# Relationship between structure and thermal stability of fluorosilicones at high temperatures

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. Relationship between structure and thermal stability of fluorosilicones at high temperatures. Polymer Degradation and Stability, 2025, 238, pp.111340. 10.1016/j.polymdegradstab.2025.111340. hal-05032827

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Submitted on 14 Apr 2025

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# Relationship between structure and thermal stability of fluorosilicones at high temperatures

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

## ARTICLE INFO

**Keywords:**

Fluorosilicone
Thermal stability
TGA
Degradation

## ABSTRACT

The thermal stability of fluorosilicones at high temperatures has been investigated with thermogravimetric analysis (TGA) based on their chemical structure (copolymerization and crosslinking), and compounding (incorporation of silica and titanium dioxide particle fillers). This allowed the comparison of thermal degradation kinetics of model samples and commercial fluorosilicones rubbers. Increasing dimethylsiloxane content in copolymers enhances thermal stability, while the presence of vinyl groups accelerates degradation by promoting depolymerization and crosslinking reactions at an earlier stage. In model filled fluorosilicone rubbers (FSR), untreated silica particles were found to destabilize the polymer matrix in the short term, whereas commercial filled FSRs exhibited greater thermal stability due to treated silica surfaces. Additionally, commercial samples crosslinked with DCP mixed with CaCO₃ resulted in the highest thermal stability. FTIR transmission analysis of copolymer residues revealed that degradation does not favor the preferential release of any monomer (dimethylsiloxane or trifluoropropylmethylsiloxane). Elemental analysis of a model homopolymer suggested that depolymerization is the dominant degradation mechanism, but side group scissions also contribute to this process.

## 1. Introduction

Fluorosilicone rubbers (FSR) are a class of polysiloxane elastomers mainly used as sealing products in contact with non-polar fluids and subjected to extreme temperature variations [1]. FSRs primarily consist of the 3,3,3-(trifluoropropyl)methylsiloxane unit in their chemical structure. However, commercial FSRs are formulated with a more complex chemical composition to achieve optimal properties for their applications.

Common modifications to the chemical structure of FSR include: (1) incorporating dimethylsiloxane units to reduce the high cost of fluorosilicone products and potentially balance the benefits between the two polymers [2]; (2) incorporating a small number of methylvinylsiloxane moieties into the fluorosilicone structure to enhance the efficiency of peroxide crosslinking and create the rubber network structure [3,4]; (3) adding silica particles (SiO₂) as reinforcing fillers due to their good compatibility with the silicon-oxygen backbone and high thermal stability [5]. Additional minor components, such as antioxidants or

pigments, can also be present. Inorganic nano-additives are commonly used to also address the requirement for high-temperature resistance [6]. This modified material exhibits different properties compared to the basic homopolymer commonly described in the literature. The impact of each added component on thermal stability is not well documented. This knowledge is important to explain changes over time and develop strategically enhanced material formulations for improved thermal stability.

The literature describes that under inert atmosphere and high temperatures, polysiloxanes primarily degrade through depolymerization [7]. The pronounced polarity of the Si-O linkage and flexibility of the -(Si-O)ₓ segments facilitate inter and intramolecular redistribution reactions, leading to the release of cyclic and volatile siloxane groups, which are thermodynamically favored compared to linear chains [8].

This degradation process can be significantly influenced by the chemical nature of the chain end-groups. Hydroxyl-terminated groups, such as silanol (Si-OH), accelerate degradation via a "back-biting" mechanism, promoting intramolecular rearrangement reactions with

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

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

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: emmanuel.richaud@ensam.eu (E. Richaud).</footnote>
the end of the chains. In contrast, end-capped polysiloxanes, such as those with trimethylsiloxy terminal groups (-O-Si(CH₃)₃), primarily depolymerize through random rearrangements within the polymer chain. As a result, end-capped polysiloxanes exhibit higher thermal stability compared to their hydroxyl-terminated counterparts [9]. Moreover, for end-capped polysiloxanes, the molecular weight of the sample is less critical than for hydroxyl-terminated polysiloxanes and does not significantly impact the onset of depolymerization, as supported by Thomas and Kendrick, who demonstrated that the depolymerization activation energy for polydimethylsiloxane (PDMS) samples with molecular weights ranging from 9600 to 4 × 10⁶ g mol⁻¹ remained nearly constant [10].

Regarding depolymerization of polysiloxane copolymers, Grassie and Francey reported that during the depolymerization of a poly (dimethyl-co-methylphenylsiloxane), dimethylsiloxane units are preferentially and randomly eliminated from the polymer chain through intra and intermolecular reorganization, leading to an enrichment of methylphenyl units in the remaining residue [11]. They also noted degradation through intermolecular processes and the release of mixed cycles containing different side groups from monomers of different chemical composition. To the best of our knowledge, the released products of a fluorosilicone copolymer are not documented in the literature.

In addition to volatile siloxanes, some authors have identified volatile fluorinated groups arising from the thermal scission of lateral groups. Thomas and Kendrick reported that 25% of the volatile products from an isotherm at 400 °C under vacuum were due to the cleavage of the Si–C bond, producing difluoropropylene (C₃H₄F₂) [12]. You et al. identified trifluoropropene (C₃H₃F₃) only at 600 °C using Py-GC-MS, with 6.4% of C₃H₃F₃ detected from their pure methyl-terminated polymer. This value increased to 12.7% with the addition of 10 wt. % silica particles [9]. Indulekha et al. studied a vinyl-terminated PTFPMS filled with cristobalite silica by TG-MS (10 °C/min) and found that degradation began at 340 °C, with C₃H₄F₂ as the major decomposition

product [13].

The influence of silica particles on thermal stability is debated in the literature. Some studies suggest that SiO₂ enhances thermal stability by forming physical interactions with the polymer matrix, reducing polymer mobility, and limiting the formation of degradation products [14, 15]. However, others propose that silanol groups on the surface of the particles can promote reactions, leading to main chain scissions, where one part of the chain grafts onto the silica while the other is released as a hydroxyl-terminated polysiloxane, thereby accelerating depolymerization [16,17]. Another inorganic particle present in peroxide compounds, CaCO₃, was described by Thomas as a neutralizer for basic polymerization residues, with the released CO₂ providing an acidic environment that improves stability, a benefit he also attributed to the acidic nature of silica [18]. Conversely, Hamdani et al. demonstrated that in the absence of basic residues, the addition of CaCO₃ to silicone rubber did not influence thermal stability below 600 °C [19].

Therefore, given that thermal degradation mechanisms of polysiloxanes are primarily based on PDMS literature, the impact of replacing methyl groups with trifluoropropyl groups in fluorosilicones is significant but underexplored. Most studies have focused on fluorosilicone homopolymers or complex FSR formulations, leaving gaps in understanding the specific thermal degradation mechanisms of fluorosilicones. To address this, thermogravimetric analysis (TGA) was employed to compare various model samples and commercial rubbers, examining how different chemical structures and formulations, such as copolymerization and silica particle inclusion, influence their thermal stability. All model polymers were synthesized with Si-(CH₃)₃ end-groups to avoid the back-biting mechanism. Molecular weight differences among the studied samples were not considered a significant factor affecting thermal stability. Additionally, chemical differences were analyzed using elemental analysis for a model homopolymer and FTIR-Transmission for a fluorosilicone copolymer. These analyses aimed to complement the gaps in the literature regarding the scission of side groups and the depolymerization in copolymer cases.

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

Material specifications. (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) units.

<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><br>(g.mol<sup>-1</sup>)<sup>b</sup>/ Polydispersity index D</th></tr></thead><tbody><tr><td>HP</td><td>(M) Homopolymer</td><td>100/0/0</td><td>–</td><td>80,000 / 3</td></tr><tr><td>CP1</td><td>(M) Copolymer</td><td>92/8/0</td><td>–</td><td>86,000 / 1.8</td></tr><tr><td>CP2</td><td>(M) Copolymer</td><td>64/36/0</td><td>–</td><td>62,000 / 1.7</td></tr><tr><td>TP1</td><td>(M) Terpolymer</td><td>90/8/2</td><td>–</td><td>84,000 / 1.8</td></tr><tr><td>TP2</td><td>(M) Terpolymer</td><td>80/10/10</td><td>–</td><td>60,000 / 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>84,000 / 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>300,000 / 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>300,000 / 2.2</td></tr><tr><td>XXTP3+F</td><td>(C) Crosslinked and filled terpolymer</td><td>85/13/2</td><td>0.75 phr of DBPH + 37 phr of SiO<sub>2</sub></td><td>300,000 / 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>300,000 / 2.2</td></tr><tr><td>XXTP3+F+N</td><td>(C) Crosslinked and filled terpolymer</td><td>85/13/2</td><td>1.625 phr of DCP + 37 phr of SiO<sub>2</sub></td><td>300,000 / 2.2</td></tr></tbody></table>

<figure><img src="image_4.png" /><figcaption>**Fig. 1.** Illustrative chemical structures of the studied polymers and peroxides.</figcaption></figure>

<footnote>a Determined by ¹H NMR. b Determined by GPC in THF at 1 mL.min⁻¹ using PMMA narrow standard for column calibration for uncrosslinked rubber.</footnote>
<figure><img src="image_5.png" /></figure>

<figure><img src="image_6.png" /><figcaption>**Fig. 2.** Schematic representation of material synthesis and compounding.</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>**Fig. 3.** TGA temperature ramp under N₂ for model (M) and commercial (C) materials (see Table 1).</figcaption></figure>

## 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 correct proportions of each monomer were added to a Schlenk flask equipped with a mechanical stirrer under nitrogen flow. Then, 0.5 mL of a solution of diphenylsilsanediol 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. Triazabicyclodecene (1,5,7-triazabicyclo[4.4.0]dec-5ene, 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

<figure><img src="image_8.png" /><figcaption>**Fig. 4.** (a) Weight loss, and (b) logarithmic weight loss over time in isotherm at 400 °C under N₂ with model samples synthesized in this work (see Table 1).</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>**Fig. 5.** TGA under N₂ in isotherms at 350 °C (a and b) and at 400 °C (c and d), showing unfilled model samples (left) and filled samples (right). Data were corrected for taking into account the filler weight ratio (Eq. (1)).</figcaption></figure>

<caption>**Table 2**</caption>

Time to reach 5 wt. % weight loss (t₅%) at 350 °C, time to reach 10 wt. % weight loss (t₁₀%) at 350 °C, and maximum kinetic rate constant (k) at 350 °C.

<table><thead><tr><th>Material</th><th>t<sub>5%</sub> at 350 °C (min)</th><th>t<sub>10%</sub> at 350 °C (min)</th><th>k x 10<sup>2</sup> (wt. % min<sup>-1</sup>)</th></tr></thead><tbody><tr><td>HP</td><td>76</td><td>116</td><td>-13.0</td></tr><tr><td>CP1</td><td>120</td><td>194</td><td>-9.6</td></tr><tr><td>CP2</td><td>166</td><td>306</td><td>-3.5</td></tr><tr><td>TP1</td><td>81</td><td>156</td><td>-5.3</td></tr><tr><td>TP2</td><td>36</td><td>117</td><td>-6.8</td></tr><tr><td>XTP1+F</td><td>62</td><td>170</td><td>-4.9</td></tr><tr><td>TP3+F</td><td>219</td><td>593</td><td>-1.6</td></tr><tr><td>XTP3+F</td><td>224</td><td>571</td><td>-1.7</td></tr><tr><td>XTP3+F+N</td><td>269</td><td>764</td><td>-1.3</td></tr></tbody></table>

°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 and ²⁹Si NMR data are given in Appendix A. More details of the synthesis process can be found in the studies of Corfa [20].

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 samples 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 non-post-cured samples was determined using a swelling test (see Appendix A). For XTP3+F (crosslinked with DBPH), ν was 0.13 mol kg⁻¹, while for XTP3+F+N (crosslinked with DCP), ν was 0.08 mol kg⁻¹. 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. Thermogravimetric analyses (TGA)

TGA was performed under pure nitrogen atmosphere (50 mL min⁻¹) using a Q50 (TA Instruments). Tests of temperature ramp were carried out with a heating rate of 10 °C min⁻¹ from 25 to 800 °C maximum. The isotherm temperatures were reached with a temperature ramp of 20 °C min⁻¹. In this case, the initial sample mass (w₀) was taken at the beginning of the isotherm. 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:

<caption>**Table 3**</caption>

Temperature of 10 wt. % weight loss (T₁₀ %) for end-capped PDMS and PTFPMS (HP in Table 1) evaluated by TGA temperature ramp under inert condition at a heating rate of 10 °C/min.

<table><thead><tr><th colspan="2">PDMS</th><th colspan="2">PTFPMS</th></tr><tr><th>T<sub>10</sub> % (°C)</th><th>Description</th><th>T<sub>10</sub> % (°C)</th><th>Description</th></tr></thead><tbody><tr><td>432–433 [7,22–24]</td><td>Pure polymer</td><td>415–426 [9,25]</td><td>Pure polymer</td></tr><tr><td>533 [26]</td><td>Filled (40 phr of fumed silica)</td><td>450 [27]</td><td>Filled (35 phr of fumed silica)</td></tr><tr><td>454 [28]</td><td>Filled (30 phr of fumed silica)</td><td>394 [29]</td><td>Filled (35 phr of hydrophilic silica)</td></tr></tbody></table>
<figure><img src="image_10.png" /><figcaption>Fig. 6. Chain scission and grafting mechanism on the silica surface [33].</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>Fig. 7. TGA under N₂ during isotherms at 400 °C, showing filled samples and considering the weight loss of the polymer matrix.</figcaption></figure>

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

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.

### 2.2.2. Elemental analyses

Carbon (C) and hydrogen (H) measurements were conducted using a homemade Carbon/Hydrogen elemental micro-analyzer (Institut des Sciences Analytiques, UMR5280 CNRS, Villeurbanne). Approximately 1.5 mg of sample was weighed in silver cups and combusted in a flow of 95 mL/min of pure oxygen. The combustion system, equipped with two furnaces at 1050 °C and 850 °C, contained a vertical glass tube filled with CuO oxidation catalyst and silver wool to trap halogens. Carbon and hydrogen were converted into CO₂ and H₂O, respectively, and measured using a CO₂/H₂O non-dispersive infrared detector (Rosemount NGA 2000). Calibration was performed with 99.9 % pure standards and two reference standards were measured every eight analyses to ensure calibration stability. Measurement uncertainties were ± 0.20 for H and ± 0.30 for C. Standards used for calibration and stability controls, purchased from Sigma Aldrich Corp. (Saint-Louis, Missouri, USA), included Methionine, Glycine, Taurine, Cystine, Valine, Tryptophan, O-Phospho-DL-Serine, L-Ascorbic Acid, and Uracil, selected to match the expected %C and %H of the samples.

The analyzed samples were the virgin and thermally degraded homopolymer (HP). Degradation was carried out using the Q50 (TA Instruments) in order to reach a weight loss of approximately 10, 20, and 30 wt. % in isothermal condition at 350 and 400 °C under N₂.

### 2.2.3. Fourier-transform infrared spectroscopy (FTIR)

The FTIR analyses in transmission mode were performed with a Perkin Elmer infrared spectrometer and data treatment was done with the Spectrum software. Each spectrum was the result of 16 accumulated scans, measured with a resolution of 2 cm⁻¹, between 400 and 4000 cm⁻¹. Potassium bromide (KBr, Sigma-Aldrich 221,864) was used to prepare KBr pellets with a concentration of 1 wt. % for each sample. The measured absorbances (A) were normalized by the thickness (l) of the corresponding KBr pellet in order to quantify the chemical functions, according to Beer-Lambert law: A = ε × l × C, where ε is the molar extinction coefficient (absorption coefficient) and C is the concentration of the absorbing species.

## 3. Results and discussion

### 3.1. Comparison of thermal stability between the materials

TGA temperature ramps were conducted to determine the onset temperatures of polymer degradation and evaluate the residual products, as shown in Fig. 3. The results reveal that all model samples without fillers reached a zero plateau at the end, indicating complete degradation through depolymerization. In contrast, the filled samples presented a final plateau corresponding to the initial percentage of inorganic particles in their formulation. Comparing the temperature required to achieve a 10 wt. % weight loss, the samples ranged from 437 °C (TP2) to 481 °C (XTP3+F+N).

However, this method was not sufficient to distinguish the thermal stability of the materials, particularly among the model samples, since they exhibited similar thermodynamic behavior at the onset of degradation. Therefore, isothermal analyses were conducted to compare the degradation kinetics of the polymers and better identify differences in their thermal stability. By performing an isothermal analysis at 400 °C on model samples and considering the typical exponential behavior of polymer degradation at high temperatures, Fig. 4 shows that ln(m/m₀) does not follow a linear relationship with time (t). Typically, a linear behavior would suggest that degradation follows a first-order reaction, governed by a single global kinetic rate constant (k), as described by Eq. (2) [21], leading to a linear decay and total weight loss.

$$ \frac{d\left(\frac{m}{m_0}\right)}{dt} = -k\left(\frac{m}{m_0}\right) \quad (2) $$

Fig. 3b reveals a clear non-linear trend with auto-acceleration. The deviation from this first-order behavior indicates that the samples undergo multiple reactions and degradation stages, each with distinct kinetic rates. The early stages of aging can involve the formation of highly reactive species, while the residue observed at longer aging times points to the creation of a highly stable product that resists depolymerization. Moreover, this result clearly demonstrates a significant difference between unfilled and filled materials, with the latter showing much slower degradation kinetic.

To further highlight the differences between the materials, Fig. 5 shows isotherms at 350 °C and 400 °C, for unfilled and filled materials.
<figure><img src="image_12.png" /><figcaption>**Fig. 8.** (a) FTIR spectra of CP2 residue over time at 400 °C under N₂, (b) normalized absorbance ratios of bands at 1267 cm⁻¹ normalized to 1211 cm⁻¹, and at 1017 cm⁻¹ normalized to 1211 cm⁻¹ over time.</figcaption></figure>

Since the lower temperature provides a clearer criterion for differentiation, Table 2 details the time required to reach 5 % and 10 % weight loss, as well as the maximum degradation rate for each sample at 350 °C.

First, when comparing the homopolymer with the copolymers, we observe that an increase in dimethylsiloxane (D) content enhances thermal stability by delaying the onset of degradation and slowing down the process. Experimental data from other authors support that pure methyl-terminated polydimethylsiloxane (PDMS) is more thermally stable than pure methyl-terminated poly[(3,3,3-trifluoropropyl)methylsiloxane] (PTFPMS, also named HP in Table 1), as shown in Table 3. This finding contradicts the theory applied to aromatic groups, where large lateral groups with electron-donating effects are known to enhance thermal stability of polysiloxanes [11]. We propose that these factors alone are insufficient to account for the experimental behavior observed in fluorinated silicones. A more plausible explanation considers that the polymer with these larger substituents is less thermodynamically stable than cyclic oligomers, favoring depolymerization [16]. This instability can be attributed to increased steric interactions within the polymer

structure, where repulsive forces are significantly amplified [17].

Secondly, when comparing CP1 with TP1, the introduction of 2 mol. % of vinyl groups negatively impacted the onset of degradation, reducing the time to reach a 5 % mass loss but lowering the mass loss rate at longer time (see for example the schematization in Fig. 5a). Comparing TP1 and TP2, we see that a higher vinyl content further reduces this onset time.

Since vinyl groups are primarily located near the chain ends (see Section 2.1), they can initially participate in depolymerization, which could account for the rapid mass loss observed at early aging stages. However, alternative reaction pathways could occur if vinyl groups are involved in crosslinking. Assuming that depolymerization proceeds more slowly in crosslinked samples than in uncrosslinked ones, this would result in a lower kinetic rate constant compared to CP1.

To understand the crosslinked samples, it is important to note that the crosslinking process reduces the number of vinyl groups by reacting with peroxide to form crosslinking junctions, with most of these junctions believed to consist of trimethylene bridges (—Si—(CH₂)₃—Si—)
<figure><img src="image_13.png" /><figcaption>Fig. 9. Carbon (C) and hydrogen (H) contents in homopolymer HP (see Table 1) over weight loss at 350 and 400 °C under N₂.</figcaption></figure>

[30]. The quantity of crosslinking junctions is significantly influenced by the amount of vinyl groups when vinyl-specific peroxide, like DBPH and DCP, are used [4,31]. However, Thomas demonstrated that fluoro-silicones are more susceptible to peroxide attack than PDMS, resulting in a larger number of crosslinks than expected, even in rubbers without vinyl groups [32]. This suggests the possibility of simultaneous crosslinking reactions, leading to a more complex network structure. Nevertheless, residual vinyl groups can remain after the crosslinking process depending on its efficiency.

Examining the model filled crosslinked polymer (XTP1+F), we observe that the onset of degradation (t₅%) occurs earlier than that of uncrosslinked unfilled (TP3+F). The reaction between siloxane chains and silanol groups on the silica surface possibly contributes to the initial degradation stage, leading to the release of hydroxyl-terminated chains that accelerate depolymerization via backbiting mechanism, as illustrated in Fig. 6. However, over longer periods, XTP1+F exhibits slower degradation kinetics, indicating that the formed silanol groups have a limited influence. These groups participate in additional reactions, such as condensation, adsorption or reaction with the silica surface that create covalent bonds between the filler and the polymer chain, acting as new crosslink nodes and slowing down the depolymerization process.

When examining the commercial samples (XTP3+F and XTP3+F+N), we observe greater thermal stability compared to the model XTP1+F. This improvement is likely due to the presence of silica particles that have undergone at least moderate surface treatment, which reduces the number of silanol groups on their surface. With a less active silica surface, the primary influence of the particles shifts to restricting polymer mobility, thereby reducing depolymerization and enhancing thermal stability [16,34]. Additionally, moderate surface modification is beneficial for industrial applications, as it lowers the viscosity of filled suspensions, simplifying processing [35].

When comparing the uncured commercial stock (TP3+F) to its crosslinked counterpart (XTP3+F) using pure peroxide (DBPH), XTP3+F exhibits greater thermal stability at 400 °C pointing out the positive effect of crosslinking on thermal stability. Fig. 7 shows that increasing the concentration of DBPH slows down thermal degradation, indicating that a higher quantity of crosslinking junctions for this material improves thermal stability. XTP3+F+N, which is crosslinked with DCP-CaCO₃, demonstrates the highest thermal stability among all the tested materials. This improved stability is not due to a higher crosslinking density, as XTP3+F+N had a lower initial crosslinking density. Therefore, we believe that this result points out that CaCO₃ and TiO₂ contribute positively to the enhanced thermal stability of this material. Both particles can neutralize HF [19,36], as indicated by the

mechanisms below:

$$CaCO_3 + 2 HF \rightarrow CaF_2 + CO_2 + H_2O$$

$$TiO_2 + 2 HF \rightarrow TiOF_2 + H_2O$$

This strong acidic species has the potential to attack Si-O bonds in the polymer backbone, generating ≡Si-OH and ≡Si-F (this species was proposed by Thomas and Kendrick [12]). By a chain scission mechanism, accelerating depolymerization by the following reaction:

$$≡Si-O-Si≡ + HF \rightarrow ≡Si-OH + ≡Si-F$$

Therefore, since both XTP3+F and XTP3+F+N contain TiO₂, the key difference lies in the presence of CaCO₃ in the latter formulation. This demonstrates the neutralization effect of CaCO₃, even under a nitrogen atmosphere.

Coming back to the effect of crosslinking, Fig. 7 shows that uncrosslinked TP3+F degrades faster than its crosslinked analogues. However, it also shows that rubber crosslinked with a higher concentration of DCP (XXTP3+F+N) does not exhibit the beneficial impact seen with XTP3+F. This suggests that differences in chemical composition and/or in the distribution of crosslinking junctions of these materials lead to different effects on thermal stability.

## 3.2. Investigations on the thermal degradation mechanism

To determine if a fluorosilicone copolymer, with F and D-units, preferentially releases a specific monomer during depolymerization, FTIR transmission analysis was performed on CP2 over time at 400 °C under N₂ (Fig. 8a) by mixing residue from TGA analysis with KBr. It seemed to us better to consider the ratio between two absorbance bands, expressing the relative residual concentrations of corresponding groups in aged samples. Fig. 8b shows two absorbance ratios: the bands at 1267 cm⁻¹ (C–H) relative to 1211 cm⁻¹ (C–F), denoted as A(1267)/A(1211), and the band at 1017 cm⁻¹ (Si–O–Si) relative to 1211 cm⁻¹ (C–F), denoted as “A(1017)/A(1211)”.

The results indicate that up to 150 min of degradation, there is no significant change to suggest a preference for the release of a specific monomer. If such a preference existed, the A(1267)/A(1211) ratio would increase with the elimination of F-units or decrease with the elimination of D-units. Therefore, depolymerization in this case randomly eliminates individual cyclic siloxane units, such as F₄ and D₃, along with mixed structures like F₂D, F₃D, and F₂D₂, as illustrated in Fig. 10a. Larger cyclic structures can also form in smaller quantities, as reported in the literature [9,37]. Additionally, although FTIR suggests that depolymerization is the primary degradation mechanism, with no change in the chemical nature of the aged samples, the presence of a persistent residue after prolonged isothermal treatment (Fig. 5c) and a degradation trend that deviates from first-order kinetics for all model materials (Fig. 4b) suggest that secondary reactions also occur. Therefore, the increase in the A(1017)/A(1211) ratio, observed only at later stages of degradation (180 min), can indicate the cleavage of side groups. This may be followed by the formation of new crosslinks through radical recombination, leading to a higher concentration of siloxane bands (Si–O–Si) in the structure.

To explore the possibility of side group scission during degradation, elemental analyses were conducted on HP, the model homopolymer without crosslinks or fillers. Carbon (C) and hydrogen (H) contents were tracked to determine whether loss of side groups actually occurred (decreasing C and H) or whether only depolymerization contributed to the observed weight loss (maintaining the same ratio of C and H in the polymer residue). Fig. 9 shows the percentage values of these elements during degradation at 350 and 400 °C under N₂.

The observed reduction in carbon (C) and hydrogen (H) indicates the presence of side group scission, while their small variations suggests that depolymerization remained as the predominant degradation
<figure><img src="image_14.png" /><figcaption>**Fig. 10.** Thermal degradation mechanisms: (a) intermolecular depolymerization of copolymers (F/D), (b) scission of side groups, and (c) release of main volatile compounds.</figcaption></figure>

mechanism. These scission processes have been reported by other authors, and recently studied by Wang et al. [38] as illustrated in Fig. 10b. According to the literature, this process can release volatile byproducts such as methane ($CH_4$) [37–41], trifluoropropene ($C_3H_3F_3$) [9] (also identified by a Py-GC-MS analysis – see Supplementary Information – Fig. SI-3), and difluoropropene ($C_3H_4F_2$), [12,13] (Fig. 10c).

As previously said, a residue persisted by the end of the isothermal treatments. For instance, after 50 h at 350 °C, HP left a residue of 7 wt. %, with an FTIR spectrum closely resembling that of the original sample

(see Supplementary Information – Fig. SI-4). However, the sample underwent a color change, from transparent to brown and eventually black in highly degraded samples, indicating a chemical modification (see Supplementary Information – Fig. SI-5). Some authors have associated the formation of black residue under inert conditions to silicon-oxycarbide (SiOC), generated from an alternative degradation path, probably involving side group scission as an intermediate step [42,43].

This residue points to the occurrence of crosslinking processes and to the presence of very short chains that are unable to depolymerize at the
same temperature. Crosslinking reactions likely arise from side chain scission, leading to the coupling of reactive species such as $≡Si°$, $≡Si-O°$, and $≡Si-CH_2°$, generating $≡Si-Si≡$, $≡Si-O-Si≡$, and $≡Si-CH_2-Si≡$. The release of methane, in particular, has been associated with crosslinking, leading to the formation of $Si-CH_2-Si$ linkages [39], but also with a complex array of scission reactions, resulting in the production of cyclic siloxanes and short-chain fragments [38]. These concurrent reactions, including depolymerization and crosslinking, could account for the non-linear behavior of the weight loss curves, as different reactions occur simultaneously.

## 4. Conclusions

The investigation into the thermal degradation mechanisms of fluorosilicones revealed several key insights. First, increasing the dimethylsiloxane concentration in copolymers enhanced thermal stability of fluorosilicones. The addition of vinyl groups promoted an earlier onset of degradation and slower kinetic rate constant, which might be correlated to depolymerization and crosslinking reactions driven by scissions of side groups. The model filled rubber revealed that untreated silica particles have a destabilizing effect in the short term. In contrast, the commercial samples showed greater thermal stability, likely due to treated silica surfaces that minimize the number of silanol groups, restrict polymer mobility, and inhibit depolymerization. Among these, XTP3+F+N showed the highest thermal stability due to the presence of $CaCO_3$. These findings highlight the critical roles of structural modifications, filler content, and crosslinking density in determining the thermal stability of fluorosilicones.

FTIR analyses of the copolymer residue during degradation indicated that depolymerization does not favor the release of any specific monomer, but instead leads to the release of both individual and mixed cyclic siloxanes. In the later stages of degradation, the scission of side groups contributes to a modification in the copolymer chemistry. Elemental analysis of the model homopolymer supported this possibility, as evidenced by the slight reduction in carbon (C) and hydrogen (H) content, although depolymerization remains the dominant mechanism.

### CRediT authorship contribution statement

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

The author is an Editorial Board Member/Editor-in-Chief/Associate Editor/Guest Editor for this journal and was not involved in the editorial review or the decision to publish this article.

### 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.2025.111340.

## Data availability

Data will be made available on request.

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