP, which allows to compare directly the magneto-thermal effect measured on the different nanocomposites. Beyond the primary temperature increase generated by magnetic hysteresis loss, we evidence that the OMF irradiation triggers a second heating mechanism from the iPP melting. This phenomenon, which was assigned to nanoparticles magnetization and subsequent rotation causing high-frequency mechanical friction, is here investigated in a systematic way. Our results indicate that while the specific power generated by nanoparticles friction is (expectedly) proportional to the irradiation time, it is independent of the nanoparticles content as long as the temperature is well-above the polymer melting point. These observations suggest therefore that the (local) filler-polymer interfacial rheology dictates the amount of heat generated through friction. From an application point of view, 7 wt.% of nanoparticles is found to be enough to induce iPP melting from magneto-thermal effect, which enables the post-processing of hot-pressed and 3D-printed specimen through "cicatrization" and "smoothing" experiments. In the former case, rewelding samples cut into two pieces is found to provide similar Young modulus and yield point as in native hot-pressed samples (exhibiting however a lower strain at failure). In the latter case and beyond the improved specimen appearance, smoothing is found to double both the stress and strain at failure of large 3D-printed samples that present nevertheless significantly lower properties than their hot-pressed counterparts.

**Keywords:** magnetic hyperthermia, magnetothermal effect, nanocomposites, healing, polypropylene, magnetite, 3D-printing
Introduction

Beyond every types of *recycling* including chemical and mechanical approaches, increasing
the lifespan of common plastic and rubber-based materials is one of the greatest challenges to
achieve a more sustainable development. In this context, spreading *reusing* and/or *repairing*
philosophies is crucial and has attracted the attention of many research groups in the two
last decades. A well-known example (from 2008) is that of self-healing soft matter where the
control of inter-molecular bonds provides "supramolecular" networks with the astonishing
ability to auto-repair at room temperature without external stimulus. ¹ Another emblematic
example from the same group relies on inorganic (silica) particles serving as stitches to cica-
trize gels and biological tissues. ² Since then, a wide range of supramolecular networks based
on small molecules have emerged, including vitrimers, covalent adaptable networks, and dis-
sociative networks, ³ offering vast possibilities in terms of mechanical/rheological properties.
In addition, polymer-based supramolecular networks known as "sticky-" or "associative-"
polymers have also been extensively studied for their rich rheological behavior, 4,5 opening
promising perspectives both in terms of processing and self-repairability. 6-8

It seems nevertheless that these major advances in chemistry did not materialize yet
into daily applications where self-healing materials remain sparsely used. While this lack
of societal impact certainly resides in the cost/complexity of the materials involved, other
issues regard their stability vs. thermo-oxidation⁹ and humidity. ¹⁰ In addition, the required
molecular mobility ensuring self-healing is most often antagonistic to high mechanical per-
formances, particularly at high deformation, which drastically limits the fields of application.
An interesting example is that of Bouteiller et al. who have tentatively addressed these lim-
itations by modifying chemically a commercial (relatively strong and tough) thermoplastic
polyurethane in order to enable its self-healing without major loss of mechanical properties. ¹¹

Apart from increasing the complexity of chemical architectures, materials healing can
be envisaged through the application of various physical stimuli including but not limited
to convective heat, microwaves,¹²,¹³ laser radiations¹⁴,¹⁵ and magnetic oscillatory fields.¹⁶⁻²⁰
In the latter case, stimulus-healing is achieved by irradiating micro- or nanoparticles (NPs) embedded into a thermoplastic polymer matrix with an oscillatory magnetic field (OMF) of several hundreds of kHz and 5-50 mT, generating heat through magnetic hysteresis loss. The physical mechanism at the origin of this loss depends on the particles' size, shape and nature.²¹ It consists of Néel (and possibly Brown) relaxation(s)²²,²³ into single magnetic domain particles whereas it is dominated by walls motion into multi-domain ones.²⁴ On the other hand, eddy currents only appear in the case of macroscopic electric conductors (> 1 mm). The advantages of magnetic hyperthermia vs. other healing methods reside in the combination of the following facts: i) the treatment can be performed in a contactless manner, including without visual path, ii) the heat can be controlled by tuning both the OMF characteristics and the sample formulation, iii) the absence of susceptor ensures low thermal inertia and minimize energy consumption and iv) the temperature field can be modified based on the NPs localization.

While magnetic mechanisms are activated as soon as the OMF is turned ON and keep on producing heat as long as it is not turned OFF, mechanical friction between NPs and the polymer matrix has been identified as an extra-heating mechanism. It was first suspected in hydrogels,²⁵ semi-crystalline homopolymers²⁶ and supramolecular networks,²⁷ and recently clarified through X-ray scattering techniques in thermoplastic polyurethanes (TPU).²⁸ In every cases, the sudden drop of viscosity (e.g. melting) following the temperature rise primarily caused by magnetic hysteresis loss triggers the particles motion, enabling them to rotate at the frequency of the OMF. In addition, their progressive polarization may also lead to translation resulting into the formation of dipolar chains.¹⁷,²⁸ Overall, melting-triggered friction modifies the time evolution of the material's temperature, denoted $T(t)$ hereafter, from a simple "Box-Lucas" exponential rise profile²⁹ to a "double-bump" one.¹⁷ The combination of both magnetic and mechanical effects on $T(t)$ has been recently theorized by Helbig et al.³⁰ in the case of superparamagnetic, isolated and spherical NPs dispersed into a Newtonian fluid, which in spite of the advanced physical description, represents a canonical (i.e., simple)
example of responsive material. Extending this model to nanocomposites made of a rheologically complex matrix like a homopolymer melt filled with possibly non-spherical³¹,³² and aggregated NPs³³ represents a great challenge for the coming years. Quantitative predictions of $T(t)$ curves is actually expected to strongly contribute to the fundamental understanding of NPs rotational diffusion in complex soft matter.

In this work, we focus on the peculiar case of isotactic polypropylene (iPP) that is massively used in the transport industry for its low cost, high mechanical properties and lightness to design car bumpers and body elements.³⁴,³⁵ Conversely to our previous work dealing with a relatively short TPU matrix (weight averaged molar mass $\overline{M_w}$ = 120 kg mol⁻¹ and molar mass dispersity, D= 2) characterized by a low zero-shear viscosity of about $\eta_0$ = 100 Pa s,¹⁷ the typical average molecular weight used for iPP based engineering goods is $\overline{M_w}$ = 200 – 300 kg mol⁻¹ and D ≈ 4, resulting in a melt viscosity of a few kPa s right above the melting point. From an application point of view, this difference is crucial since an efficient cicatization can only be performed on the condition that the polymer mobility is high enough, corresponding to a rapid macroscopic flow. We demonstrate below that iPP healing is indeed achievable, opening the way to further industrial development.

The article is organized as follows: the materials and the wide range of experimental techniques used to characterize the composites are described in section II. The results obtained on the neat iPP matrix and four composites containing a growing content of NPs are presented and discussed in section III. Thermal investigations by DSC, magnetic hyperthermia curves ($T(t)$) and rheological dynamic frequency sweeps first allow to characterize i) the ability of the material to generate heat under the OMF and ii) flow rapidly enough above the melting point. Based on these fundamental results, we then focus on the most industrially relevant formulation on which we perform both cicatization and smoothing procedure. The former consists of re-welding specimens cut into two pieces while the latter aims to reduce the surface roughness of 3D-printed parts. The performance of these two "healing" methods are evaluated through uniaxial tensile tests and scanning electronic microscopy.
# Materials and Methods

## Materials

### Polymer matrix

The polymer matrix is a commercial isotactic polypropylene (iPP) graciously provided by Total (France). Its weight average molar mass is $\overline{M_w}$ = 199 kg mol$^{-1}$ and its molar mass dispersity is $D= 4.5$. The corresponding steric exclusion chromatography elution curve is provided in Supporting Information (SI) **Section 1**.

### Magnetic fillers

"PGL" (commercial name) magnetite nanoparticles ($Fe_3O_4$) were purchased from Hymagin' (France). They mainly consist of elementary particles of $d_{50}$ = 75 nm forming unbreakable aggregates of a few hundreds of nm in diameter. The main interest of this powder resides in its synthesis route involving steel waste as a raw material. The patented technology$^{36}$ is entirely based on reactions with water and does not require any volatile organic solvent, strong acid nor flammable gases, fitting well the sustainable development philosophy developped in the present work.$^{37}$ The powder's X-ray diffraction pattern, a thermogravimetric analysis and a Fourier transfrom infra-red spectrum are provided in **SI Section 1**. The two latter analyses suggest the presence of small amount of organic species at the powder surface in spite of the organics-free synthesis route, possibly improving the powder dispersion in the iPP matrix.

### Nanocomposites compounding (small scale)

iPP/$Fe_3O_4$ nanocomposites were prepared by twin-screw extrusion in a 15 cm$^3$ chamber (DSM Xplore, The Netherlands). Appropriate amount of polymer and magnetic fillers were handed-mixed and poured into the extruder so that the nominal mass fraction of magnetite within the nanocomposites was 2, 7, 16 and 22 wt.%, corresponding to ca. 0.3, 1.3, 3.1 and 4.5 vol.%. The extruder chamber was pre-heated at 210 °C ensuring the quick melting of the
polymer. The rotor speed was set to 40 rpm and the mixing time to 6 min. The so-prepared nanocomposites were then hot-pressed at 210 °C for 3 min and slowly cooled down in air.

The film thickness was fixed to 1 mm, serving to prepare samples for magnetic hyperthermia characterization as well as rheological and tensile tests.

## Nanocomposites compounding (large scale) and additive manufacturing

For large-scale 3D-printing (see below), larger quantities of nanocomposites were prepared using a corotating twin-screw extruder EUROLAB 16 (Thermofisher, United-Kingdom). The screw diameter and the length/diameter ratio were $D = 16$ mm and $L/D = 40$, respectively. The device was divided into 10 sectors starting from the *feeding* zones to the 4 mm diameter *exit die*. The barrel temperature was fixed to 210 °C, with a gradual increase from room temperature to 150 °C between the *feeding* zones and zone 2. iPP pellets and $Fe_3O_4$ NPs were simultaneously introduced at feeding rates of 880 g h$^{-1}$ and 120 g h$^{-1}$ ($x_{NP} \approx 7$ wt.%) using two twin-screw gravimetric feeders (Brabender and Coperion K-Tron). A specific screw profile, consisting of various forwarding screw elements (*feeding* and *conveying* zones) and two successive *mixing* zones was designed to optimize the mixing procedure. Details on the twin-screw configuration are reported in **SI Section 2**. After cooling at room temperature, the extruded strand was finally pelletized to feed the 3D-printer.

Pellets were processed through the additive manufacturing system Lines SONIC, (Lines Manufacturing, France), capable of producing parts up to 2500 mm in all $\vec{X}, \vec{Y}, \vec{Z}$ directions. This three-axis gantry high-speed additive manufacturing production machine was equipped with a Fused Granular Fabrication (FGF) technology consisting of a raw material processing unit coupled with a screw extrusion setup and a melted material deposition device that enables high speed printing up to a theoretical value of 2000 mm s$^{-1}$. The actual deposition speed for iPP based nanocomposites was however set to 800 mm s$^{-1}$. Similarly as for the compounding step, the screw temperature was fixed to 210 °C. Once melted, the material entered the deposition unit where the temperature (210 °C), flow rate (125 cm$^3$ min$^{-1}$) and
flow pressure (16 bars) were carefully monitored and regulated before coming out of the 3.5
mm diameter nozzle.

## Methods

### Differential Scanning Calorimetry (DSC)

Experiments were run on a DSC-8000 (PerkinElmer, USA) under a nitrogen flow of 20 mL min⁻¹ at 5 °C min⁻¹. Samples were first cooled down to 20 °C and heated up to 250 °C (50 – 200 °C displayed). This cycle was then repeated once, resulting in the measurements displayed in Figure 1. Additional thermograms for different thermal history including 3D-printed samples are available in **SI Section 3**.

### X-ray Diffraction (XRD)

X-ray diffraction measurements were performed at room temperature on a D8 advance apparatus (Bruker, USA) using Cu Kα₁ radiation. The diffraction pattern of the PGL Fe₃O₄ powder was recorded in the 2θ range of 20 – 90 °.

### Scanning Electron Microscopy (SEM) and Focused Ion Beam (FIB)

For large magnification microstructural observations (Figure 6), sectioning and imaging of the nanocomposites were investigated using a dual column focused ion beam (FIB)-scanning electron microscope (SEM) ZEISS NVision40. A Ga²⁺ ion beam accelerated at 30 kV was used. A two-step milling procedure was employed on bulk nanocomposites that were metalized with gold beforehand. First, a bulk trapezoid was milled at high current beam (13 nA) so that the shorter face could be imaged by the electron beam up to a minimum depth of 25 µm. Then, a final polishing of the observed surface was carried out with a fine current beam (150 pA). SEM images of the polished surface were then recorded with an accelerating voltage of 5 kV, using an in-lens secondary electron (SE) detector. Such imaging conditions
allow observation of insulating materials with minimal charging effects, high resolution, and
a good contrast between the inorganic and organic phases.³⁸

Low magnification micrographs of the 3D-printed specimens (Figure 8) were obtained using a VEGA3 scanning electron microscope (TESCAN, USA), operating at a 15 kV acceleration voltage with a secondary electron detector. Transverse cross-sections of fractured tensile test specimens were analysed with or without magnetic induction smoothing. Additionally, longitudinal sections were observed on a cryocut sample after undergoing smoothing treatment.

## X-ray tomography

X-ray tomography was carried out at room temperature on a V-TomeX device (GE Phoenix X-ray GmbH, USA) with a 2520 V detector from Varian. We used 80 kV and 280 µA as X-ray parameters, placing the sample vertically with a voxel size of 3 µm. Nine hundreds radiographs with an exposure time of 333 ms and 3 radiographs averaged at each step were recorded during the 360° rotation of the sample for a total acquisition time of 20 min.

## Rheology

Dynamic frequency sweep measurements were performed with a strain-controlled rheometer (ARES, 2KFRTN1 from Rheometric Scientific, currently TA, USA) using stainless steel parallel plates of 8 mm diameter. The samples thickness was set to *ca.* 1.0 mm. The temperature was fixed to 210 °C. The angular frequency was varied from $\omega = 10^2$ to $10^{-1}$ rad s$^{-1}$. The strain amplitude was fixed to $\gamma = 3 \%$ for both iPP and M7 samples.

## Tensile Tests

Uniaxial tensile tests were conducted on dog-bone shaped specimen at room temperature with a constant displacement rate of 2 mm min$^{-1}$ using a 1/ME (MTS, USA) apparatus equipped with a 5 kN load cell. Hot-pressed specimens were punched out from thin films with
working dimensions 26.0 × 4.0 × 1.0 mm³. Similarly, 3D-printed specimen were punched
out from thick sheets resulting in dimensions of 26.0 × 4.0 × 3.4 mm³. A minimum of three
tensile tests was conducted for each type of sample. The variability of the results is displayed
through shaded rectangles in Figures 5 and 7.

Induction heating and thermal imaging

Induction heating was performed with a master controller v3+ apparatus (CEIA, Italy) connected to a two-turns water-refrigerated flat spiral inductor. A 0.1 mm thick Teflon-based composite tape covered the inductor to protect it from material flowing. The frequency of the magnetic field was 855 kHz and its maximum amplitude in the sample was estimated to 10 mT through Comsol simulations.¹⁷ The magnetic field was generated by a current of 19 A in the coil and delivered through pulses of duration *t*<sub>pulse</sub> every *t*<sub>period</sub>. The apparent power of the inductor was varied from *t*<sub>pulse</sub>/*t*<sub>period</sub> ranging from 17 to 100 %. Disk-shaped samples (6 mm diameter - 1 mm thick) were submitted to the magnetic field and their temperature was recorded with a PI 640i infrared camera (Optris, Germany) equipped with a 15 × 11° lens characterized by an optical resolution of 640 × 480 pixels and a measurement rate of 32 Hz. Temperatures reported in this work systematically refer to the maximal temperature recorded by the camera (at the center of the sample). Data were processed with the software *OptrisPIX Connect* in the 20 – 250 °C range. We used the default camera calibration from Optris and checked that the temperature well agreed with the value provided by a thermocouple (±2%).³⁹ Cicatrization of hot-pressed dumbbell specimens consisted of cutting samples into two identical pieces and placing them in a silicone mold of corresponding dimensions equipped with an anti-adhesive bottom part. After a minute of induction heating the healed specimens were extracted from the mold and tested (tensile test or SEM). Smoothing of 3D-printed specimens consisted of a 10 min superficial induction heating performed over the whole operational length and on both sides of the samples.
# Results and Discussion

## Magnetothermal properties

DSC thermograms of the neat iPP and corresponding nanocomposites are reported in Figure 1, demonstrating that 2-22 wt.% (1-5 vol.%) of Fe₃O₄ NPs do not affect significantly the crystallization of the polymer, for which the melting temperature *T*<sub>m</sub> and the crystalline ratio *X*<sub>c</sub> vary erratically within 161 – 164 °C and 45 – 50 % intervals, respectively. (*X*<sub>c</sub> estimations are based on a specific enthalpy of neat crystalline iPP of Δ*H*<sup>*</sup> = 168 J g<sup>-1</sup>, see ref. ⁴⁰). Nevertheless, while the melting peak is rather bimodal at low NPs content, the higher temperature side of the peak seems to disappear progressively with increasing the fraction of NPs, suggesting a slightly hindered spherulites growth. Overall, the weak dependence of the thermograms upon varying the filler content is expected because of the limited amount of fillers and the polar surface of the Fe₃O₄ NPs embedded into a quite apolar iPP matrix.

<figure><img src="image_2.png" /></figure>

Figure 1: Second heating ramp differential scanning calorimetry thermograms (endo up). Impact of the weight fraction of NPs. Nanocomposites data were vertically shifted for clarity, the shift value (between 0 and 2 Wg⁻¹) is indicated for each sample. The heating rate was fixed to 5 °C min⁻¹. Raw thermograms were normalized by the mass of iPP contained in the samples.
The almost unchanged thermal behavior of the matrix (heat capacity and melting point) makes it possible to compare directly the evolution of the $T(t)$ curves obtained through magnetic hyperthermia, i.e., when the sample is submitted to an oscillatory magnetic field starting from room temperature. The corresponding results are displayed in Figures 2a-d emphasizing the impact of the pulse duration and the NPs content (2d-inset). As expected, increasing the ratio $t_{pulse}/t_{period}$ or the weight fraction of NPs leads to a faster rise of the temperature that possibly reaches the melting point of the matrix (grey zone). In this case, an extra heating mode originating from NPs friction in the viscous polymer system- atically emerges,²⁸ which makes deviate the $T(t)$ curves from a simple exponential rise (or "Box-Lucas" profile) corresponding to the solution of the 0th dimension heat equation for a constant energy flux.²⁹ All the $T(t)$ curves can be satisfyingly fitted by using the following set of equations

$$ 
\begin{align} 
T(t < t_m) &= T_0 + \Delta T_1 [1 - \exp(-(t/\tau_1)^\beta)] \tag{1} \\ 
T(t > t_m) &= T_0 + \Delta T_1 [1 - \exp(-(t/\tau_1)^\beta)] + \Delta T_2 [1 - \exp(-(t-t_m)/\tau_2)] \nonumber 
\end{align} 
$$

where $t_m$ is the time required to activate the second heating mode, i.e., to soften enough the iPP to trigger NPs motion, $T_0$ is the room temperature, $\Delta T_1$ and $\Delta T_2$ are the temperature amplitudes related to magnetic losses and friction respectively, $\tau_1$ and $\tau_2$ are the correspond- ing characteristic times, and $\beta$ (always comprised within the 0.85–1 interval) is an exponent enabling to slightly "stretch" the first exponential term providing somewhat an error bar on the $T(t)$ profile. All the fit parameters are gathered into Table 1. Here, it is worth to note that the second heating mode systematically emerges from ca. 140 °C, i.e., 10 °C below the onset of the melting peak measured in DSC (Figure 1), which suggests that NPs rotation does not require the matrix to be completely molten. In fact, NPs motion below the melting point could arise from crystalline re-ordering and metastable phases denoted as "rotator phases" or "mesophases" as observed in canonical n-alkanes 6–8 °C below $T_m$ 41–43 as well as in iPP and similar polymers.44,45 In addition, the second mode is particularly well-visible
in comparison to our previous work focusing on TPU-based nanocomposites,¹⁷ most likely because of the relatively narrower melting peak and the higher degree of crystallinity of the iPP matrix that implies a gigantic and sudden drop of viscosity at the melting point. By applying the two versions of Equation 1 (with $\beta = 1$), one can then extract the specific power [also called specific absorption rate (SAR) in W g⁻¹] associated to each heating mode that can be estimated through

$$ SAR_i = \frac{c_{p-i} \Delta T_i}{x_{NP} \tau_i} \qquad (2) $$

where $c_{p-i}$ is the specific heat capacity of the nanocomposites that is measured by DSC at 25 °C and 180 °C, corresponding to the solid- and liquid-states of iPP respectively relevant to quantify the two heating mechanisms. At 25 °C, $c_{p-1} = 1.85, 1.84, 1.76, 1.76$ and $1.65$ J g⁻¹ K⁻¹ while at 180 °C, $c_{p-2} = 3.24, 3.12, 3.12, 3.24$ and $3.05$ J g⁻¹ K⁻¹, respectively for iPP, M2, M7, M16 and M22 specimens. The former values are in fair agreement with the room temperature specific heat capacity tabulated for neat iPP ($c_p^{PP} = 1.62$ J g⁻¹ K⁻¹)⁴⁶ and magnetite ($c_p^{Fe_3O_4} = 0.65$ J g⁻¹ K⁻¹)⁴⁷ as well as with the "composites values" provided by a simple mixing rule. Besides, $\frac{\Delta T_i}{\tau_i} = \frac{dT}{dt}|_{t=0}$, the latter term being generally used to evaluate SARs graphically without any data fitting. Note that although $\Delta T_i$ and $\tau_i$ do not systematically show monotonic trends with $x_{NP}$ and $t_{pulse}/t_{period}$ in Table 1 because of their relative freedom in achieving a good fit on a restricted temperature range, the corresponding SARs that quantify the two physical mechanisms do evolve in a consistent way with the pulse duration and the NPs content (Figure 3).

As one may have anticipated, Figure 3a indicates that SAR₁ is systematically proportional to $t_{pulse}/t_{period}$ and does not vary significantly with the NP content. The proportionality was expected since the magnetic losses are cumulated based on the number of magnetization vs. incident field ($M = f(H)$) hysteresis loops achieved by the system, itself being directly proportional to the duration of the pulse, i.e., to the $t_{pulse}/t_{period}$ ratio. On the other hand, the constancy of SAR₁ as a function of $x_{NP}$ (Figure 3c) indicates that the
cyclic magnetization of the NPs is not impacted by their dipolar interactions, which seems
reasonable given the low NP volume fraction ($\Phi_{NP} < 5$ vol.%). From a quantitative point of
view, the quite high values of $SAR_1$ reaching up to ca. 1000 W g$^{-1}$ are undoubtedly related
to the multi-domain ferrimagnetic nature of the nanoparticles and the high frequency of the
OMF.

At first sight, it appears that the same logic can be applied to $SAR_2$ in Figure 3b. In
this case, the specific power generated by friction is seen to be proportional to the pulse
duration suggesting that the NPs continuously rotate within the polymer matrix as long
as the magnetic field is ON, and stop otherwise. In addition, $SAR_2$ is also observed to be
independent of the NPs content above $x_{NP} = 7$ wt.% (i.e., in M7, M16 and M22), which
indicates that NPs are not interfering mechanically with each other for experiments shorter
than two minutes. Conversely, when non-being 0 (absence of melting), $SAR_2$ values are
much lower in M2, being typically five times less than in the other nanocomposites. This
observation can alternatively be made in Figure 3d, where $SAR_2$ values increase very much
between $x_{NP} = 2$ and 7 wt.% before stabilizing. In spite of the lack of direct experimental
evidence, we assign this strong deviation to the fact that, for $t_{pulse}/t_{period} = 83$ and 100 %, the
steady state temperature in M2 mainly belongs to the polymer melting zone measured from
DSC (see Figure 2a), i.e., that the iPP matrix is not completely molten in the experimental
time window, limiting therefore the rotational diffusion of NPs.

Interestingly, although $SAR_1$ is systematically ca. twice as higher than $SAR_2$, the fact
that the two parameters are within the same order of magnitude indicates that the fric-
tion mechanism is highly relevant to control the temperature of polymeric fluid submitted
to high-frequency magnetic hyperthermia. Said differently, $SAR_2$ represents here ca. 1/3
of the whole heat generated by the material submitted to the OMF. (Note however that
$SAR_2$ values are hardly predictable since they depend on the amplitude of the multi-domain
NPs deflection angle and the rotational friction coefficient that are not directly accessible
experimentally in non-Newtonian fluids).<sup>28,30</sup>
<figure><img src="image_3.png" /><figcaption>Figure 2: Temperature vs. time during induction heating for different values of $t_{pulse}/t_{period}$ in a) M2, b) M7, c) M16 and d) M22 nanocomposites. Inset: Impact of $x_{NP}$ for $t_{pulse}/t_{period} = 33\%$. Grey solid lines are fit to the data with Eq. 1. Grey zones represent the temperature range corresponding to the melting peak measured in DSC.</figcaption></figure>
<figure><img src="image_4.png" /><figcaption>Figure 3: SAR values corresponding to magnetic hysteresis losses ($SAR_1$) and NPs friction ($SAR_2$) for various $t_{pulse}/t_{period}$ ratios (a-b) and various content in NPs (c-d). The black solid line in (a) is a fit with an affine function. Note that NPs friction is not detected in M2 for $t_{pulse}/t_{period} < 83\%$, i.e., $T_m$ is not reached, resulting in $SAR_2 = 0$ in (b). Error bars are estimated from thermogravimetric measurements only, their higher and lower limits corresponds to the average value of $x_{NP}$ plus or minus the standard deviation, respectively.</figcaption></figure>
<table><caption>Table 1: Samples characteristics and fit parameters obtained from applying Eq. 1 on $T(t)$ data presented in Figure 2. $T_0$ (not reported here) is a free parameter varying erratically between 18.4 and 20.5 °C. $x_{NP}$ is measured by thermogravimetric analysis and averaged on three samples. $\Phi_{NP}$ is deduced from $x_{NP}$ assuming iPP and $Fe_3O_4$ densities of 0.9 and 5.1 g cm$^{-3}$, respectively. Star symbols (*) point data fits providing "surprising" and meaningless individual $\Delta T_i$ and $\tau_i$ values because of the limited observable temperature range ($T < 250$ °C) related to the IR camera restrictions and iPP degradation. Their ratio $\Delta T_i/\tau_i$, being proportional to $SAR_i$ (see Eq. 2) is however a relevant physical parameter to study, varying in a consistent way with $x_{NP}$ and $t_{pulse}/t_{period}$ (see Figure 3).</caption><thead><tr><th>Sample</th><th>x<sub>NP</sub></th><th>Φ<sub>NP</sub></th><th>X<sub>c</sub></th><th>t<sub>pulse</sub>/t<sub>period</sub></th><th>ΔT<sub>1</sub></th><th>ΔT<sub>2</sub></th><th>τ<sub>1</sub></th><th>τ<sub>2</sub></th><th>t<sub>m</sub></th><th>β</th></tr><tr><th>—</th><th>wt.%</th><th>vol.%</th><th>%</th><th>—</th><th>°C</th><th>°C</th><th>s</th><th>s</th><th>s</th><th>—</th></tr></thead><tbody><tr><td rowspan="5">iPP</td><td rowspan="5">0.0</td><td rowspan="5">0.0</td><td rowspan="5">50</td><td>17</td><td>33.2</td><td>—</td><td>14.2</td><td>—</td><td>—</td><td>0.88</td></tr><tr><td>33</td><td>65.2</td><td>—</td><td>17.2</td><td>—</td><td>—</td><td>0.88</td></tr><tr><td>50</td><td>96.8</td><td>—</td><td>15.8</td><td>—</td><td>—</td><td>0.85</td></tr><tr><td>67</td><td>115.2</td><td>—</td><td>17.2</td><td>—</td><td>—</td><td>0.85</td></tr><tr><td>83</td><td>132.4</td><td>25.9</td><td>15.2</td><td>58.7</td><td>76.4</td><td>0.89</td></tr><tr><td rowspan="5">M2</td><td rowspan="5">1.9</td><td rowspan="5">0.3</td><td rowspan="5">46</td><td>100</td><td>147.0</td><td>37.7</td><td>13.7</td><td>47.2</td><td>40.0</td><td>0.88</td></tr><tr><td>17</td><td>102.4</td><td>—</td><td>14.8</td><td>—</td><td>—</td><td>0.86</td></tr><tr><td>33</td><td>135.7</td><td>49.3</td><td>9.8</td><td>25.3</td><td>29.4</td><td>0.98</td></tr><tr><td>50</td><td>144.6</td><td>122.0</td><td>7.0</td><td>20.5</td><td>16.0</td><td>1.0</td></tr><tr><td>67</td><td>145.5</td><td>244.1</td><td>5.6</td><td>29.1</td><td>12.0</td><td>1.0</td></tr><tr><td rowspan="5">M7</td><td rowspan="5">6.9</td><td rowspan="5">1.3</td><td rowspan="5">46</td><td>83</td><td>203.7*</td><td>146.1*</td><td>5.6*</td><td>14.4*</td><td>7.8</td><td>0.8</td></tr><tr><td>100</td><td>146.4*</td><td>1054.2*</td><td>3.2*</td><td>71.6*</td><td>7.1</td><td>1.0</td></tr><tr><td>17</td><td>134.3</td><td>81</td><td>8.3</td><td>29.8</td><td>23.4</td><td>0.99</td></tr><tr><td>33</td><td>147.2</td><td>169.0</td><td>4.7</td><td>16.3</td><td>10.2</td><td>1.0</td></tr><tr><td>50</td><td>155.1</td><td>309.7</td><td>3.2</td><td>18.1</td><td>5.8</td><td>1.0</td></tr><tr><td rowspan="5">M16</td><td rowspan="5">15.4</td><td rowspan="5">3.1</td><td rowspan="5">45</td><td>67</td><td>158.5</td><td>337.0</td><td>2.5</td><td>14.6</td><td>4.3</td><td>1.0</td></tr><tr><td>83</td><td>164.9</td><td>984.5*</td><td>2.0</td><td>36.0*</td><td>3.3</td><td>1.0</td></tr><tr><td>100</td><td>173.2</td><td>1156.9*</td><td>1.7</td><td>40.0*</td><td>3.0</td><td>1.0</td></tr><tr><td>17</td><td>143.3</td><td>140.5</td><td>6.2</td><td>28.9</td><td>13.2</td><td>1.0</td></tr><tr><td>33</td><td>153.3</td><td>491.6</td><td>3.7</td><td>33.2</td><td>6.9</td><td>1.0</td></tr><tr><td rowspan="5">M22</td><td rowspan="5">21.2</td><td rowspan="5">4.5</td><td rowspan="5">46</td><td>50</td><td>160.4</td><td>372.7</td><td>2.5</td><td>15.6</td><td>4.4</td><td>1.0</td></tr><tr><td>67</td><td>161.5</td><td>367.5</td><td>1.8</td><td>10.4</td><td>3.1</td><td>1.0</td></tr><tr><td>83</td><td>156.5</td><td>2000*</td><td>1.9</td><td>52.2*</td><td>3.5</td><td>1.0</td></tr><tr><td>100</td><td>165.1</td><td>2000*</td><td>1.3</td><td>47*</td><td>2.1</td><td>1.0</td></tr></tbody></table>
Rheological behavior

Table 1 allows to identify the M7 sample as the most promising material for industrial applications since it appears that 1.3 vol.% of NPs is enough to trigger the iPP melting (almost) regardless of the "inductor power" (~ t$_{pulse}$/t$_{period}$ ). In other words, M7 has a *priori* the ability to be cicatrized and smoothed while being the formulation requiring the least amount of NPs, making it therefore the cheapest to produce and the most likely to flow rapidly. The latter property is actually investigated by performing rheological dynamic frequency sweeps on both the neat iPP and M7. Results are presented in Figure 4 where the storage (G') and loss (G'') dynamic moduli vs. the angular frequency measured at 210 °C are reported. As one may have anticipated, adding 1.3 vol.% of NPs is seen to have no significant effect on the rheological behavior of iPP, ensuring therefore M7 to flow at the same rate as the neat homopolymer, corresponding to a zero-shear viscosity η₀ ≈ 800 Pa s. The sample aspect is on the other hand drastically different, passing literally from white to black upon the addition of (only) 1.3 vol.% of Fe₃O₄ NPs (inset of Figure 4), which possibly limits the applications of these materials when used as visible parts.
<figure><img src="image_5.png" /><figcaption>Figure 4: Storage ($G'$, solid symbols) and loss ($G''$, open symbols) dynamic shear moduli vs. angular frequency for neat iPP (squares) and M7 (triangles) samples measured at 210 °C. The canonical $G'$ and $G''$ slopes corresponding to a Maxwellian fluid are displayed for comparison. Inset: photograph of the neat iPP and the M7 sample emphasizing the drastic change in color with only 1.3 vol.% in Fe₃O₄ NPs.</figcaption></figure>
Cicatrization of hot-pressed specimen

Based on the nanocomposites ability to i) heat above the melting point and ii) flow rapidly, we chose the M7 formulation to produce hot-pressed dumbbell specimens with the aim to test mechanical properties before and after a cicatrization procedure. The latter consisted of cutting a specimen into two pieces with a razor blade and rewelding the material confined into a silicon mold by using localized magnetic hyperthermia similarly as in ref. 17 Tensile tests performed on series of "hot-pressed" samples and their "healed" counterparts are presented in Figure 5. On one hand, "hot-pressed" samples exhibit a Young modulus close to 1 GPa, a yield stress of 18 MPa and an average strain at failure of about 5%. On the other hand, although "healed" samples present similar characteristics in the linear regime, which confirms that a complete cicatrization of the material has been achieved, their failure occur right from the yield point (2-3%), suggesting that in spite of the absence of porosity (see X-ray tomography slides in Figure 6), the iPP chains did not have enough time to fully relax and recrystallize as in the bulk material. In fact, the failure systematically happens in the healed zone whereas it was randomly located in the "hot-pressed" samples - see inset Figure 5 where 1-5 and 6-10 specimen refer to "hot-pressed" and "healed" samples, respectively. At first sight, improving the strain at failure could therefore consist of increasing the healing time or decreasing the viscosity of the polymer. We expect that optimizing these parameters can certainly help but may not be sufficient to recover the original properties. In fact, we demonstrated that by using a similar procedure, responsive thermoplastic polyurethanes of ca. 10 times lower viscosity than the present iPP, also exhibited a significant decrease of their strain at failure after 1 min of cicatrization, passing from ca. 180% down to 120-150% according to the type of NPs. 17 The choice of the filler and their ability to diffuse within the matrix actually represents another pathway to explore. 16,28

The microstructure of a "hot-pressed" and a "healed" specimen investigated through X-ray tomography and FIB-SEM are presented in Figure 6. As mentioned above, low magnification X-ray tomography slices taken in the (X̄, Ȳ) plane presented in Figures 6a-b first
<figure><img src="image_6.png" /><figcaption>Figure 5: a) Tensile tests of hot-pressed and healed M7 samples. Shaded orange and blue zones indicate the dispersion of the strain at failure. The inset represents a series of hot-pressed (top) and healed (bottom) samples at failure. b-c) M7 dumbbell specimen configuration during cicatrization. Photographs are taken with white and IR light, respectively (see Section II for details).</figcaption></figure>
evidence that the NPs are well distributed within the iPP matrix after hot-pressing and
that this structure is preserved in the healed zone. Also, they do not show any particular
porosity as one may have expected from the healing of (partly) hydrophilic materials due
to sudden moisture evaporation.¹⁷ Besides, cross-section FIB-SEM micrographs taken in the
$(\vec{X}, \vec{Z})$ plane at a much higher magnification enabling single NP resolution suggest that the
healing procedure involving induction heating may have oriented NPs aggregates according
to the magnetic field lines. In fact, while NPs do not present any particular organization in
the hot-pressed sample (Figure 6c), the healed version indicate a slight vertical alignment of
the aggregates (Figure 6d) that is likely to become evident with increasing the filler content
(see e.g. refs¹⁷,²⁸).

<figure><img src="image_7.png" /><figcaption>Figure 6: Microstructure of the *hot-pressed* and *healed* M7 samples at various lengthscales. a-b) X-ray tomography slices and c-d) FIB-SEM micrographs. The magnetic field lines during healing are oriented vertically in (d).</figcaption></figure>
Smoothing 3D-parts

Beyond the production of hot-pressed (or equivalently injected) parts, additive manufactur-
ing achieved through Fused Deposition Modeling (FDM) of thermoplastic polymers repre-
sents a promising alternative to produce complex architectures likely to modify the design of
entire engineering systems in the transport industry. This technology however suffers from a
severe compromise between the quality of the printed parts and the time required to produce
them, being critically longer than with injection-based methods for large items (> 100 g).
Increasing the size of the nozzle and/or the speed of filament deposition to accelerate the
manufacturing are notably expected to lead to a larger number of critical defects possibly
resulting in catastrophic and premature failure of the operating materials. 48–50 Most common
defects include macroscopic anisotropy, 51 residual stress and porosity. 52,53 In this context,
magnetic hyperthermia represents an efficient and solvent-free method to smooth the rough
surface of 3D-printed parts with the aim to enhance both their mechanical performances and
their appearance.

Here, we have 3D-printed kilograms of M7 nanocomposite sheets at a speed relevant for industrial applications from which we have extracted 3.4 mm thick dumbbell specimen. Tensile tests were performed both on these ‘as printed’ samples and on their ‘smoothed’ counterparts and are presented in Figure 7. The first striking observation is the dramatically lower mechanical properties of the ‘as printed’ specimens with respect to their ‘hot-pressed’ analogs: i) their Young modulus is close to 700 MPa (vs. 1 GPa), ii) the stress at failure is always about 5.5 MPa (vs. 18 MPa), where a fragile (vs. ductile) rupture is observed, and iii) the strain at failure belongs to the 0.75-1.2% interval (vs. 5-7%). The impact of the smoothing procedure seems however quite significant, doubling both the strain and the stress at failure in a systematic way (3 specimens). For the sake of clarity, we have reported the main tensile tests characteristics of all the samples in Table 2.

These results can be rationalized through the scanning electronic microscopy experiments
presented in Figure 8, where an ‘*as printed*’ and a partly smoothed sample are investigated.
<figure><img src="image_8.png" /><figcaption>Figure 7: a) Tensile tests of ‘*as printed*’ and ‘*smoothed*’ M7 samples (dashed lines). Analog data for the hot-pressed sample is reported for comparison (solid line). Inset: ‘*as printed*’ and ‘*smoothed*’ specimen at failure. Shaded zones emphasize the dispersion of strain at failure. b) Photograph of the setup configuration used to smooth the 3D-printed samples. The processed zone visibly reflects more light.</figcaption></figure>
<table><caption>Table 2: Representative tensile properties of M7 samples including i) *hot-pressed* and *healed* samples, and ii) *as printed* and *smoothed* ones.</caption><thead><tr><th>processing</th><th>E<br>(MPa)</th><th>σ<sub>f</sub><br>(MPa)</th><th>ε<sub>f</sub><br>(-)</th></tr></thead><tbody><tr><td>hot-pressed</td><td>1095</td><td>18.0</td><td>0.051</td></tr><tr><td>healed</td><td>895</td><td>18.0</td><td>0.021</td></tr><tr><td>as printed (3D)</td><td>685 ±21</td><td>5.5 ±0.8</td><td>0.008 ±0.001</td></tr><tr><td>smoothed (3D)</td><td>463 ±0.6</td><td>8.9 ±0.1</td><td>0.019 ±0.00035</td></tr></tbody></table>

Figures 8a-b focuses on the cross section of the ‘*as printed*’ dumbbell specimen, clearly evidencing the presence of a large number of ca. 100 µm pores. This observation confirms therefore that obtaining a homogeneous material from rapid and large scale 3D-printing is far from being trivial and that, apart from the optimization of the printing parameters, a post-treatment is most often required to enhance the quality of the part. From a mechanical point of view, the porosity is expected to reduce the Young modulus and lower both the yield stress and yield strain as observed in Figure 7 (see ‘*hot-pressed*’ vs. ‘*as printed*’ samples’ performance). It is however interesting to remark that the critical defect size of iPP is expected to be $a_c \approx K_{Ic}^2/(\pi\sigma_f^2) \approx 1$ mm for mode I solicitation,⁵⁴ where $K_{Ic}$ is the corresponding plane strain fracture toughness, close to 4 MPa m¹/² and $\sigma_f$ is the yield stress. The value of $a_c$ is thus significantly larger than the pores’ size, which should result in a plastic deformation of the ‘*as printed*’ M7 sample before its catastrophic failure. The fact that the opposite trend is observed in Figure 7a suggests therefore that the pores are likely to be connected (forming larger defects) or that their inter-wall distance is smaller than their diameter (modifying the local stress intensity factor). In fact, we remind the reader that the nozzle diameter is 3.5 mm and that in spite of the fused filament compression during the printing, ca. 1 mm roughness remains at the surface of the specimen as evidenced in Figures 8c-d.

Remarkably, the same Figures however emphasize that these large surface heterogeneities are mostly erased when the samples are post-treated with the magnetic field (orange zones). In addition, the internal porosity is not detectable anymore, indicating that the smoothing procedure does act in an efficient way to enhance the quality of the specimen. Figure 8e
focuses on the interface between the healed and native zones of the specimen, revealing a clear
difference in terms of fracture pattern. While the (healed) top part exhibits some roughness
at large scale (≈ 100 µm), the (native) bottom part is clearly less homogeneous, with a
roughness lengthscale closer to 10 µm. These results support the mechanical tests presented
in Figure 7a where the dumbbell specimen were treated on both faces, undoubtedly reducing
the size of the largest defects present in the material and enhancing therefore its resistance
to failure. Nevertheless, the similar Young modulus observed on ‘as printed’ and ‘smoothed’
samples indicate that apart from the residual geometrical defects, the iPP structure might
have been modified by the additive manufacturing. (No significant difference is however
observed from DSC thermograms of ‘as printed’ and ‘hot-pressed’ samples - see SI Section
3).

<figure><img src="image_9.png" /><figcaption>Figure 8: SEM micrographs of 3D-printed samples after tensile failure. a-b) Cross-section of an ‘as printed’ sample that exhibits a uniform spatial distribution of milliporosities (blue). c-d-e) Partly smoothed sample highlighting the structural difference between native and smoothed zones (orange).</figcaption></figure>
Conclusion

iPP-based nanocomposites filled with various contents of Fe₃O₄ NPs (xNP = 2-22 wt.%) have been prepared through twin-screw extrusion with the aim to identify a relevant material for the transport industry endowed with the ability to be cicatrized and/or smoothed easily from magnetic hyperthermia. The first step consisted of a systematic characterization of the specific power generated by the NPs under a high frequency oscillatory magnetic field varying both the effective time pulse and the NPs content. As expected, magnetic hysteresis losses were observed to i) scale in a linear way with the former parameter and ii) be independent of the latter. More interestingly, our experiments unambiguously confirmed the presence of a second heating mechanism previously assigned to NPs-polymer friction triggered by the melting of the matrix. The specific power generated by friction was characterized for the first time evidencing its linear evolution with the time pulse while being quasi-independent of the NP content above a certain temperature threshold. Of a particular interest is the fact that the specific power generated through friction was found to be (only) three times lower than the one obtained through magnetic hysteresis losses evidencing that NPs motion is highly relevant to control the temperature of polymeric fluids submitted to magnetic hyperthermia.

From a practical point of view, we produced hot-pressed and 3D-printed specimen for tensile tests based on a nanocomposite containing 7 wt.% of Fe₃O₄ NPs, ensuring both rapid heating up to the iPP melting point and rapid flowing. The ability to heal was evaluated through a rewelding step performed after having cut the specimen into two parts. Tensile experiments evidenced similar stress at failure prior and after the cicatrization whereas the corresponding strain was systematically lower suggesting incomplete inter-diffusion of the polymer at the interface prior to its recrystallization. On the other hand, smoothing 3D-printed specimens (of significantly weaker properties than their hot-pressed counterparts) was found to double both their strain and stress at failure, representing a quite promising result for the development of 3D-printed parts post-processing.
**Associated Content**

**Supplementary Information:** 1) Raw materials: XRD pattern of the Fe₃O₄ NPs and SEC of the iPP, 2) Details on the twin screw extruder for large scale production, 3) DSC thermograms of samples with various thermal history and formulations.

**Authors information**

**Corresponding Author** : *G.P.B.: guilhem.baeza@insa-lyon.fr

**Acknowledgement**

All the authors thank Sophie Desmoures (So Sciences) for her help as manager of the "*Smart Polymer Bumper*" project as well as Nathalie Veiga (PO), Philippe Gilotte (PO), Manuel Collet (ECL), Rauf Garip (LINES), Emerik Poursilie (LINES) and Nassim Aliche (LINES) for fruitful discussions and technical support. G.P.B. thanks Lisa-Marie Pouilly, Camille Crouzet and Philippe le Bouteiller (Hymag'in) for their advice on the choice of the magnetite powder. P.G., J.C. and G.P.B. thank the *So Sciences Foundation* for funding allocated to MATEIS within the project *Smart Polymer Bumpers*. This work was supported as part of MANIOC project (Micro-Magnetic Manipulation of Block-Copolymers) funded by the french ANR [ANR-22-CE06-0006-01] as well as the french ministry of research and higher education (L.E. PhD thesis funding).
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