# Thermal stability of fluorosilicones under oxidative conditions

Leticia Pires, Alice Corfa, Vincent Ladmiral, Sylvain Caillol, Sébastien Roland, Emmanuel Richaud

► To cite this version:

Leticia Pires, Alice Corfa, Vincent Ladmiral, Sylvain Caillol, Sébastien Roland, et al. Thermal stability of fluorosilicones under oxidative conditions. Polymer Degradation and Stability, 2025, 243, pp.111772. 10.1016/j.polymdegradstab.2025.111772. hal-05370201

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# Thermal stability of fluorosilicones under oxidative conditions

Leticia Pires<sup>a</sup>, Alice Corfa<sup>b</sup>, Vincent Ladmiral<sup>b</sup>, Sylvain Caillol<sup>b</sup>, Sébastien Roland<sup>a</sup>, Emmanuel Richaud<sup>a,* ID</sup>

## ARTICLE INFO

Keywords:
Fluorosilicone
TGA
Degradation
Thermo-oxidation

## ABSTRACT

Using thermogravimetric analysis (TGA) this research investigates how polymer composition, temperature, and oxygen concentration influence degradation mechanisms of fluorosilicones at high temperatures. The findings reveal that even low oxygen levels significantly impact the degradation of model unfilled fluorosilicone homopolymer (HP), with further formulation changes reducing this effect. Overall, the copolymerization with dimethylsiloxane and vinylmethylsiloxane lower degradation kinetics. The stability is increased for model rubbers in which silica particles are incorporated, indicating that silica provides an antioxidative effect, and in cured samples obtained by peroxide crosslinking using dicumyl peroxide (DCP) combined with $CaCO_3$ or pure 2,5-dimethyl-2,5-di-(tert-butylperoxy)hexane (DBPH). Analysis of macromolecular changes using modulated differential scanning calorimetry (MDSC) on HP reveals distinct degradation behaviors under oxygen at 250 °C and 280 °C possibly due to the appearance of a heterogeneous mechanism. These findings offer valuable insights into the thermo-oxidative stability of fluorosilicones.

## 1. Introduction

Fluorosilicone rubbers are polysiloxane elastomers commonly used for sealing applications in environments with non-polar fluids and extreme temperature variations [1]. These rubbers are primarily composed of poly(3,3,3-(trifluoropropyl)methylsiloxane) (PTFPMS) but are often modified with additional components to optimize their performance. Typical modifications include introducing dimethylsiloxane and methylvinylsiloxane units (via copolymerization), and the addition of silica particles ($SiO_2$) as reinforcing filler. Inorganic nanoparticles and other additives can also be incorporated to improve high-temperature resistance, mechanical properties, or antioxidant performance [2-7]. These modifications result in materials with properties distinct from the basic homopolymer described in the scientific literature. While our previous work investigated the influence of each component on the thermal stability of fluorosilicones [8], this study aims to advance the understanding of their oxidative resistance by evaluating different formulations.

Purely thermal degradation primarily proceeds via depolymerization, driven by the high polarity of the Si-O bond and the flexibility of the -(Si-O)ₙ- segments [9], and it is accelerated at high temperatures by backbiting [10]. These features promote inter and intramolecular

redistribution reactions, releasing volatile cyclic siloxanes, which are thermodynamically favored over linear chains [11]. In contrast, thermo-oxidative degradation, typically studied via TGA temperature ramps, progresses through at least two main stages. First, at low temperature (between 320 and 420 °C [12]), oxidation and elimination of side groups (-CH₃ and -CH₂CH₂CF₃), yielding volatile products (e.g., $CH_2O$ and $C_6H_5F_3O$) and generating macroradicals that induce crosslinking. Second, at high temperature (above 420 °C [12]), partial thermal depolymerization occur [11,13]. The final product of this extreme degradation is pure silica [11].

Under isothermal conditions, intermediate behaviors are observed depending on aging temperatures. Xu et al. quantitatively investigated the competing mechanisms of oxidation and depolymerization through elemental analysis of a fluorosilicone homopolymer aged at 350 °C for 6 h under air [14]. Their results showed a 42.5 % reduction in the relative atomic percentage of silicon due to the volatilization of cyclic siloxanes, while the remaining 57.5 % participated in oxidation reactions. Their study demonstrated that below 350 °C, the predominant mechanism involves oxidative side-group cleavage followed by macroradical crosslinking, generating stable products that effectively hinder the depolymerization process. Other studies also show that side-group oxidation simultaneously reduces organic content and promotes the

<footnote><sup>a</sup> Laboratoire PIMM, Arts et Métiers Institute of Technology, CNRS, CNAM, 151 boulevard de l'Hopital, Paris 75013, France</footnote>

<footnote><sup>b</sup> ICGM, Univ Montpellier, CNRS, ENSCM, Montpellier, France</footnote>

<footnote>* Corresponding author.
E-mail address: emmanuel.richaud@ensam.eu (E. Richaud).</footnote>
formation of new ≡Si-O-Si≡ intermolecular linkages [15,16]. In our previous study, we demonstrated that during the aging of a commercial fluorosilicone rubber under air at temperatures ranging from 200 to 250 °C, both chain scission and crosslinking occur simultaneously, with the latter resulting mainly from oxidation reactions [17]. Later, the effect of structure and formulation was first investigated on the thermal stability of fluorosilicones by comparing various fluorosilicone compositions through non-isothermal and isothermal thermogravimetric analyses under nitrogen [8]. We verified that depolymerization mainly occurred at high temperature, but was accompanied by side reactions involving lateral groups scission. Here, the aim is to pursue those investigations by elucidating the interdependent effects of polymer chemistry, temperature, and oxygen concentration on degradation kinetics, and the balance of depolymerization over oxidation. Since this latter determines the competition between chain scissions and crosslinking, this study will additionally explore macromolecular changes depending on the degradation temperature under O₂, providing valuable insights into inhomogeneities during the crosslinking induced by thermal oxidation process.

## 2. Experimental

### 2.1. Materials

The materials are detailed in Table 1, and Fig. 1 shows the chemical structures of the studied polymers and crosslinking agents.

Model samples were synthesized by organocatalyzed anionic ring opening polymerization of hexamethylcyclotrisiloxane (D₃), 1,3,5-trifluoropropyl-1,3,5-trimethylcyclotrisiloxane (F₃) and 1,3,5-trivinyl-1,3,5-trimethylcyclotrisiloxane (V₃). The targeted proportions of each monomer were added to a Schlenk flask equipped with a mechanical

<caption>**Table 1**</caption>

Summary of materials under study. (M) is for model samples, (C) is for commercial samples, N is for neutralizer (CaCO₃), and F/V/D is the ratio of 3,3,3-trifluoropropylmethylsiloxane (F), dimethylsiloxane (D) and methylvinylsiloxane (V).

<table><thead><tr><th>Name</th><th>Description</th><th>Polymer molar ratio: F/D/V<sup>a</sup></th><th>Additional components</th><th>Average number molar mass M<sub>n</sub> (g.mol<sup>-1</sup>)<sup>b</sup> and polydispersity D<sup>b</sup></th></tr></thead><tbody><tr><td>HP</td><td>(M) Homopolymer</td><td>100/0/0</td><td>-</td><td>M<sub>n</sub> = 80,000<br>D = 3.0</td></tr><tr><td>CP1</td><td>(M) Copolymer</td><td>92/8/0</td><td>-</td><td>M<sub>n</sub> = 86,000<br>D = 1.8</td></tr><tr><td>CP2</td><td>(M) Copolymer</td><td>64/36/0</td><td>-</td><td>M<sub>n</sub> = 62,000<br>D = 1.7</td></tr><tr><td>TP1</td><td>(M) Terpolymer</td><td>90/8/2</td><td>-</td><td>M<sub>n</sub> = 84,000<br>D = 1.8</td></tr><tr><td>TP2</td><td>(M) Terpolymer</td><td>80/10/10</td><td>-</td><td>M<sub>n</sub> = 60,000<br>D = 1.9</td></tr><tr><td>XTP1+F</td><td>(M) Crosslinked and filled terpolymer</td><td>90/8/2</td><td>0.6 phr of DBPH + 25 phr of SiO<sub>2</sub></td><td>M<sub>n</sub> = 84,000<sup>c</sup><br>D = 1.8</td></tr><tr><td>TP3+F</td><td>(C) Filled terpolymer</td><td>85/13/2</td><td>0.6 phr of DBPH + 37 phr of SiO<sub>2</sub></td><td>M<sub>n</sub> = 300,000<br>D = 2.2</td></tr><tr><td>XTP3+F</td><td>(C) Crosslinked and filled terpolymer</td><td>85/13/2</td><td>0.6 phr of DBPH + 37 phr of SiO<sub>2</sub></td><td>M<sub>n</sub> = 300,000<sup>c</sup><br>D = 2.2</td></tr><tr><td>XTP3+F +N</td><td>(C) Crosslinked and filled terpolymer</td><td>85/13/2</td><td>1.3 phr of DCP + 37 phr of SiO<sub>2</sub></td><td>M<sub>n</sub> = 300,000<sup>c</sup><br>D = 2.2</td></tr></tbody></table>

<figure><img src="image_3.png" /><figcaption>**Fig. 1.** Chemical structures of the studied polymers and peroxides.</figcaption></figure>

stirrer under nitrogen flow. Then, 0.5 mL of a solution of diphenylsilanediol in THF (0.4 M) was added to the monomer mixture. Experimental [M]₀/[I]₀ was determined before the catalyst addition by ¹H NMR spectroscopy. 1,5,7-triazabicyclo[4.4.0]dec-5-ene (TBD) was solubilized in THF at 60 °C until a perfectly clear solution was obtained. 0.5 mL of a solution of TBD in THF (1.2 M) was then added to start the polymerization. The polymerization was carried out under nitrogen. The resulting polymers were precipitated twice from THF in methanol, washed three times with methanol and dried overnight at 100 °C under reduced pressure (6.5 × 10⁻² mbar). Due to the difference in monomer reactivities, terpolymers are expected to display an almost triblock structure with a fluorinated central block (containing few D or V units) and two external blocks rich in V and D units. Their ¹H NMR data are given in Fig. SI-1. More details of the synthesis process can be found in the studies of Corfa et al. [18].

The model crosslinked and filled sample (XTP1+F) contained 20 wt. % of fumed silica (Sigma-Aldrich, S5130). Commercial samples contained around 27 wt. % of inorganic particles, mostly fumed silica particles and a small percentage of titanium dioxide (TiO₂, less than 1 wt. %).

The crosslinked films (200 µm thick) were prepared by curing in a hot press at 170 °C during 4 min in the presence of peroxides. Two peroxides were used: 2,5-dimethyl-2,5-di-(tert-butylperoxy)hexane (DBPH) (purity ≥ 92 %), and dicumyl peroxide (DCP) containing 60 wt. % of CaCO₃.

The peroxide loadings were adjusted to the same molar amount (19 mmol kg⁻¹) but had different active oxygen concentrations: 2.28 mmol kg⁻¹ for DBPH and 1.14 mmol kg⁻¹ for DCP. The initial crosslinking density (ν) of post-cured samples was determined using swelling test (see **Supplementary Information SI-2**). For XTP3+F (crosslinked with DBPH), ν was 0.18 mol/kg, while for XTP3+F+N (crosslinked with DCP), ν was 0.12 mol/kg. Based on their initial active oxygen content, DCP demonstrated higher crosslinking efficiency compared to DBPH. However, there is no evidence to suggest that the chemical nature of the crosslink bridges differs between the two peroxides.

The overall strategy for elaborating model and commercial materials is given in Fig. 2.

### 2.2. Characterization

#### 2.2.1. Size-exclusion chromatography (SEC)

SEC was performed using a triple-detection GPC system (Agilent Technologies) with dedicated software for multidetector analysis. Two PL1113–6300 ResiPore columns (300 × 7.5 mm) were used with THF as the eluent at a flow rate of 1 mL/min. Detection was carried out with a 390-LC PL0390–0601 refractive index detector. The system was maintained at 35 °C. Calibration was done using PMMA standards ranging from 540 to 2210,000 g/mol. Samples were prepared by diluting 0.1 mL (1.8 M) of polymer in 1 mL of THF.

#### 2.2.2. NMR

NMR samples were prepared in CDCl₃ and analyzed on a Bruker Avance I 400 MHz spectrometer at 25 °C. ¹H NMR spectra were recorded with a spectral width of 8 kHz, transmitter offset of 3 kHz, acquisition time of 4 s, and four scans.

<footnote>a Determined by ¹H NMR.</footnote>

<footnote>b Determined by GPC in THF at 1 mL/min using PMMA narrow standards for column calibration.</footnote>

<footnote>c For crosslinked polymers, Mₙ corresponds to the value measured before peroxide crosslinking process.</footnote>
<figure><img src="image_4.png" /><figcaption>Fig. 2. Schematic representation of material synthesis and compounding.</figcaption></figure>

### 2.2.3. Thermogravimetric analyses (TGA)

TGA was performed under pure oxygen atmosphere (50 mL/min) using a Q500 (TA Instruments). Different oxygen concentrations were obtained with a Q50 coupled with a GasMix, mixing the desired amount of oxygen with nitrogen to achieve the same rate of gas flow. Temperature ramp analyses were carried out with a heating rate of 10 °C/min from 25 to 800 °C. The isotherm temperatures were reached with a temperature ramp of 20 °C/min under nitrogen, switching the gas afterwards in case of oxidizing atmosphere. In this case, the initial sample weight ($w_0$) was taken at the beginning of the isotherm. For the model unfilled samples, isotherms were conducted three times for each material, while the other materials and conditions were tested at least twice. In order to better estimate the degradation of rubber part in filled samples, raw data from isothermal treatments of filled samples were corrected by subtracting the filler content from the sample mass to better isolate polymer degradation by the following equation:

$$ \frac{m_{\text{rubber}}(t)}{m_{\text{rubber0}}} = \frac{m(t) - w_{f0} \cdot m_0}{m_0 \cdot (1 - w_{f0})} \quad (1) $$

Where $w_{f0}$ is the initial mass ratio of filler, $m(t)$ and $m_0$ correspond to residual and initial mass of filled sample, $m_{\text{rubber}}(t)$ and $m_{\text{rubber0}}$ correspond to residual and initial mass of the rubber. The maximum degradation kinetics were calculated after the initial 5 wt. % polymer mass loss, excluding the onset phase.

### 2.2.4. Modulated differential scanning calorimetry (MDSC)

Measurements using MDSC were conducted with a Q1000 series calorimeter from TA Instruments under nitrogen atmosphere, using standard aluminum pans. The samples were first equilibrated at -80 °C, followed by a 3-minute isotherm, and then heated at a rate of 1 °C/min up to 40 °C, with temperature modulation of ± 0.50 °C every 60 s. This technique enables the separation of the total heat flow signal versus temperature into two components: a reversible signal associated with thermal transitions such as the glass transition, and a kinetic signal associated with other phenomena such as molecular relaxation [19]. The instrument was calibrated using sapphire samples to ensure accurate calorimetric capacity values.

## 3. Results and discussion

### 3.1. Influence of oxygen concentration

To evaluate the impact of varying oxygen levels on degradation, TGA temperature ramps were conducted under different $O_2$ concentrations. Fig. 3 compares the homopolymer (HP) and the commercial material XTP3+F+N, demonstrating their contrasting behaviors as the two materials exhibiting the most divergent properties in our study. For HP, even a small amount of oxygen (5 %) significantly lowers the onset temperature of degradation. The 5 % weight loss temperature ($T_{5\%}$) shifts from 447 °C under $N_2$ to 392 °C under 5 % $O_2$. For XTP3+F+N, this shift was less pronounced, $T_{5\%}$ only decreased from 434 °C under $N_2$ to 430 °C under 5 % $O_2$. For both materials, this effect became more pronounced at higher oxygen concentrations. These lower onset temperatures highlight the fact that, in the presence of oxygen, reactions driven by oxygen attack are more likely to initiate degradation. Indeed, the activation energy of thermo-oxidation has been shown to be lower than that of depolymerization [13].

The progression of this degradation can vary, either further oxidizing the material or creating a pathway of degradation that also promotes depolymerization. An indicator of oxidation is the final plateau observed in temperature ramps. Under $N_2$, HP reached a zero-final plateau, indicating complete depolymerization of the polymer, while the residual weight of XTP3+F+N is close to its initial filler content (27 wt. %). However, higher oxygen concentrations resulted in higher plateau values for both materials, suggesting that a portion of the material oxidized, forming a stable residue. Thus, the polymer transforms into two distinct fractions: a volatile fraction and a residual solid fraction. As earlier described, the volatile fraction is composed of cyclic siloxanes formed by depolymerization, as well as oxidized side groups [11,13]. Several mechanisms may be involved, depolymerization leading to monomers $D_3$, $D_4$ [11], $F_3$, $F_4$ and possibly $F_xD_y$, and products coming from direct cleavage of lateral groups (such as 3,3,3-trifluoropropene revealed from Py-GC-MS - see Fig. SI-2), while the residual solid is primarily composed of $SiO_2$. FTIR spectra of XTP3+F+N residues after TGA temperature ramps under $O_2$ and $N_2$ are given in Supplementary Information Fig. SI-3 and is very close to the spectra of $SiO_2$ present in literature [20].

Fig. 4 presents our proposed degradation mechanism for the formation of $SiO_2$ after total oxidation. According to Thomas [21], ≡Si•
<figure><img src="image_5.png" /><figcaption>Fig. 3. TGA temperature ramp under various oxygen concentration with (a) HP, and (b) XTP3+F+N.</figcaption></figure>

radicals can be formed through oxidation of α-position -CH₂- groups, adjacent to Si atoms, generating hydroperoxides (-CH(OOH)-) that decompose and create ≡Si• radicals (Fig 4-a). At high temperature, another possibility to form ≡Si• radicals is through the direct side-group scission (Fig. 4-b) as suggested by a Py-GC-MS analysis of XTP3+F (see Fig. SI-2). According to Bond Dissociation Energies, ≡Si-CH₂CH₂CF₃ bond would be weak enough and actually could undergo a homolysis reaction (see Suppl. Info 4). The formation of ≡Si-O-Si≡ crosslinks can have two origins. First, the recombination of ≡Si• and OH• radicals could form silanols ≡Si-OH, followed by their condensation to create crosslinks (Fig. 4-c). However, this pathway requires sufficient radical stability for recombination, which remains speculative. Second, ≡Si• can react with O₂ to form peroxyl radicals ≡Si-OO•, which decompose to siloxyl radicals ≡Si-O• via a process analogous to hydrocarbon oxidation (Fig. 4-d) [22]. Finally, crosslinking arises through combination of ≡Si• and ≡Si-O• radicals or with silica surfaces, creating inactive products (Fig. 4-e). Through degradation pathways analogous to those shown in Fig. 4a-4e, side-chain oxidation ultimately yields SiO₂ (Fig. 4f).

In summary, the reaction process proposed in Eq. (2) accounts for the experimental results in Fig. 3.

$$ \text{Polymer } \frac{\Delta, O_2}{\beta Volatiles + \gamma Residue} \quad (2) $$

Based on this mechanism, and assumption that the residue is primarily composed of SiO₂ and does not contain carbides (as supported, for instance, by Supplementary Information Fig. SI-3), we can assess the extent of oxidation and depolymerization as a function of oxygen concentration. Considering only the atomic quantity of silicon (Si), the initial molar amount of Si is converted into the molar fraction of released cyclic siloxanes (β) and the molar fraction of oxidized polymer that forms SiO₂ (γ). Therefore, here, β represents the fraction of the material involved in depolymerization, neglecting other minor thermal decomposition pathways, and γ represents the fraction involved in oxidation, calculated as follows:

$$ \gamma = \left( \frac{w_{\text{SiO}_2} - w_{\text{f0}}}{w_{\text{polymer}}} \right) \frac{M_{\text{repeat unit}}}{M_{\text{SiO}_2}} \quad (3) $$

$$ \beta = 1 - \gamma \quad (4) $$

Where wSiO₂ (wt. %) is the final plateau at the end of TGA temperature ramp, wf0 is initial filler content (wt. %) when present, and wpolymer is the initial weight percentage of polymer matrix, excluding fillers. Mrepeat unit and MSiO₂ are the theoretical molecular weight of the monomers associated with each polymer and silica, respectively. For example, HP under 100 % O₂ reached a final plateau at 33 wt. % (wSiO₂ = 33 %, wf0 = 0 %, wpolymer = 100 %, MHP = 156 g/mol, MSiO₂ = 60 g/mol), yielding γ=0.86 and β=0.14, for the contribution of oxidation and

depolymerization, respectively.

Fig. 5 shows how the contribution of each mechanism changes with increasing oxygen concentration. For HP, up to 20 % O₂, the relationship is nearly linear, with each 1 % increase in oxygen reducing depolymerization by 3 %. In contrast, XTP3+F+N showed no oxidized residue at 5 % O₂, and between 5 and 20 % O₂, each 1 % increase in oxygen reduced depolymerization by 1.5 %. Additionally, the observed reduction in oxidation kinetic rate at higher oxygen levels is a phenomenon already observed by other studies on thermo-oxidation of thermoplastics [23,24], indicating a critical oxygen concentration beyond which the reactions become independent of further increases in O₂. Moreover, this critical oxygen concentration increases at lower temperatures [23,24]. Therefore, under 1 bar O₂ pressure, samples undergo full oxidation at higher degradation temperatures (corresponding to the onset temperature in Fig. 3). In this condition, the system can be considered to be in the “oxygen-excess regime”, meaning that further increasing the oxygen concentration no longer accelerates oxidation. From a kinetic perspective, this implies that the oxygen concentration is high enough to neglect reactions between two Si° radicals, for example. For polyolefins, it has been shown that this critical oxygen pressure increases as the temperature decreases [25]. However, there is no clear evidence that silicon-based polymers exhibit the same behaviour. Nonetheless, at moderate temperatures (e.g., below 220 °C, as discussed later), depolymerization is expected to become negligible, and degradation should be dominated by oxidation. In this regime, the overall degradation rate is expected to remain sensitive to oxygen pressure.

## 3.2. Influence of temperature

The extent of oxidation is also influenced by temperature, as demonstrated in Fig. 6 showing multiple isothermal analyses of the homopolymer (HP) and the commercial rubber (XTP3+F+N) under pure O₂. This finding reveals that by decreasing temperature, there is a higher final residue weight, indicating enhanced crosslinking mechanism. The oxidation of lateral groups, generating ≡Si-CH•-R and ≡Si• radicals, dominates over direct scissions of ≡Si-CH₂CH₂CF₃ groups due to its lower activation energy. The stable final residue contains primarily new ≡Si-O-Si≡ bridges, and minor ≡Si-CH₂-Si≡ and ≡Si-CH₂-O-Si≡ crosslinks. Therefore, the observed final plateau levels also indicate that oxidation dominates over depolymerization at lower temperatures, consistent with its lower activation energy [13]. This process forms a stable network, with increased crosslinking and reduced mobility, thereby inhibiting further degradation reactions under the same conditions.
<figure><img src="image_6.png" /></figure>

<figure><img src="image_7.png" /></figure>

<figure><img src="image_8.png" /></figure>

<figure><img src="image_9.png" /></figure>

<figure><img src="image_10.png" /></figure>

<figure><img src="image_11.png" /><figcaption>(f)</figcaption></figure>

<caption>**Fig. 4.** Oxidation of the fluorinated side group (a), release a volatile product (b), formation of silanol and subsequent crosslinking (c), proposed formation of siloxyl (d), radicals' recombination (e), and proposed mechanism of SiO₂ formation (f).</caption>

### 3.3. Effect of structure and formulation on thermo-oxidative stability

To further emphasize the impact of chemical structure on the thermo-oxidative stability, the analysis presented in Section 3.1 was extended to other materials at 20 % oxygen concentration (similar to air condition). Fig. 7 illustrates this comparison, highlighting the susceptibility of each material to its dominant degradation mechanism and

estimating the relative contributions of depolymerization and oxidation in the overall degradation process.

This comparative analysis shows that materials with a higher degree of depolymerization tend to exhibit greater thermo-oxidative stability. Basing our analysis on the residual weight at the end of the temperature ramp test, it seems to us that when depolymerization dominates, little residue remains. In contrast, oxidation produces more residue because it
<figure><img src="image_12.png" /><figcaption>Fig. 5. Representative contribution of each mechanism over increase in oxygen concentration with (a) HP, and (b) XTP3+F+N for samples heated at 10 °C min⁻¹ under pure oxygen up to total decomposition of polymer.</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>Fig. 6. Weight loss of HP (a) and XTP3+F+N considering the weight loss of the polymer matrix (normalized by filler ratio) (b) in isothermal condition at different temperatures under pure O₂.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>Fig. 7. Relative contribution of each mechanism from TGA temperature ramp at 20 % of oxygen for different materials.</figcaption></figure>

promotes crosslinking, which limits depolymerization and generates a stable polymer network. Since oxidation has a lower activation energy than depolymerization (30 vs. 43 kcal mol⁻¹ for PDMS [13]), it is expected to occur more easily at lower temperature and leave higher

residual weights. Hence, **assuming that oxidation and depolymerization are two competing processes**, the following interpretation can be proposed:

(i) if depolymerization strongly predominates over thermal oxidation irrespectively to temperature, the final residue would tend towards 0.

(ii) Reversely, basing on the previous observation that low temperature oxidation favors crosslinking and crosslinking usually slows down depolymerization, the final of the plateau will be increased.

Here, our result show that the contributions from each mechanism may vary depending on polymer structure and formulation. Considering that, at 20 % O₂, the homopolymer (HP) exhibits greater oxidation susceptibility than to its D-units-containing analog (CP1), this difference cannot be attributed to oxygen permeability, as both silicone samples show high permeability, with pure PDMS being slightly more permeable than pure PTFPMS [26]. Instead, the results demonstrate that copolymerization enhances oxidative stability, supporting the theoretical expectation that PTFPMS is more prone to oxidation than PDMS due to the higher reactivity of -CH₂- groups compared to -CH₃ groups [21, 27]. On the other hand, a higher degree of depolymerization does not imply faster reaction kinetics. Previous studies conducted under an inert atmosphere have shown that copolymers depolymerize more slowly than fluorosilicone homopolymers. This highlights the interest of this analysis in minimizing the influence of oxidation and favor the pathway
that involves higher activation energy. Additionally, compared to the
unfilled polymers, the filled rubbers show an even lower tendency to
oxidation, indicating that the inorganic particles also have an antioxi-
dant effect.

Under 100 % O₂, maximizing the oxidative effects, Fig. 8 shows the TGA temperature ramps for the materials, indicating that the temperature required to reach a 10 wt. % weight loss under O₂ ranged from 326 °C (TP2) to 380 °C (XTP3+F), whereas under N₂, the temperature varied from 437 °C (TP2) to 481 °C (XTP3+F+N), as showed in our previous work [8]. The final plateau levels also varied, reflecting the impact of chemical structures on their oxidation.

Although temperature ramps are a straightforward and effective method for highlighting differences between materials, they represent a more extreme condition of degradation. In contrast, isothermal analyses provide a closer representation of the conditions under which the materials are used. For example, aerospace industry standards use a testing temperature of 200 °C [28,29]. Accelerated tests, such as isothermal measurements at higher temperatures, are also valuable for building kinetic models and estimating material lifetimes, although results should be interpreted with caution when comparing across different temperature ranges. In this study, we use TGA under isothermal conditions as an accelerated test, allowing multiple materials to be compared in a simpler way.

Therefore, to investigate and compare the thermo-oxidative stability of the materials, Fig. 9 presents isothermal analyses under O₂ at 300 °C for unfilled model samples and at 320 °C for filled samples. We consider that depolymerization plays a minor role in these conditions, as evidenced by the limited weight loss (1.7 % after 120 min at 320 °C under N₂) shown for sample XTP3+F+N in **Supplementary Information Fig. SI-4**. Consequently, our analysis primarily reflects oxidative degradation effects. The higher temperature for filled samples was selected because the commercial samples exhibited minimal weight change at 300 °C during the analyzed period, indicating their enhanced thermo-oxidative stability. Table 2 details the time required to reach a 10 wt. % weight loss at both 300 °C and 320 °C, as well as the maximum degradation rate for each sample at these temperatures.

Among the model samples, we can consider that HP and CP1 exhibit similar kinetic behaviors at 300 °C. Thus, only a significant amount of b-units enhanced the resistance to thermo-oxidation, as observed for CP2, but not for CP1. TP1 and TP2 showed slow degradation kinetics at 300 °C, potentially due to the oxidation of vinyl groups and recombination of macroradicals being favored.

XTP1+F demonstrates very good thermo-oxidative stability compared to the other model samples, indicating a potential anti-oxidative effect of the silica particles. Fumed silica is mainly

incorporated into silicones both to facilitate processing, given their low molecular weight, and to provide reinforcement, thereby enhancing the tensile strength of the final product. SiO₂ could help stabilize the material by slowing down oxygen diffusion. However, we consider that the oxidation of thin fluorosilicone films (200 µm) is not governed by diffusion, as partly confirmed by the results in Fig. SI-5. Since the results of oxygen permeability measurements for thin samples (200 µm) aged at 250 °C under air presented no significant difference over time (see Table SI-1), we will neglect the role of oxygen diffusion (and the possible effect of SiO₂ on it) as an explanation of our results in the following of this paper.

The ability of SiO₂ to scavenge hydroxyl radicals (HO•) has also been documented in the scientific literature [30,31]. Rusevova Crincoli et al. reports that the decomposition of rhodamine B dye (RhB) in the presence of H₂O₂ (decomposing into HO• and HOO• radicals) proceeds more slowly when SiO₂ is present [29]. In this work, we assume that silanol groups on the silica surface can react with HO• radicals, forming siloxyl radicals ≡Si-O• and H₂O [32]. Such radicals would probably react with others organic radicals (as ≡Si• radicals formed through the pathway illustrated in Fig. 4-a) or other surface particles, creating stable products, as proposed in Fig. 10. As explored in other references [33,34], the antiradical activity of SiO₂ could be confirmed by direct observations (using ESR, for example), although the lifetime of radicals may be very short under the ageing conditions of the present study. Additionally, quantifying trace amounts of water could be considered, provided that the reactions shown in Fig 10 generate a significant amount of water compared with other reactions involved in FVMQ degradation (such as the reaction of HO• radicals with other abstractable C—H sites).

The commercial samples showed a substantial difference compared to the model samples, with 320 °C being a more suitable temperature for comparing filled materials. The isotherms of the commercial samples presented an initial plateau that was absent in the model filled rubber (XTP1+F), suggesting that this effect was not related to silica particles.

As showed in **Supplementary Information Fig. SI-6**, non-post-cured samples had a longer initial plateau period compared to post-cured ones, and samples crosslinked with higher concentrations of peroxide exhibited a shorter plateau. This suggests that another component in the initial formulation may have an antioxidative role, influencing both the crosslinking process and the early stages of degradation.

A similar behavior was seen by Xu et al. during isothermal tests on FSR, where the addition of CeO₂ functioned as a radical scavenger [27]. Given the presence of a small amount of TiO₂ in the commercial samples, this might similarly account for the observed differences, as Englert et al. showed that even 0.5 wt. % of TiO₂ could enhance the thermo-oxidative stability of PDMS rubber [3]. This improvement occurs through a redox reaction where TiO₂ captures free electrons, forming Ti³⁺ ions, preventing the formation of HO• radicals [31]. Moreover, his highlights the substantial influence of inorganic particles, even at low concentrations, on the thermo-oxidative stability of fluorosilicones, surpassing the effects of changes in chemical structure.

Therefore, considering the maximum kinetic rate of degradation ($k_{320}$) rather than $t_{10%}$ to compare the commercial samples, TP3+F showed lower $k_{320}$ than XTP3+F, likely due to crosslinking involving the available vinyl groups, consistent with observations in other terpolymers. XTP3+F+N exhibits a lower $k_{320}$ than XTP3+F. The stabilizing effect of CaCO₃ was first observed under inert conditions [8] and explained based on previous research [35]. In the case of ageing under air, two complementary explanations can be proposed: either this effect is caused by residual vinyl groups participating in crosslinking reactions during testing, or to the chemical composition of XTP3+F+N crosslinks being less susceptible to oxidation. In TP3+F, allylic CH₂ groups are located at the vicinity of silicon atoms, making them more sterically hindered and therefore less reactive than the central methylene groups connected through the crosslink nodes ≡Si-CH₂-CH₂-CH₂-Si≡.

<figure><img src="image_15.png" /><figcaption>**Fig. 8.** TGA temperature ramp under pure O₂ with all the studied materials.</figcaption></figure>
<figure><img src="image_16.png" /><figcaption>**Fig. 9.** TGA under pure O₂ in isotherms (a) at 300 °C showing model unfilled samples, and (b) at 320 °C with filled materials and HP, considering the weight loss of the polymer matrix (normalized by filler ratio).</figcaption></figure>

**Table 2**

Time to reach 10 wt. % weight loss (t₁₀%) at 300 °C and 320 °C, and maximum apparent kinetic rate (k) at same temperatures. NB: mass loss rates were normalized using filler ratio.

<table><thead><tr><th>Material</th><th>t<sub>10%</sub> at 300 °C (min)</th><th>k<sub>300</sub> x 10 (wt. %/min)</th></tr></thead><tbody><tr><td>HP</td><td>11 ± 1</td><td>-9 ± 1</td></tr><tr><td>CP1</td><td>9 ± 0.2</td><td>-12 ± 0.1</td></tr><tr><td>CP2</td><td>15 ± 2</td><td>-7 ± 1</td></tr><tr><td>TP1</td><td>17 ± 4</td><td>-6 ± 2</td></tr><tr><td>TP2</td><td>18 ± 10</td><td>-5 ± 2</td></tr><tr><td>XTP1+F</td><td>33</td><td>-3</td></tr><tr><td>Material</td><td>t<sub>10%</sub> at 320 °C (min)</td><td>k<sub>320</sub> (wt. %/min)</td></tr><tr><td>XTP1+F</td><td>9</td><td>-1.3</td></tr><tr><td>TP3+F</td><td>32</td><td>-0.7</td></tr><tr><td>XTP3+F</td><td>39</td><td>-1.8</td></tr><tr><td>XTP3+F +N</td><td>26</td><td>-1.6</td></tr></tbody></table>

**3.4. Macromolecular consequences of chemical degradation**

To further investigate the impact of temperature on polymer mobility, we performed modulated differential scanning calorimetry (MDSC). Fig. 11 illustrates the reverse heat capacity (Cₚ) in the temperature range for the glass transition temperature (Tg) of the materials, comparing the effect of temperatures at 250 and 300 °C under O₂. At 250 °C, there is a noticeable shift in Tg to higher temperatures, a widening of the temperature range, and a decrease in ΔCₚ. These observations are characteristic of a typical crosslinking process that increases network rigidity and possibly affects the distribution of crosslink nodes. However, at 300 °C, we observe almost no shift in Tg and a significant decrease in ΔCₚ. Table 3 summarizes the MDSC results, including conditions at 220 °C and 280 °C. These analyses indicate different macro-molecular evolution depending on the temperature range.

According to Logeais [36], the observed changes in the thermograms can be explained as follows: oxidation induces additional crosslinking and leads to the formation of “mineralized” regions (composed of pure SiO₂ or possibly Si-O-C zones). At 250 °C, these regions remain at low

<figure><img src="image_17.png" /><figcaption>**Fig. 10.** Mechanism of stabilization by SiO₂.</figcaption></figure>

<figure><img src="image_18.png" /><figcaption>**Fig. 11.** Reverse Cₚ over temperature from MDSC, showing the glass transition evolution of HP under pure O₂ at (a) 250 °C and (b) 300 °C.</figcaption></figure>
<table><caption>**Table 3**
Summary of MDSC data at different temperature analyzing thermo-oxidative HP degradation.</caption><thead><tr><th>Aging temperature (°C)</th><th>Aging time (min)</th><th>Residual weight (%)</th><th>T<sub>g</sub> (°C)</th><th>ΔC<sub>p</sub></th><th>Ratio ΔC<sub>p</sub></th></tr></thead><tbody><tr><td rowspan="3">220</td><td>0</td><td>0</td><td>-70</td><td>0.24</td><td>1</td></tr><tr><td>600</td><td>91</td><td>-61</td><td>0.19</td><td>0.78</td></tr><tr><td>1200</td><td>83</td><td>-50</td><td>0.17</td><td>0.71</td></tr><tr><td rowspan="3">250</td><td>45</td><td>93</td><td>-63</td><td>0.19</td><td>0.84</td></tr><tr><td>90</td><td>88</td><td>-56</td><td>0.11</td><td>0.50</td></tr><tr><td>180</td><td>80</td><td>-43</td><td>0.068</td><td>0.30</td></tr><tr><td rowspan="3">280</td><td>30</td><td>84</td><td>-68</td><td>0.076</td><td>0.28</td></tr><tr><td>60</td><td>73</td><td>-67</td><td>0.048</td><td>0.18</td></tr><tr><td>90</td><td>69</td><td>-68</td><td>0.039</td><td>0.16</td></tr><tr><td rowspan="3">300</td><td>15</td><td>85</td><td>-69</td><td>0.094</td><td>0.42</td></tr><tr><td>30</td><td>69</td><td>-69</td><td>0.033</td><td>0.15</td></tr><tr><td>45</td><td>66</td><td>-69</td><td>0.038</td><td>0.11</td></tr></tbody></table>

concentration, as the oxidation rate is moderate and occurs throughout the entire sample volume. In this regime, a single-phase network is formed, and the distribution of crosslinking nodes gradually broadens over time, as indicated by the widening of the $T_g$ range. Basing on a hypothesis inspired by classical literature on heterogeneous oxidation [37], we believe that as the temperature increases, ageing becomes progressively more heterogeneous, and the formation of “mineralized” regions becomes predominant over random crosslinking. The material evolves into a two-phase heterogeneous network, consisting of a highly crosslinked phase, while the other remains a relatively mobile domain, as illustrated in Fig. 12. As a result, $T_g$ remains nearly constant.

The presence of a rigid phase can have a similar effect as adding silica particles to a siloxane matrix, which reduces overall mobility, lowering $ΔC_p$, without significantly affecting $T_g$ [38–41]. On the other hand, even though few additional crosslinks have been formed in the mobile domain, this might have minimal impact on $T_g$, since the high flexibility of siloxane polymer chains requires a substantial number of new crosslinks to noticeably shift $T_g$ [42]. Furthermore, if this mobile fraction undergoes depolymerization in a subsequent stage, this process also does not significantly influence $T_g$, as observed for HP aged at 400 °C under $N_2$, presented in **Supplementary Information Fig. SI-7**.

## 4. Conclusions

The analysis of degradation under different oxygen concentrations revealed that even small amounts of oxygen have a notable impact on the degradation process of model HP, while having a smaller impact on the commercial material XTP3+F+N. Below 20% of $O_2$ for HP, each 1% increase in oxygen concentration led to a 3% reduction in depolymerization. In contrast, XTP3+F+N decreased 1.5% in depolymerization per 1% increase in oxygen concentration. The temperature ramp method demonstrated that the susceptibility to oxidation and depolymerization is dependent on material composition, with filled samples showing a stronger tendency toward depolymerization. Oxidized units may accelerate depolymerization rather than forming a stable oxidized network. A chemical mechanism was proposed, assuming that part of ≡Si• radicals may react with oxygen and generate $SiO_2$. An open question is whether this mechanistic scheme can accurately simulate ageing kinetics under different oxygen pressures, as has been done for polyolefin oxidation [21].

This study also explored how other chemical modifications and formulations impact the thermo-oxidative degradation of fluorosilicones under isothermal conditions. HP and CP1 displayed similar degradation kinetics, but HP was more susceptible to oxidation, while CP1 (F/D = 92/8) tended toward greater depolymerization. Only high D-unit content in CP2 (F/D = 64/36) effectively mitigated degradation. TP1 and TP2 exhibited slower degradation rates, likely due to the oxidation of vinyl groups and macroradical recombination. Commercial samples differed significantly from model samples, with inorganic particles, especially

<figure><img src="image_19.png" /><figcaption>**Fig. 12.** Schematic evolution of the macromolecular structure of HP aged under $O_2$ (I) up to 250 °C, and (II) above 280 °C.</figcaption></figure>

$SiO_2$, significantly enhancing thermo-oxidative stability in the early stages of degradation. Additionally, using DPC as a crosslinking agent instead of DBPH slowed down the degradation kinetics of FSR. The crosslinking junctions formed by DPC might be less susceptible to oxidation, but further investigations are needed to confirm this.

Temperature also plays a critical role, with lower temperatures significantly favoring oxidation over depolymerization. At a macromolecular level, differences in behavior depending on the temperature ranges were noted. Below 250 °C, a single-phase network forms, leading to an increase in crosslinks and their distribution over time. In contrast, above 280 °C, a two-phase heterogeneous network develops, consisting of a highly crosslinked phase and a more mobile domain that is also susceptible to depolymerization.

### CRediT authorship contribution statement

**Leticia Pires:** Writing – original draft, Methodology, Investigation, Formal analysis. **Alice Corfa:** Resources, Investigation. **Vincent Ladmiral:** Writing – review & editing, Supervision, Conceptualization.

**Sylvain Caillol:** Writing – review & editing, Supervision, Methodology.

**Sébastien Roland:** Writing – review & editing, Supervision, Methodology. **Emmanuel Richaud:** Writing – review & editing, Supervision, Methodology, Funding acquisition, Conceptualization.

### Declaration of competing interest

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

### Acknowledgements

We gratefully acknowledge the Agence Nationale de la Recherche (ANR) for their financial support through the FLuorosilicones Ageing Mechanisms and Synthesis (FLAMS ANR-20-CE06-0024) project. We also extend our thanks to ITC Elastomères for providing the commercial materials used in this research (XTP3+F and XTP3+F+N).

### Supplementary materials

Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.polymdegradstab.2025.111772.
# Data availability

Data will be made available on request.

# References

[1] J.G. Drobny, Fluoroelastomers, in: introduction to fluoropolymers, Elsevier (2013) 149-230, https://doi.org/10.1016/B978-1-4557-7442-5.00010-3.

[2] R. Han, Y. Li, Q. Zhu, K. Niu, Research on the preparation and thermal stability of silicone rubber composites: a review, Compos. C: Open Access 8 (2022) 100249, https://doi.org/10.1016/j.jcomc.2022.100249.

[3] M. Englert, F. Minister, A. Moussaoui, W. Pisula, Mechanical properties of thermo-oxidative aged silicone rubber thermally stabilized by titanium oxide based fillers, Polym. Test. 115 (2022) 107726, https://doi.org/10.1016/j.polymertesting.2022.107726.

[4] A. Zheng, Y. Huang, Y. You, J. Hu, D. Wei, X. Xu, Y. Guan, Boron particles acting as antioxidants for fluorosilicone rubber due to their radical scavenging activity, Polym. Degrad. Stab. 158 (2018) 168-175, https://doi.org/10.1016/j.polymdegradstab.2018.09.017.

[5] L.C. Sim, S.R. Ramanan, H. Ismail, K.N. Seetharamu, T.J. Goh, Thermal characterization of Al2O3 and ZnO reinforced silicone rubber as thermal pads for heat dissipation purposes, Thermochim. Acta 430 (2005) 155-165, https://doi.org/10.1016/j.tca.2004.12.024.

[6] Z. Xu, Y. Zhang, J. Zhou, M. Qi, J. Shi, J. Zhang, Study on high-temperature composite properties of fluorosilicone rubber with NANO-SB 2 O 3, J Appl. Polym. Sci. 137 (2020) 49302, https://doi.org/10.1002/app.49302.

[7] Y. Guan, R. Yang, Y. Huang, C. Yu, X. Li, D. Wei, X. Xu, Multi-walled carbon nanotubes acting as antioxidant for fluorosilicone rubber, Polym. Degrad. Stab. 156 (2018) 161-169, https://doi.org/10.1016/j.polymdegradstab.2018.06.018.

[8] L. Pires, A. Corfa, V. Lademiral, S. Caillol, S. Roland, E. Richaud, Relationship between structure and thermal stability of fluorosilicones at high temperatures, Polym. Degrad. Stab. 238 (2025) 111340, https://doi.org/10.1016/j.polymdegradstab.2025.111340.

[9] G. Camino, S.M. Lomakin, M. Lazzari, Polydimethylsiloxane thermal degradation part 1, Kinet. Asp. Polym. 42 (2001) 2395-2402, https://doi.org/10.1016/S0032-3861(00)00652-2.

[10] R.J. Pazur, A. Tilley, B. Zaman, C. Porter, Thermal degradation of fluorosilicone rubber up to 250 °C, Rubber Chem. Technol. 98 (2025) 308-328, https://doi.org/10.5254/RCT.24.00052.

[11] R.G. Jones, eds, in: W. Ando, J. Chojnowski (Eds.), Silicon-Containing Polymers: The Science and Technology of Their Synthesis and Applications, Springer Netherlands, Dordrecht, 2004, https://doi.org/10.1007/978-94-011-3939-7.

[12] T.H. Thomas, T.C. Kendrick, Thermal analysis of polydimethylsiloxanes. I. Thermal degradation in controlled atmospheres, J. Polym. Sci, A-2 Polym, Phys. 7 (1969) 537-549, https://doi.org/10.1002/pol.1969.160070308.

[13] T.H. Thomas, T.C. Kendrick, Thermal analysis of polydimethylsiloxanes. I. Thermal degradation in controlled atmospheres, J. Polym. Sci, A-2 Polym, Phys. 7 (1969) 537-549, https://doi.org/10.1002/pol.1969.160070308.

[14] X. Xu, J. Liu, P. Chen, D. Wei, Y. Guan, X. Lu, H. Xiao, The effect of ceria nanoparticles on improving heat resistant properties of fluorosilicone rubber, J. Appl. Polym. Sci. 133 (2016) 44117, https://doi.org/10.1002/app.44117 app.

[15] J. Wu, J. Dong, Y. Wang, B.K. Gond, Thermal oxidation ageing effects on silicone rubber sealing performance, Polym, Degrad, Stab. 135 (2017) 43-53, https://doi.org/10.1016/j.polymdegradstab.2016.11.017.

[16] T. Kaneko, S. Ito, T. Minakawa, N. Hirai, Y. Ohki, Degradation mechanisms of silicone rubber under different aging conditions, Polym, Degrad, Stab. 168 (2019) 108936, https://doi.org/10.1016/j.polymdegradstab.2019.108936.

[17] L. Pires, A. Corfa, V. Lademiral, S. Caillol, M. Dekar, C. Beauvais, S. Roland, E. Richaud, Thermal and thermo-oxidative accelerated aging of fluorosilicone rubber, Polym, Degrad, Stab 232 (2025) 111150, https://doi.org/10.1016/j.polymdegradstab.2024.111150.

[18] A. Corfa, Fluorosilicones: De La Polymérisation à La Dépolymérisation Pour Le Recyclage, Université de Montpellier, 2024. PhD Thesis, https://theses.fr/s368534.

[19] A. Toda, Kinetics of enthalpy recovery studied by temperature-modulated differential scanning calorimetry, Thermochim. Acta 716 (2022) 179330, https://doi.org/10.1016/j.tca.2022.179330.

[20] T.N. Tran, T.V. Anh Pham, M.L. Phung Le, T.P. Thoa Nguyen, V.M. Tran, Synthesis of amorphous silica and sulfonic acid functionalized silica used as reinforced phase for polymer electrolyte membrane, Adv. Nat. Sci.: Nanosci, Nanotechnol. 4 (2013) 045007, https://doi.org/10.1088/2043-6262/4/4/045007.

[21] D.K. Thomas, High temperature stability in fluorosilicone vulcanisates, Polymer (Guildf) 13 (1972) 479-484, https://doi.org/10.1016/0032-3861(72)90087-0.

[22] D. Lindsay, J.A. Howard, E.C. Horswill, L. Iton, K.U. Ingold, T. Cobbley, A. Li, The bimolecular self-reactions of secondary peroxy radicals. Product studies, Can, J. Chem. 51 (1973) 870-880, https://doi.org/10.1139/v73-131.

[23] E. Richaud, F. Farcas, P. Bartoloméo, B. Fayolle, L. Audouin, J. Verdu, Effect of oxygen pressure on the oxidation kinetics of unstabilised polypropylene, Polym. Degrad. Stab. 91 (2006) 398-405, https://doi.org/10.1016/j.polymdegradstab.2005.04.043.

[24] O. Okamba-Diogo, E. Richaud, J. Verdu, F. Fernagut, J. Guilment, B. Fayolle, Molecular and macromolecular structure changes in polyamide 11 during thermal oxidation, Polym. Degrad, Stab. 108 (2014) 123-132, https://doi.org/10.1016/j.polymdegradstab.2014.05.028.

[25] E. Richaud, Durabilité Des Géotextiles en Polypropylène, Arts et Métiers ParisTech, 2006, https://pastel.hal.science/pastel-00002053.

[26] W.L. Robb, Thin silicone membranes-their permeation properties and some applications, Ann, N. Y. Acad, Sci. 146 (1968) 119-137, https://doi.org/10.1111/j.1749-6632.1968.tb20277.x.

[27] X. Xu, J. Liu, P. Chen, D. Wei, Y. Guan, X. Lu, H. Xiao, The effect of ceria nanoparticles on improving heat resistant properties of fluorosilicone rubber, J. Appl. Polym. Sci. 133 (2016), https://doi.org/10.1002/app.44117 app.44117.

[28] AFNOR, Série aérospatiale - Caoutchoucs fluorosilicones - Classe 61 - caractéristiques. https://www.boutique.afnor.org/fr-fr/norme/nf1l7261/serie-aerospatiale-caoutchoucs-fluorosilicones-classe-61-caractéristiques/fa192253/80084, 2018.

[29] AFNOR, Série aérospatiale - élastomères fluorosiliconés - Classe 63 - Caractéristiques. https://www.boutique.afnor.org/fr-fr/norme/nf1l7163 /serie-aerospatiale-élastomères-fluorosilicones-classe-63-caractéristiques/fa191 203/79889, 2017.

[30] K. Rusevova Crincoli, S.G. Huling, Hydroxyl radical scavenging by solid mineral surfaces in oxidative treatment systems: rate constants and implications, Water. Res. 169 (2020) 115240, https://doi.org/10.1016/j.watres.2019.115240.

[31] M. Suh, P.S. Bagus, S. Pak, M.P. Rosynek, J.H. Lunsford, Reactions of hydroxyl radicals on titania, silica, alumina, and gold surfaces, J. Phys. Chem. B 104 (2000) 2736-2742, https://doi.org/10.1021/jp993653e.

[32] S. Ferraris, I. Corazzari, F. Turci, A. Cochis, L. Rimondini, E. Vernè, Antioxidant activity of silica-based bioactive glasses, ACS Biomater. Sci, Eng. 7 (2021) 2309-2316, https://doi.org/10.1021/acsbiomaterials.1c00048.

[33] M. Kawakubo, K. Tsunoda, H. Yajima, T. Ishii, H. Kaidou, A. Ahagon, Reactions of radicals in filled rubber compounds: I. Detection of free radicals in rubbers by ESR, Rubber Chem. Technol. 78 (2005) 644-658, https://doi.org/10.5254/1.3547903.

[34] E. Troncoso-Ortega, P. Castano-Rivera, R. Romero, A. Henriquez, C. Méndez, E. Schott, D. Contreras, Organic radicals stabilization in natural rubber: discerning the influence of thermo-oxidation using chemically modified or unmodified lignin as antioxidant, J. Ind. Eng. Chem. 137 (2024) 435-447, https://doi.org/10.1016/j.jiec.2024.03.029.

[35] S. Hamdani, C. Longuet, J.-M. Lopez-Cuesta, F. Ganachaud, Calcium and aluminium-based fillers as flame-retardant additives in silicone matrices. I. Blend preparation and thermal properties, Polym. Degrad. Stab. 95 (2010) 1911-1919, https://doi.org/10.1016/j.polymdegradstab.2010.04.013.

[36] C. Logeais, Influence De La Synthèse Et Du Vieillissement Sur Les Propriétés Mécaniques Des Fluorosilicones, Université Paris sciences et lettres, 2024. https://pastel.hal.science/tel-05106850.

[37] B. Fayolle, L. Audouin, G.A. George, J. Verdu, Macroscopic heterogeneity in stabilized polypropylene thermal oxidation, Polym. Degrad. Stab. 77 (2002) 515-522, https://doi.org/10.1016/S0141-3910(02)00110-6.

[38] P. Klonos, A. Panagopoulou, L. Bokobza, A. Kyritsis, V. Peoglos, P. Pissis, Comparative studies on effects of silica and titania nanoparticles on crystallization and complex segmental dynamics in poly(dimethylsiloxane), Polymer (Guildf) 51 (2010) 5490-5499, https://doi.org/10.1016/j.polymer.2010.09.054.

[39] D. Fragiadakis, P. Pissis, Glass transition and segmental dynamics in poly(dimethylsiloxane)/silica nanocomposites studied by various techniques, J. Non. Cryst. Solids. 353 (2007) 4344-4352, https://doi.org/10.1016/j.jnoncrysol.2007.05.183.

[40] D. Fragiadakis, P. Pissis, L. Bokobza, Glass transition and molecular dynamics in poly(dimethylsiloxane)/silica nanocomposites, Polymer (Guildf) 46 (2005) 6001-6008, https://doi.org/10.1016/j.polymer.2005.05.080.

[41] G. Tsagaropoulos, A. Eisenberg, Dynamic mechanical study of the factors affecting the two glass transition behavior of filled polymers. Similarities and differences with random ionomers, Macromolecules 28 (1995) 6067-6077, https://doi.org/10.1021/ma00122a011.

[42] S. Fanse, Q. Bao, Y. Zou, Y. Wang, D.J. Burgess, Effect of crosslinking on the physicochemical properties of polydimethylsiloxane-based levonorgestrel intrauterine systems, Int. J. Pharm. 609 (2021) 121192, https://doi.org/10.1016/j.ijpharm.2021.121192.