# Formation of skin-core structure in filled room-temperature vulcanized polydimethylsiloxane after long exposure to moderate temperatures

Florian Feyne, Éric Le Bourhis, Florian Lacroix, Laurence Autissier, Julie Pepin, Olga Smerdova

► To cite this version:

Florian Feyne, Éric Le Bourhis, Florian Lacroix, Laurence Autissier, Julie Pepin, et al. Formation of skin-core structure in filled room-temperature vulcanized polydimethylsiloxane after long exposure to moderate temperatures. Polymer Testing, 2025, 153, pp.109022. 10.1016/j.polymertesting.2025.109022. hal-05363555

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https://hal.science/hal-05363555v1

Submitted on 13 Nov 2025

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# Formation of skin-core structure in filled room-temperature vulcanized polydimethylsiloxane after long exposure to moderate temperatures

Florian Feyne<sup>a,†</sup>, Eric Le Bourhis<sup>a</sup>, Florian Lacroix<sup>b</sup>, Laurence Autissier<sup>c</sup>, Julie Pepin<sup>b</sup>, Olga Smerdova<sup>a,†</sup>

## ABSTRACT

This study highlights the formation of a skin-core effect inside a filled silicone rubber exposed to air for durations up to 600 days at temperatures between 30 °C and 70 °C. Coupling mechanical and physico-chemical characterisations reveal that the volatile components, such as crosslinking by-products and oligomers, evaporate from the elastomeric network during the exposure, resulting in the shrinkage of the sample. This phenomenon gives rise to the formation of a composition gradient from the sample's surface towards its core. All the samples become more brittle with ageing, but the tensile modulus increase only slightly at the longest durations and highest temperatures. More fillers are found on the outer surface while the infrared spectra in the core of the sample remain unaffected by neither the temperature nor the duration. The nanoindentation highlights gradients of elastic modulus on the cross-section and provides an estimation of the skin thickness. By wrapping the samples in aluminum foil, the evaporation of volatile components is prevented. The measured properties remain constant with the duration of thermal exposure even at the highest temperature.

## 1. Introduction

Silicone elastomers are widely used in a broad range of industries, including automotive, electronics, healthcare, aerospace, and mould making, among others [1]. Their good mechanical properties make them particularly well suited for encapsulation and sealing applications. While silicone elastomers include a large range of polymers, poly-dimethylsiloxane (PDMS) elastomers are most commonly used for research and industry purposes. PDMS elastomers can be filled using various fillers to tailor specific mechanical properties, while their chemistry can be changed by adjusting the crosslinker and the stoichiometry of the components.

Most of PDMS materials are cured under high temperatures, thus called HTV (high-temperature vulcanised) PDMS. However, some of them, RTV (room-temperature vulcanised) PDMS, can be vulcanised at room temperature after establishing the contact between the crosslinker, the catalyst and the PDMS chains. One of the main interests of using RTV PDMS in the industry is their pourability after mixing. Due to this property, they are particularly suitable for sealing applications with complex geometries. The RTV PDMS gaskets can be formed before sealing and comply with the joint parts by mechanical action, or be

injected directly in the space to fill, thus adhering to both parts during crosslinking. In these materials, the activation energy for starting the crosslinking process is usually brought by air humidity, oxidizing the catalytic specie. The crosslinking process can follow a polyaddition or polycondensation reaction, latter one producing a residual specie from the condensation reaction. However, RTV PDMS can also vulcanise by thermal activation, which reduces the crosslinking duration [1].

The ageing and thermal degradation mechanisms depend largely on the chemical composition [2] of the PDMS elastomers and are still not well understood for filled RTV PDMS elastomers. Although initial mechanical properties are lower for RTV PDMS comparing to HTV PDMS, both type of materials exhibit some decrease in elongation at break and tensile strength, together with an increase in hardness and modulus at 100 % extension after natural ageing for up to 20 years [3]. The PDMS elastomers are known for their thermal and oxidative resistance, attributed to the high dissociation energy of the Si-O bond in the Si-O-Si backbone [4,5]. The thermal degradation of these PDMS elastomers is well-studied at high temperatures and is attributed to successive oxidation phenomena [6–9]. Depending on the crosslinker molecular formula, on the terminal groups of the PDMS chains or on the ageing conditions, the degradation of the chemical structure and mechanical

<footnote><sup>a</sup> Institut Pprime, CNRS, ISAE-ENSMA, Université de Poitiers, Futuroscope Chasseneuil, 86962, France</footnote>

<footnote><sup>b</sup> Laboratoire de mécanique Gabriel Lamé (LaMé), EA7494, Université de Tours, Université d'Orléans, INSA Centre Val de Loire, 7 avenue Marcel Dassault, Tours, F-37200, France</footnote>

<footnote><sup>c</sup> CEA, DAM, Le Ripault, Monts, F-37260, France</footnote>

<footnote>* Corresponding author.
E-mail address: olga.smerdova@ensma.fr (O. Smerdova).</footnote>
properties can appear in a large range of temperatures starting from 120 °C [9]. The oxidation reaction induces chain-scission, sometimes accompanied by cross-linking [10] with different temperature-dependent kinetics [11] and possibly non-Arrhenius behaviour [12]. For instance, the thermal ageing of phenyl methyl vinyl silicone rubber in air results in cross-linking induced hardening and embrittlement of the material [13]. However, the RTV PDMS are usually not used at temperatures as high as the HTV PDMS. The condensation-cured silicone rubbers, for instance, entrap some volatile by-products and become subject to reversion when exposed to the temperature above 90 °C, which limits their temperature of application [14]. Moreover, they exhibit permanent shrinkage even at room temperature, which is supposed to stabilise after several days after cross-linking. However, very few studies have addressed the evolution of mechanical and physical properties of filled RTV PDMS bulk samples for a long period of time at the temperatures below the reversion.

A recent study by Brounstein et al. [15] of thermal ageing on commercial unfilled RTV PDMS elastomer Sylgard 184, conducted between 25 °C and 90 °C over durations ranging from 2 to 12 months, revealed an increase in the crosslinking density, associated to a post-cure of the material, accompanied by an increase in Shore A hardness. The properties stabilised after several months of ageing. An investigation by Patel and Skinner [16] examined an RTV 5370 siloxane rubber pads aged between 35 °C and 80 °C for up to 6 months. Thermo mechanical analysis (TMA) was also performed on this material for the temperatures up to 215 °C. The authors found that the stress at 25 % compression per unit density decreased significantly while sealed in moist inert atmosphere, but only slightly in open to air atmosphere. The compression set measured by TMA increased with temperature and time. The proposed degradation mechanism involves the volatilisation of catalytic species and water. However, when this volatilisation process is inhibited, the network degrades through a softening process, characterized by an increased contribution of low molecular weight species. The predominance of hydrolysis at high temperature ageing (>155 °C) in moist environment was also highlighted by Kaneko et al. [13] on phenyl methyl vinyl silicone rubber. Similar experimental tendencies regarding the increase in compression set and acceleration of the degradation by humidity were obtained by Patel et al. [17] on the foams of the Dow Corning RTV 5370 and by Labouriau et al. [18] on RTV 5370 and SX358 PDMS foams aged under compressive strain at temperatures from 30 to 110 °C up to 6 years. The suggested degradation mechanisms, however, are different. In the former study, the change in compression set is attributed to the viscoelastic reorganisation of the network without significant changes of crosslink density, as evidenced by NMR measurements. In the latter study, the degradation mechanism is associated with chemical activity of the tin(II) catalyst and its residues. The presence of 24 wt% of fumed silica fillers does not impact oxidizing behaviour of the RTV 5370 foam, as shown by Chaudhry and Billingham [19]. Using chemiluminescence results, authors suggested that the oxidation behaviour of this material is linked to the tin catalyst of its solvent.

It needs to be noted that the increase in compression set with ageing of silicone rubber impacts the sealing performance of the material by reducing the contact pressure, as demonstrated by Wu et al. [20] using coupled material study and finite element simulations of the contact with rough surface of the seal. A particular attention should thus be payed to the change of mechanical response of a rubber seal with ageing.

In addition to the above-mentioned phenomena, the ageing of thick polymer samples can introduce gradients in the cross-section related to gas diffusion, both inwards and outwards the bulk sample. It is well-known that the degradation of polymers in oxidative atmosphere can be linked to a diffusion-limited oxidation (DLO) phenomenon [21–23]. To characterise these gradients, spatially resolved techniques, such as infrared micro-spectrometer or a modulus profiler apparatus have been used by Celina et al. [22] and Gillen et al. [21,23]. Although, the oxidation gradients are not expected for RTV PDMS due to its low

temperature of use, this approach seems relevant to verify gradients of any nature in a thick sample.

In this study, the ageing behaviour of an in-house formulated bicomponent RTV PDMS material filled with silica and, in lower fraction, calcium carbonate is explored. The fillers are acting as a thickener for the PDMS matrix, which is important for the industrial application where the seal needs to be injected inside spatial voids. The targeted sealing application requires long-term stability of the material subjected to moderate temperature. An accelerated ageing protocol is used with a temperature range between 30 °C and 70 °C up to 600 days with and without access to air. The originality of this paper is in the exploration of skin-core gradients due to ageing in addition to the usual bulk characterisations. To do so, infrared spectroscopy and nanoindentation are employed on the surface and on the cross-section of the aged samples. To complete the understanding of the global mechanism, the mass loss measurements and swelling at equilibrium tests are performed on filled and unfilled PDMS materials. The impact of ageing on the bulk mechanical behaviour is verified by the tensile tests.

## 2. Materials and methods

### 2.1. Materials

The material is two-part RTV PDMS. The PDMS chains were condensed from short PDMS chains (silicone oil) and crosslinked using tetraethyl orthosilicate (TEOS) through a condensation reaction catalysed by a tin salt catalyst, representing 2 per hundred rubber (phr). The material is reinforced with SiO₂ and CaCO₃ particles at 50 phr (15.5 % of volume fraction) and 9 phr (1.5 % of volume fraction), respectively. Both fillers provide an increase of the mechanical properties of the PDMS material, while composition offers relevant injection and sealing performance of the industrial application. The SiO₂ particles have minimal, median, and maximal equivalent circular diameters of 0.7 µm, 4.5 µm, and 37 µm, respectively. The corresponding values for CaCO₃ particles are 0.7 µm, 7 µm, and 45 µm, respectively.

The mixture is degassed under vacuum for 10 min to evacuate air bubbles and then poured into a smooth Teflon®-lined mould to obtain a large 2 mm thick pad. The mould is sealed, and the crosslinking process is initiated by the interaction with the moisture from the ambient air trapped within the mould. The assembly is maintained at room temperature for at least 100 h to allow the crosslinking process to complete. Following demoulding, the samples are cut directly from the pads using H2-shaped blades according to the ISO37 [24].

The same ingredients with the same manufacturing protocol are used to produce PDMS pads without reinforcing particles. These materials are called “uPDMS”, for unfilled PDMS, in contrast to “fPDMS”, for filled PDMS.

### 2.2. Ageing protocol

The H2-shaped fPDMS and uPDMS samples were aged under ambient air using ventilated ovens, at temperatures ranging from 30 °C to 70 °C for durations between 10 days and 600 days. The impact of moisture was not in the scope of this study, thus the relative humidity (RH) was kept at 30 % during the ageing at 30 °C in a climatic chamber (Climats® EXCAL2 240 HE). The 30 °C - 30 %RH was chosen to reduce interaction with ambient moisture at 30 °C. For ageing temperatures above 30 °C, relative humidity decreases rapidly to 2 % at 70 °C as by psychrometric diagram. The furnaces employed for ageing between 40 °C and 70 °C were Memmert® UF110 and UFE400 models. All samples were hung using clips inside the ovens without any contact with the oven parts.

Some of H2 samples were wrapped in aluminium foils to prevent any direct contact with the ageing atmosphere. These samples were subjected to the same ageing protocol as the others and will be referred to as “confined”.
A Sartorius RC210D weighing scale was used to measure the mass of the samples before and after ageing, which was around 1.5 g. The measurement precision is 0.01 mg, with an uncertainty of less than 0.01 %. The relative mass loss W is defined by Equation (#1):

$$W(\%) = \frac{W_a - W_0}{W_0} \times 100 \qquad (1)$$

where W₀ is the initial mass of each sample and W_a is the mass after ageing. The mass of the samples was measured from 3 to 5 min after removal from the oven.

## 2.3. Swelling at equilibrium measurements

The rubber network was characterised using solvent swelling in methylcyclohexane [25,26], following these steps: the samples were weighed and immersed in a large volume of solvent to ensure complete submersion throughout the swelling process. Samples were allowed to swell for at least 96 h to reach swelling equilibrium, then weighed (W_s), and subsequently dried in a vacuum chamber at room temperature for 24 h. After the extraction of solvent, the dried samples were weighed again.

The soluble fraction f_s, representing the percentage of uncrosslinked species inside the material, was obtained from Equation (#2) [26]:

$$f_s = \frac{W_f - W_i \times c}{W_i \times (1 - c)} \times 100 \qquad (2)$$

where W_i is the initial mass, W_f is the final mass after drying, c is the weight fraction of insoluble fillers (SiO₂ + CaCO₃).

The swelling ratio (SR) is calculated from Equation (#3):

$$SR = \frac{W_s}{W_f} \qquad (3)$$

where W_s is the equilibrium swollen mass.

Several physical indicators can be used to measure the ability of the network to swell (swelling ratio [10], network functionality [27], Flory–Huggins interaction parameter [28], volumetric fraction of rubber [29], degree of swelling [30]). However, it is difficult to admit that the solvent uptake after ageing will be the same as for the virgin samples. If chain scission or crosslinking appears during the ageing, the volume fraction of polymer can evolve as well as the local particles volume fraction [10]. Moreover, for calculating the degree of swelling, the densities of the rubber and solvent need to be known precisely. After ageing, the density of the rubber can change due to shrinking [30]. For the volumetric fraction of rubber, the mass fraction of the fillers is required [29]. For these reasons, the swelling ratio, SR, will be used to follow the evolution of the rubber network with ageing.

## 2.4. ATR-FTIR experiments

Infrared spectroscopy measurements were conducted using a Thermo Fisher Nicolet™ iS20 spectrometer. An attenuated total reflection (ATR) diamond detector was employed, with a spectral range of 600 cm⁻¹ to 4000 cm⁻¹ using a 1 cm⁻¹ spectral resolution. The obtained spectrum is the result of 16 co-added scans. The diamond crystal, with a diameter of 2 mm. A consistent force of 1 N was applied to all samples. The analysis was performed on virgin and aged fPDMS, uPDMS and, separately, on the SiO₂ and CaCO₃ powders.

The chemical signature of virgin fPDMS samples is composed of the spectra of its individual components, as shown in Fig. 1. One major contribution originates from the Si-O-Si bonds, which vibrate in the range of 1000 cm⁻¹ and 1080 cm⁻¹ and are attributed to the backbone of the PDMS chains. The second major contribution is located at 785 cm⁻¹ and arises from the CH bonds, which link the methyl groups to the PDMS backbone. Notable differences between the spectra of uPDMS and fPDMS include the presence of an absorbance peak at 1163 cm⁻¹ and a

<figure><img src="image_2.png" /><figcaption>Fig. 1. ATR-FTIR spectra of SiO₂, CaCO₃ particles, virgin uPDMS and virgin fPDMS samples within the spectral range 1600–600 cm⁻¹.</figcaption></figure>

higher absorbance level between 1000 cm⁻¹ and 1080 cm⁻¹ for fPDMS in the fPDMS spectrum. These differences are associated with the presence of SiO₂ particles within the fPDMS samples. Table 1 refers all the wavenumbers that can be identified on both the uPDMS and fPDMS ATR-FTIR spectrums.

## 2.5. Uniaxial tensile tests

Uniaxial tensile tests were performed using an Instron® ElectroPulse E1000 equipped with a ±250 N capacity Dynacell™ load cell. The tests were carried out at 22 ± 0.5 °C with a crosshead rate of 30 mm/min. The strain was measured using marker tracking with a Mako U-130 acquisition camera. The elastic moduli were calculated using an incompressible Neohookean hyperelastic model [37,38] based on the elongation range [1.1; 1.3]. Under uniaxial tension, the engineering stress σ is related to the elongation λ as follows:

$$\sigma = 2C_{01} \left( \lambda - \frac{1}{\lambda^2} \right) \qquad (4)$$

where C₀₁ is the Neo-Hooke constant.

<table><caption>**Table 1**

Interpretation of FTIR peaks of the SiO₂ and CaCO₃ particles, uPDMS, and fPDMS samples.</caption><thead><tr><th>Material</th><th>Wavenumber (cm<sup>-1</sup>)</th><th>Group contribution</th><th>References</th></tr></thead><tbody><tr><td rowspan="6">PDMS</td><td>686</td><td>CH<sub>3</sub> rocking vibration</td><td rowspan="3">[31]</td></tr><tr><td>701</td><td>Si-OH bond vibration</td></tr><tr><td>784</td><td>CH<sub>3</sub> rocking</td></tr><tr><td>863</td><td>Si-C stretching in Si-CH<sub>3</sub></td><td rowspan="2">[32]</td></tr><tr><td>1008</td><td>Si-CH<sub>3</sub> asymmetric rocking</td></tr><tr><td>1080</td><td>Si-O-Si symmetric stretching</td><td>[33]</td></tr><tr><td>1257</td><td>Si-O-Si asymmetric stretching</td><td>[32]</td></tr><tr><td rowspan="2">CaCO<sub>3</sub></td><td>712</td><td>CH<sub>3</sub> deformation in Si-CH<sub>3</sub></td><td>[32]</td></tr><tr><td>873</td><td>C-O out of plane bending vibration</td><td>[34]</td></tr><tr><td rowspan="7">SiO<sub>2</sub></td><td>1407</td><td>C-O in-plane bending vibration</td><td>[34]</td></tr><tr><td>695</td><td>CO<sub>3</sub><sup>2-</sup> ν-asymmetric stretching</td><td>[34]</td></tr><tr><td>778</td><td>Si-O symmetric bending vibration</td><td>[35]</td></tr><tr><td>797</td><td>Si-O-Si bending vibration</td><td>[36]</td></tr><tr><td>1058</td><td>Si-O-Si stretching vibration</td><td>[36]</td></tr><tr><td>1074</td><td>Si-O-Si stretching vibration</td><td>[36]</td></tr><tr><td>1162</td><td>Si-O-Si stretching vibration</td><td>[36]</td></tr></tbody></table>
## 2.6. Nanoindentation tests

The nanoindentation machine employed was a NanoVantage indenter from MicroMaterials Ltd®. A cono-spherical diamond tip with a nominal radius of 25 µm and a 90° conical aperture angle was used. The experiments were conducted using a passive detection method, which allowed for precise determination of the contact point on the measured curves after experiment. The tests were force-controlled with a maximum force of 600 µN and a loading rate of 5 µN s⁻¹. The maximum force was maintained for 30 s, followed by unloading at a rate of 5 µN s⁻¹ until the contact is lost. The force-displacement data obtained from the indentation tests were analysed according to the Oliver-Pharr method [39]. This method provides the elastic modulus E<sub>IT</sub> using equation (#7):

$$E_{IT} = \sqrt{\pi} \times \frac{S}{2 \times \beta \times \sqrt{A_p}} \quad (7)$$

S is the contact stiffness obtained from the force-displacement data, β is a constant related to the indenter tip shape (equals to 1 for a spherical tip), A<sub>p</sub> is the projected contact area between the indenter and the material surface.

## 2.7. Scanning electron microscopy

Scanning electron microscopy (SEM) was used to observe the surfaces of the virgin and aged fPDMS. Samples surfaces were coated using Au/Pd deposits and observed using a TESCAN SEM using 5 kV accelerating tension from X800 to X1000 magnification.

# 3. Results and discussions

This section starts with the characterisations of the exposed to air fPDMS and uPDMS. Global and local mechanical and physical characterisations are presented. The results and discussions are then completed with aged under confinement samples. The ageing mechanism is summarised and schematised in the end of the section.

## 3.1. Impact of thermal ageing on bulk properties

The mass evolution of the samples after thermal ageing is presented in Fig. 2. Each result is the average over 3 samples. The mass loss increases with ageing duration for all investigated temperatures and ageing periods. The temperature accelerates the mass loss, which attains almost 8 % at the most severe condition of 100 days at 70 °C. However, even at the lowest temperature of 30 °C, 1 % of mass is lost after 100 days and more than 5 % after 600 days. There is no evidence of stabilisation of this process from the data. Although, the rate is higher at the first 300 days (4 % lost) than in the second period of ageing (around 1 % more). A double mass loss rate can be observed at all temperatures,

<figure><img src="image_3.png" /><figcaption>Fig. 2. Mass evolution of the fPDMS samples after ageing between 30 °C and 70 °C for 10–600 days and of the uPDMS at 70 °C for 10–100 days.</figcaption></figure>

symbolised with dashed lines for the first rate and dotted lines for the second rate. This double rate comes from the different evaporation energy required for each component of the network.

The mass evolutions of uPDMS and fPDMS can be compared at 70 °C up to 100 days. The fPDMS samples show a slightly higher mass loss, reaching a total decrease of 7.81 % against 6.72 % for uPDMS after 100 days. This small difference suggests that the ageing process is due to the loss of the PDMS matrix rather than the particles. The slightly higher mass loss in fPDMS samples is likely due to the detachment of particles from the fPDMS sample surface upon removal from the oven. Indeed, a layer of white powder was observed on some samples after the most severe ageing durations.

To understand the evolution of the fPDMS network with ageing, the swelling ratio (SR) and the soluble fraction (fs) were obtained from swelling measurements. Since the samples were cut from several pads manufactured at different days (from the same batch), the control measurements were done on the virgin samples from each pad. To remove possible dispersity due to manufacturing process, the soluble fractions are presented normalized by the data from these control samples. The virgin pads of fPDMS presented a swelling ratio of 2.27 ± 0.0486, its soluble fraction was 4.83 ± 0.135 %. The data presented in Fig. 3 reveal that the swelling ratios remain rather constant for all the fPDMS.

Fig. 4 displays the evolution of the normalized soluble fraction. For all ageing temperatures, the soluble fraction decreases over time, with a more pronounced decrease at higher temperatures.

While usual ageing mechanisms for silicone elastomers through chain scission increase swelling ratio and soluble fraction [10,11], the swelling ratios remain relatively unmodified for fPDMS while the soluble fraction decreases significantly. The constancy of the swelling ratio indicates that the network is not significantly affected by the thermal ageing. During the crosslinking, the TEOS molecules attach to the ends of PMDS chains from oil creating the network. The decrease of soluble fraction with temperature and time indicates the evaporation of the non-bonded to the polymer network chemical species. Given the composition of the fPDMS samples and the known crosslinking reaction [27], these species primarily consist of ethanol (C₂H₆O) and water (H₂O) and unattached to the network PDMS oligomers, which are trapped within the material after room-temperature vulcanisation. The most significant reduction in the soluble fraction is observed at 70 °C after 150 days, reaching a 50 % decrease from its original value.

The tensile response of virgin and aged fPDMS samples is presented in Fig. 5. The impact of the ageing on small strain elastic behaviour was evaluated through incompressible Neo-Hooke model, while the embrittlement was quantified by elongation at break. The virgin fPDMS samples present a NeoHooke parameter of C₀₁ = 1.03 MPa, and an elongation at break of 2.12 ± 0.18. These figures reveal that, at all ageing temperatures, the samples exhibit rather similar stress-strain behavior with a lower elongation at break after ageing. This trend is

<figure><img src="image_4.png" /><figcaption>Fig. 3. Evolution of the swelling ratios of the fPDMS after thermal ageing between 40 °C and 70 °C.</figcaption></figure>
<figure><img src="image_5.png" /><figcaption>Fig. 4. Evolution of the soluble fraction for filled PDMS after thermal ageing.</figcaption></figure>

consistent across all ageing durations: increasing the ageing duration results in a decrease of the elongation at break. For instance, ageing at 70 °C for 100 days reduces the elongation at break by 51 %. The reduction is only 6 % after the ageing at 40 °C for the same duration. The figure also indicates that the NeoHooke constants $C_{01}$ of the samples remain relatively constant regardless of the ageing conditions, with an increase of at most 20 % after 100 days at 70 °C.

Since the swelling results did not reveal change in swelling ratio, the observed embrittlement and stiffening could be attributed to an appearance of a skin-core effect. This skin-core effect can induce local gradients in mechanical properties, which could contribute to the degradation of the overall mechanical performance. These hypotheses will be explored through local measurements in the following section.

## 3.2. Characterisation of the outer surface and highlight of structural gradients in the aged samples

To characterise local gradients in aged materials, it is necessary to use highly resolved techniques with small tested volume. In this study, scanning electron microscopy (SEM) was used to observe the outer surface of virgin and aged samples. ATR-FTIR spectroscopy was used to characterise the outer layers of the material, both on the surface exposed to ageing and on the cross-section cut. These results were completed with nanoindentation across the cross-section of the cut sample to highlight the gradients and evaluate the thickness of skin layer.

Firstly, the exposed to ageing surfaces of the fPDMS were observed with SEM. Fig. 6 clearly evidences that the surfaces of aged samples (see Fig. 6(b and c-d)) are rougher than the surface of the virgin sample (see

Fig. 6(a)). It seems from these figures, that the surface is progressively covered with silica and calcium carbonate particles with ageing. To verify this hypothesis and discern the chemical composition of the outer skin, ATR-FTIR spectroscopy was used.

Fig. 7 presents the FTIR spectra obtained on the surface of virgin uPDMS and fPDMS samples, as well as four aged samples subjected to prolonged ageing durations. For uPDMS aged at 70 °C for 100 days, no discernible changes are observed in either the intensity or the wave-number of the peaks. This confirms that no oxidation phenomena occur under moderate ageing conditions and no new bonds appear after ageing. Considering high mass loss of almost 8 % for this sample (see Fig. 2), this confirms that the evaporation of the species non-attached to the network is the main ageing mechanism for this RTV PDMS.

For the aged fPDMS samples, the intensity of the C-H rocking vibrations at 865 cm⁻¹ remains unchanged. The peaks associated with Si-O-Si bonds, at 695 cm⁻¹, 1058 cm⁻¹, 1074 cm⁻¹ and 1162 cm⁻¹ (see Table 1), exhibit higher intensities for the ageing conditions of 40 °C for 400 days, 50 °C for 300 days, and 70 °C for 100 days. At 1163 cm⁻¹, the peak intensity increases for all ageing conditions. The peak located at 1257 cm⁻¹ shows a lower intensity for all ageing conditions. The peak around 1050 cm⁻¹ shifts significantly towards lower wavenumbers, and the peaks around 1150 and 690 cm⁻¹, attributed to SiO₂ and barely discernible in the unaged material, become well-defined. The single peak around 790 cm⁻¹ splits in two, increasingly resembling the double peak of silica.

These results indicate a more significant contribution from silica particles in the aged fPDMS samples compared to the unaged material. However, as for uPDMS, there is no evidence of new bonds. These measurements allow us to conclude that the skin is indeed filled with higher concentration of the particles and the matrix has shrunk due to the evaporation of entrapped in the network species.

Once the outer surface is investigated, it is necessary to characterise the core of the aged samples in order to conclude about possible gradients. To do so, nanoindentation measurements were conducted through the thickness of several thermally aged samples cut with a razor blade. This technique provides local elastic modulus at a micrometer scale and, thus, can be used for a mapping of the elastic properties through the cross-section of the sample with very high resolution.

Fig. 8 illustrates the evolution of the elastic indentation modulus of the virgin sample and three samples aged at the most severe conditions. The virgin sample exhibits rather constant modulus through the thickness confirming homogeneous distribution of particles and matrix without any gradients across the sample. There is no skin-core effect for this

<figure><img src="image_6.png" /><figcaption>Fig. 5. a) Uniaxial tensile tests (engineering stress vs elongation) at room temperature, b) $C_{01}$ parameters and c) elongation at break of the fPDMS for different ageing conditions.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>Fig. 6. SEM images of the exposed surfaces of (a) virgin fPDMS sample, (b) aged at 40 °C for 100 days sample, (c) aged at 50 °C for 100 days sample, and (d) aged at 40 °C for 400 days.</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>Fig. 7. ATR-FTIR spectra of virgin and aged uPDMS and fPDMS samples.</figcaption></figure>

sample.

In contrast, all tested aged samples display a mechanical gradient from the surface to the core. Specifically, each aged sample demonstrates an increased stiffness near the exposed to air surface and the modulus in core same as for the virgin material. The thickness of this gradient is around 300–400 nm, though difficult to identify precisely. Although the stiffness gradients are commonly observed in aged thick elastomers [22,23,40,41], this is usually due to the oxidation diffusion layer, which is not the case here. Combining the FTIR results with these nanoindentation measurements, it is clear that there is no oxidation and the mechanical gradient is only due to the different particle

<figure><img src="image_9.png" /><figcaption>Fig. 8. Evolution of the local indentation modulus measured in the thickness of the virgin and aged samples.</figcaption></figure>

concentration across the sample. The absence of new bonds in the core of the aged sample is confirmed by the ATR-FTIR measurement on its cross-section (see Fig. 9). Indeed, the spectrum of the core was similar to that of the unaged sample with only a small difference at around 1050 $cm^{-1}$. This spectrum is very different from the one measured at the external surface of the sample, already presented in Fig. 7 and discussed above.

### 3.3. Summary of the ageing mechanism and the effect of confinement

All the results presented above indicate that the change of properties is heterogeneous and is due to the evaporation of the oligomers and by-products of crosslinking. The results of characterisations of the confined fPDMS samples confirm this. In particular, the ATR-FTIR spectrum of the outer surface of wrapped in aluminum foil fPDMS exposed to 40 °C for 200 days, shown in Fig. 9, presents no difference with virgin fPDMS.
<figure><img src="image_10.png" /><figcaption>**Fig. 9.** FTIR spectra conducted on the surface, in the core of the 40 °C-400 days aged fPDMS sample as well as the outer surface of the fPDMS aged in confinement for 200 days at 40 °C.</figcaption></figure>

This means that chemical bonds of the surface are not altered by temperature. The mass loss (Fig. S1) and soluble fraction (Fig. S2) of several samples aged at temperatures from 30 to 70 °C, presented in **Supplementary document**, indicate very few changes. For instance, the mass loss is only 0.3 % after 209 days of ageing at 70 °C.

The impact of the exposure of the filled RTV PDMS to air for long durations at moderate temperature is schematised in Fig. 10. Due to evaporation of the oligomers and by-products of crosslinking reaction, the PDMS matrix shrinks. The diffusion is mostly limited to the outer layer of the sample, thus creating a distinct skin-core structure in the 2 mm thick samples. The particle fraction is higher in the skin layer, which increases local elastic modulus and global elastic modulus. This high concentration of the particles on the surface induces the embrittlement of the material. In the core, the material remains relatively unchanged by thermal exposure.

## 4. Conclusions

In this work, the thermal ageing of a room temperature vulcanised polydimethylsiloxane silicone elastomer under moderate temperatures was studied. The material is two-part RTV crosslinked by polycondensation with TEOS as crosslinking ageing and tin catalyst, reinforced with silica and calcium carbonate particles. The temperature range of interest spanned from 30 °C to 70 °C over durations of 10–600 days. The evolution of the material was examined through mass loss, swelling at equilibrium, and infrared spectroscopy measurements. Mechanical property changes were evaluated at the macroscopic scale via uniaxial tensile testing, and at the local scale through nanoindentation measurements.

The results demonstrated that the change of properties of fPDMS at these temperatures is due to the evaporation of various chemical species, including water, ethanol, and uncrosslinked PDMS oligomers. This was confirmed by the ageing in confinement, which caused no significant changes in mechanical or structural properties of fPDMS. Moreover, the exposed to air ageing resulted in skin-core structure with a higher concentration of SiO₂ and CaCO₃ particles in the skin, seen in the SEM observation of the surface. The shrinkage of the rubber induced the embrittlement of material observed in tensile tests. The tensile Neo-Hooke C₀₁ parameter increased very slightly with ageing. The structural and mechanical gradients were detected using infrared spectroscopy and nanoindentation technique. Indeed, local elastic moduli were impacted by ageing close to the outer surface but remained similar to virgin material in the core.

### CRediT authorship contribution statement

**Florian Feyne:** Writing – review & editing, Writing – original draft, Visualization, Validation, Supervision, Software, Project administration, Methodology, Investigation, Formal analysis, Data curation, Conceptualization.
**Eric Le Bourhis:** Writing – review & editing, Writing – original draft, Visualization, Validation, Supervision, Methodology, Investigation, Funding acquisition, Formal analysis, Conceptualization.

**Florian Lacroix:** Writing – review & editing, Writing – original draft, Visualization, Validation, Supervision, Methodology, Investigation, Funding acquisition, Formal analysis, Conceptualization.

**Laurence Autissier:** Writing – review & editing, Writing – original draft,

<figure><img src="image_11.png" /><figcaption>**Fig. 10.** Schematics of thermal ageing effect on fPDMS when exposed to air.</figcaption></figure>
Visualization, Validation, Supervision, Methodology, Investigation, Funding acquisition, Formal analysis, Conceptualization. **Julie Pepin:** Writing – review & editing, Writing – original draft, Visualization, Validation, Supervision, Software, Methodology, Investigation, Formal analysis. **Olga Smerdova:** Writing – review & editing, Writing – original draft, Visualization, Validation, Validation, Supervision, Software, Methodology, Investigation, Funding acquisition, Formal analysis, Data curation, Conceptualization.

## Declaration of competing interest

The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: FEYNE reports financial support was provided by French National Research Agency - ANR-11-LABEX-0017-01. If there are other authors, they declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Acknowledgments

This work was partially funded by the French Government programs "Investissements d'Avenir" LABEX INTERACTIFS (reference ANR-11-LABEX-0017-01).

Pprime Institute gratefully acknowledges "Contrat de Plan Etat - Région Nouvelle-Aquitaine" (CPER) as well as the "Fonds Européen de Développement Régional (FEDER)" for their financial support to the reported work.

## Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.polymertesting.2025.109022.

## Data availability

Data will be made available on request.

## References

[1] P. Jerschow, Silicone Elastomers, 12, Rapra Technology Limited, 2001, 5 vols.

[2] G. Deshpande, M.E. Rezac, Kinetic aspects of the thermal degradation of poly (dimethyl siloxane) and poly(dimethyl diphenyl siloxane), Polym. Degrad. Stabil. 76 (1) (Jan. 2002) 17–24, https://doi.org/10.1016/S0141-3910(01)00261-0.

[3] D. Oldfield, T. Symes, Long term natural ageing of silicone elastomers, Polym. Test. 15 (2) (1996) 115–128.

[4] Y.-R. Luo, Comprehensive Handbook of Chemical Bond Energies, CRC Press, Boca Raton, 2007, https://doi.org/10.1201/978142007282.

[5] D. Lansade, Stabilisation de résines poly(dimethylsiloxane) en environnement spatial radiatif, Université de Bordeaux, 2020. Thèse de doctorat.

[6] M.J. Hunter, J.F. Hyde, Organo-silicon polymers; the cyclic dimethyl siloxanes, J. Am. Chem. Soc. 68 (Apr. 1946) 667–672, https://doi.org/10.1021/ja01208a042.

[7] D. Atkins, C. Murphy, C. Saunders, Polymethylsiloxanes...Thermal and oxidation stabilities, Ind. Chem. 39 (11) (Nov. 1947) 1395–1401, https://doi.org/10.1021/ie50455a606.

[8] T.H. Thomas, T.C. Kendrick, Thermal analysis of polydimethylsiloxanes. I. Thermal degradation in controlled atmospheres, J. Polym. Sci. Part -2 Polym. 7 (3) (1969) 537–549, https://doi.org/10.1002/pol.1969.160070308.

[9] N. Grassie, I.G. Macfarlane, The thermal degradation of polysiloxanes—I. Poly (dimethylsiloxane), Eur. Polym. J. 14 (11) (1978) 875–884.

[10] T. Khaled, A. Rios de Anda, E. Richaud, Macromolecular and mechanical changes in aged silicones, Polym. Degrad. Stabil. 234 (Apr. 2025) 111211, https://doi.org/10.1016/j.polymdegradstab.2025.111211.

[11] L. Pires, et al., Thermal and thermo-oxidative accelerated ageing of fluorosilicone rubber, Polym. Degrad. Stabil. 232 (Feb. 2025) 111150, https://doi.org/10.1016/j.polymdegradstab.2024.111150.

[12] M. Celina, K.T. Gillen, R.A. Assink, Accelerated ageing and lifetime prediction: review of non-Arrhenius behaviour due to two competing processes, Polym. Degrad. Stabil. 90 (3) (Dec. 2005) 395–404, https://doi.org/10.1016/j.polymdegradstab.2005.05.004.

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

[14] Wacker, Room Temperature Vulcanizing (RTV) silicones. Material and processing guidelines [Online]. Available: https://www.wacker.com/h/en-us/medias/7461-EN.pdf. (Accessed 30 June 2025).

[15] Z. Brounstein, J. Zhao, D. Geller, N. Gupta, A. Labouriau, Long-Term thermal ageing of modified sylgard 184 formulations, Polymers 13 (18) (Jan. 2021) 3125, https://doi.org/10.3390/polym13183125.

[16] M. Patel, A.R. Skinner, Thermal ageing studies on room-temperature vulcanised polysiloxane rubbers, Polym. Degrad. Stabil. 73 (3) (2001) 399–402.

[17] M. Patel, S. Chinn, R.S. Maxwell, T.S. Wilson, S.A. Birdsell, Compression set in gas-blowed condensation-cured polysiloxane elastomers, Polym. Degrad. Stabil. 95 (12) (Dec. 2010) 2499–2507, https://doi.org/10.1016/j.polymdegradstab.2010.07.038.

[18] A. Labouriau, T. Robison, L. Meincke, D. Wrobleski, D. Taylor, J. Gill, Ageing mechanisms in RTV polysiloxane foams, Polym. Degrad. Stabil. 121 (Nov. 2015) 60–68, https://doi.org/10.1016/j.polymdegradstab.2015.08.013.

[19] A.N. Chaudhry, N.C. Billingham, Characterisation and oxidative degradation of a room-temperature vulcanised poly(dimethylsiloxane) rubber, Polym. Degrad. Stabil. 73 (3) (Jan. 2001) 505–510, https://doi.org/10.1016/S0141-3910(01)00139-2.

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

[21] K.T. Gillen, R.L. Clough, C.A. Quintana, Modulus profiling of polymers, Polym. Degrad. Stabil. 17 (1) (Jan. 1987) 31–47, https://doi.org/10.1016/0141-3910(87)90046-2.

[22] M. Celina, J. Wise, D.K. Ottesen, K.T. Gillen, R.L. Clough, Correlation of chemical and mechanical property changes during oxidative degradation of neoprene, Polym. Degrad. Stabil. 68 (2) (2000) 171–184.

[23] K.T. Gillen, R.L. Clough, Polymer ageing insights available from modulus profiling data, Polym. Eng. Sci. 29 (1) (1989) 29–35, https://doi.org/10.1002/pen.760290107.

[24] International organisation for standardisation, 'Rubber, Vulcanized or Thermoplastic — Determination of Tensile stress-strain Properties', ISO. [Online]. Available: https://www.iso.org/standard/86892.html.

[25] A.M. Bueche, Interaction of polydimethylsiloxanes with swelling agents, J. Polym. Sci. 15 (79) (Jan. 1955) 97–103, https://doi.org/10.1002/pol.1955.120157908.

[26] A.M. Stricher, R. Rinaldi, C. Barrès, F. Ganachaud, L. Chazeau, How I met your elastomers: from network topology to mechanical behaviours of conventional silicone materials, RSC Adv. 5 (66) (2015) 53713–53725.

[27] F.W. Van der Weij, The action of tin compounds in condensation-type RTV silicone rubbers, Makromol. Chem. Macromol. Chem. Phys. 181 (12) (1980) 2541–2548.

[28] E. Delebecq, F. Ganachaud, Looking over liquid silicone rubbers: (1) network topology vs chemical formulations, ACS Appl. Mater. Interfaces 4 (7) (Jul. 2012) 3340–3352, https://doi.org/10.1021/am300502r.

[29] J.L. Valentín, J. Carretero-González, I. Mora-Barrantes, W. Chassé, K. Saalwächter, Uncertainties in the determination of cross-link density by equilibrium swelling experiments in natural rubber, Macromolecules 41 (13) (Jul. 2008) 4717–4729, https://doi.org/10.1021/ma8005087.

[30] A. Simon, J. Pepin, D. Berthier, S. Méo, Degradation mechanism of FKM during thermo-oxidative ageing from mechanical and network structure correlations, Polym. Degrad. Stabil. 208 (2023) 110271.

[31] K. Jeronimo, V. Koutsos, R. Cheung, E. Mastrapaolo, PDMS-ZnO piezoelectric nanocomposites for pressure sensors, Sensors 21 (17) (2021) 5873.

[32] L.M. Johnson, et al., Elastomeric microparticles for acoustic mediated bioseparations, J. Nanobiotechnol. 11 (1) (Dec. 2013) 22, https://doi.org/10.1186/1477-3155-11-22.

[33] F.A.B. Silva, F.A. Chagas-Silva, F.H. Florenzano, F.L. Pissetti, Poly (dimethylsiloxane) and Poly[vinyltrimethoxysilane- co -2-(dimethylamino) ethyl methacrylate] based cross-linked organic-inorganic hybrid adsorbent for Copper (II) removal from aqueous solutions, J. Braz. Chem. Soc. (2016), https://doi.org/10.9535/0103-5053.20160110.

[34] L. Ogbesta, et al., Rapid and application-tailored assessment tool for biogenic powders from crustacean shell waste: fourier transform-infrared spectroscopy complemented with X-ray diffraction, scanning electron microscopy, and nuclear magnetic resonance spectroscopy, ACS Omega 6 (42) (Oct. 2021) 27773–27780, https://doi.org/10.1021/acsomega.1c03279.

[35] B. Saikia, Spectroscopic estimation of geometrical structure elucidation in natural SiO 2 crystal, J. Mater. Phys. Chem 2 (2) (2014) 28–33, https://doi.org/10.12691/jmpc.2-2-3.

[36] N. Tahiri, L. Khouchaf, M. Elaatmani, G. Louarn, A. Zegzouti, M. Daoud, Study of the thermal treatment of SiO2 aggregate, in: IOP Conference Series: Materials Science and Engineering, IOP Publishing, 2014 012002 [Online]. Available: http://iopscience.iop.org/article/10.1088/1757-899X/62/1/012002/meta.

[37] L.R.G. Treloar, The elasticity of a network of long-chain molecules. I, Trans. Faraday Soc. 39 (1943) 36–41.

[38] L.R.G. Treloar, The elasticity of a network of long-chain molecules—Ii, Trans. Faraday Soc. 39 (1943) 241–246.

[39] W.C. Oliver, G.M. Pharr, An improved technique for determining hardness and elastic modulus using load and displacement sensing indentation experiments, J. Mater. Res. 7 (6) (1992) 1564–1583.

[40] C. Fradet, F. Lacroix, G. Berton, S. Méo, E. Le Bourhis, Instrumented indentation of an elastomeric material, protocol and application to vulcanisation gradient, Polym. Test. 81 (2020) 106278.

[41] C. Ovalle, L. Laiarinandrasana, R. Borisov, G. Miquelard-Garnier, B. Fayolle, Laminated rubber bearings subjected to combined shear and compressive loading: Laminated rubber bearings subjected to combined shear and compressive loading:
heterogeneous ageing and crack growth in mode II, in: Constitutive Models for Rubber XII, first ed., CRC Press, London, 2022, pp. 46–52.