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The Temperature Effect in the High-Density Polyurethane Foams: A Comparison Between the Petroleum and Soy-Based Materials

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28 September 2026

Posted:

30 September 2026

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Abstract
The combined influence of chemical structure and cellular morphology on the performance of high-density polyurethane foams has received limited attention. In this work, three foams based on a commercial petrochemical polyol (SD441) were compared with three foams produced from a mixture of epoxidized soybean-oil polyol and glycerol at densities ranging from approximately 100 to 300 kg/m³. The biobased foams exhibited substantially lower compressive strength than the petrochemical materials, with longitudinal strength ranging from 0.60 to 4.07 MPa compared with 1.03 to 8.92 MPa for the petrochemical foams. Electron microscopy revealed larger cells, greater anisotropy, and pronounced morphological heterogeneity in the biobased materials, particularly at higher densities. Despite their lower initial storage modulus, the biobased foams exhibited greater resistance to thermally induced softening: at 170 °C, they retained approximately 32–39% of their initial storage modulus, compared with only 6–8% for the petrochemical foams. A second DMA heating run, together with FTIR and DSC results, indicated irreversible structural changes during the first heating cycle. The reduction in residual isocyanate content and changes in the carbonyl and amide II regions were consistent with further network development during heating, suggesting additional post-curing in the biobased foams. DSC identified two exothermic events in the biobased materials: a broad event between 60 and 245 °C, consistent with additional reactions, and a second event between approximately 255 and 300 °C that overlapped with the onset of thermal degradation observed by TGA. In contrast, the type of blowing agent strongly influenced the degradation temperature but had no significant effect on mechanical performance within the investigated density range. Overall, the results demonstrate a distinct trade-off between mechanical and thermomechanical performance in high-density biobased polyurethane foams, highlighting the influence of polyol molecular architecture, network development, and cellular morphology on their behaviour.
Keywords: 
high density polyurethane foams
;  
biobased foams
;  
soy-oil polyurethane foams
;  
thermal dynamic mechanical stability
;  
glass transition temperature
Subject: 
Chemistry and Materials Science  -   Materials Science and Technology

Introduction

Sandwich panels are widely used in construction and vehicle applications. The development of biobased cores has gained increasing attention in recent years. Accordingly, Oliveira et al. presented a comprehensive review on biobased sandwich panels, highlighting the importance of replacing the three main components—core, adhesive, and skin—with biobased alternatives [1]. Among core materials, the main alternatives reported include cork (0.2–0.7 MPa, 140–250 kg/m³), polylactic acid (PLA) foams (2.4 MPa, 320 kg/m³), and balsa wood (5.5–22 MPa, 100–280 kg/m³). In this work, we focus on an alternative core material based on high-density soy-oil polyurethane. The use of high-density foams is supported by the well-established relationship between mechanical properties and density: Ashby presented simple models describing the increase in elastic modulus and other mechanical properties with density [2], and, in the same direction, Saint-Michel et al. experimentally linked changes in mechanical performance to density and aging processes [3]. In this context, increasing the density represents a potential strategy for improving the elastic modulus and compressive strength of polyurethane foams. Therefore, the development of high-density soy-based foams was proposed as a means of combining the mechanical advantages associated with increased density with the use of a biobased raw material.
Regarding polyurethane formulation, the reactions occurring during curing are critical, as they directly influence the composition of the final material. Polymerization kinetics are governed by two main reactions: urethane formation from the polyol–isocyanate mixture, and the blowing reaction between water and isocyanate, which produces CO₂, an amine, and releases heat. This heat drives the evaporation of physical blowing agents. However, three additional reactions condition the final chemical structure: the formation of urea (from amine and isocyanate), biuret (from urea and isocyanate), and allophanate (from urethane and isocyanate) [4]. The urea and biuret formation has received considerably less attention in low- and medium-density rigid foams, where cellular morphology and cell-wall integrity are generally regarded as the dominant factors governing mechanical performance [5,6,7]. In high-density foams, however, this effect becomes more relevant: as cell volume decreases and wall thickness increases, the relative importance of cell integrity diminishes while the chemical composition of the cell wall gains greater weight. In particular, the post-cure cycle can promote the progression of these secondary reactions. Increasing temperature can also trigger foam decomposition; in conventional polyurethane foams, this has been associated with temperatures between 170°C and 200°C [8]. In recent studies on polyisocyanurate foams, the use of soy oil has been linked to increased resistance to thermal degradation and flammability, with degradation onset around 250°C [9]. A similar result was reported for foams based on epoxidized rapeseed oil polyol, where the authors also compared this polyol with a petrochemical polyester polyol [10]. Along the same lines, Guo et al. evaluated degradation and tan δ for three polyurethanes based on heteroaromatic, polyester, and polyether polyols, finding significant changes in degradability and transition temperature depending on the chemical structure of the polyol [11]. In medium-density foams, other authors reported a 46°C increase in Tg upon substituting petrochemical polyol with epoxidized soy oil [12].
Several modifications have been proposed to increase the reactivity of biopolyol-based systems [13]. Epoxidation has attracted significant interest due to its low process cost [14]. The lower reactivity of biopolyols stems mainly from two factors: the number of hydroxyl groups and their reduced accessibility for reaction with isocyanate groups. The latter is limited both by molecular weight [15] and by the position of hydroxyl groups within the macromolecular structure, which is particularly relevant in polyurethane synthesis given the short processing window available for these materials. For epoxidized soy oil, other authors have reported hydroxyl numbers around 250 mg KOH/g⁻¹ and molecular weights between 954 and 1300 [16,17,18]. Accordingly, the system under study is expected to exhibit lower reactivity compared to the petrochemical system. One possible solution is to use mixtures of biobased polyols with glycerin, which introduces three significant effects: its low molecular weight and three hydroxyl groups make it an excellent crosslinker [19], thereby increasing the crosslinking density of the system. However, glycerin competes with the biopolyol for reaction with isocyanate, and its low molecular weight also acts as a plasticizer. In the case of linseed-oil-based polyurethane, the authors reported an increase in the stability of the elastic modulus along with two transitions at 60°C and 175°C [20]. The glycerin content has a significant influence on the thermomechanical properties of rigid polyurethanes; however, higher glycerin additions substantially increase the required isocyanate index, consequently reducing the proportion of the base polyol that can be replaced. For mixtures of hydrolyzed soy oil and glycerin at 20% and 30%, tested with different water contents, the authors concluded that density was the most significant factor affected, and that glycerin contents above 20% did not produce a significant change in the crosslinking of the system [21].
Once polymerized, polyurethane structures are commonly described in terms of soft and hard segments. Soft segments are associated with the polyol backbone, particularly in the case of long-chain polyols. Hard segments, on the other hand, are related to urethane, urea, biuret, and isocyanurate groups. The distribution of soft and hard segments is closely linked to phase separation behavior. In particular, the addition of glycerin to soy-based polyurethane thermosets has been associated with the formation of two interpenetrating polymer networks, attributed to the dissimilar molecular weights and reactivities of the polyol mixture components, as well as their competition for reaction with residual isocyanate groups [22]. The formation of these polymer networks has been linked to significant changes in thermomechanical properties, with domain distribution playing a major role [23]. Segment separation has been widely reported in the literature and is particularly relevant for tailoring thermomechanical and gas-barrier properties. Synthesis temperature strongly influences segment separation, with higher temperatures generally associated with poorer segment distribution [24]. Furthermore, microphase separation has been associated with a significant increase in Tg, and long-chain, high-molecular-weight polyols show a marked tendency to form such microstructures [25].
This work focuses on the use of mixtures of epoxidized soybean oil and glycerin as an alternative core material for sandwich panels. To the best of our knowledge, no other studies have reported similar materials at densities between 100 kg/m³ and 300 kg/m³. The main precedent is the fabrication of thermosetting prepregs from mixtures of glycerin and epoxidized soy oil, where the heterogeneity of the polyol system promoted the formation of an interpenetrating polymer network. However, no other studies have reported biobased polyurethane foams within this density range. In cellular materials, an additional challenge is the effect of reaction kinetics on the curing cycle. The results presented here show a significant increase in the thermal stability of the storage modulus for the biobased materials relative to the reference petrochemical foams. In addition, irreversible processes activated by thermal cycling were observed and corroborated by DMA and DSC. These results were linked, through FTIR, to a post-cure process and to changes in the heterogeneous structure formed by soft and hard segments. Another notable result, observed by SEM, was the formation of nodular structures in the biobased materials, possibly associated with changes in the microstructural distribution of the domains.

Materials and Methods

Material

The biobased formulations were prepared using OHE290P as the biopolyol (hydroxyl number of 290 mg KOH·g⁻¹ and functionality of 2.9), which was compared with formulations based on the petrochemical polyol SD441 (Huntsman; hydroxyl number of 440 mg KOH·g⁻¹, functionality of 4.3). Glycerin was added to the biobased formulations to increase the functionality and hydroxyl number of the biopolyol, with the additional effects of crosslinking. The glycerin content was kept constant at 20 parts by weight (pbw) across all biobased formulations. AK8805 was used as the silicone surfactant, and dibutyltin dilaurate (DBTDL) and N,N-dimethylcyclohexylamine (DMCHA) were used as catalysts. The blowing agents were water and Opteon 1100, a hydrofluoroolefin (HFO). A polymeric methylene diphenyl diisocyanate (pMDI), commercially designated Suprasec 5005 (Huntsman), was used in all formulations.

Sample Preparation Methods

All components except the diisocyanate were first mixed together; the pMDI was then added, and the resulting blend was mixed at 2500 rpm for 20 s in a cylindrical mold 70 mm in diameter. Table 1 reports the amount of each component; the surfactant and both catalysts were kept constant across all formulations.
The foams were cured for one week and subsequently cut into cubic specimens with a side length of 25 mm. Specimens were characterized in terms of apparent density, Fourier-transform infrared spectroscopy (FTIR), compression behavior, dynamic mechanical analysis (DMA), thermogravimetric analysis (TGA), and scanning electron microscopy (SEM).
As shown in Table 1, formulations F1–F3 are based on the petrochemical polyol, while F4–F6 are based on soy oil. Both material families span a density range of approximately 100 to 300 kg/m³. For the lowest-density formulations, a physical blowing agent (HFO) was used, whereas water was used as the blowing agent for the remaining formulations. This substitution was necessary because, in high-density foams, the amount of physical blowing agent required is very small, which makes accurate dosing difficult.

Characterization Methods

Density, compression strength and SEM. Apparent density was measured on ten cubic specimens (25 mm side) following ASTM D1622. Specimens were cut from the interior of each foam block, avoiding the outer skin, to prevent the effect of surface densification. The same specimens were subsequently used to determine compressive strength and elastic modulus. Prior to testing, all specimens were post-cured at 70°C for 5 h, following the standard post-cure conditions typically applied to polyurethanes. Compression tests were performed in accordance with ASTM D1621 using an Instron 5985 universal testing machine equipped with a 100 kN load cell. Five specimens were tested in the foam-rise direction (hereafter, longitudinal, L) and five in the perpendicular direction (hereafter, transverse, T), at a crosshead speed of 2.5 mm/min. Compressive strength (Rc) was defined as the stress at 10% deformation, and the elastic modulus (M) was calculated from the slope of the linear (elastic) region of the stress–strain curve. The anisotropy factor (AF) was calculated as the ratio between the longitudinal and transverse compressive strength. Compressive strength values were also normalized by apparent density to obtain specific strength. All results are reported in Table 2. The cellular structure of the materials was examined by scanning electron microscopy (SEM). Samples were gold-coated prior to imaging, and micrographs were acquired at 600× magnification. Cell dimensions were measured using ImageJ software.
DMA. Rectangular specimens (5 mm thickness × 9 mm width × 50 mm length) were tested in dual-cantilever configuration using a PerkinElmer DMA 8000. Specimens both with and without the post-cure cycle (70°C for 5 h) were evaluated. After confirming that the material response remained within its linear viscoelastic range, the following test conditions were selected: frequency of 0.7 Hz, amplitude of 0.03 mm, heating rate of 2°C/min, and a temperature range of 20–220°C. A comparison between the post-cured and non-post-cured specimens is provided in the Supporting Information. For the non-post-cured specimens, two consecutive heating runs were performed under identical conditions to assess the effect of thermal cycling. Figure 2a,b show the storage modulus (E′) and damping factor (tan δ), respectively.
FTIR. FTIR pellets were prepared with an approximate composition of 0.8 mg of sample and 114 mg of potassium bromide (KBr); five pellets per formulation were tested to obtain the average values reported in Table 3. To ensure homogeneous mixing, the material was grated and ground in an agate mortar before KBr was added; the mixture was then ground again and pressed into a pellet using a Pike hydraulic press at 7 t for 4 min. Spectra were collected using a Shimadzu IRAffinity-1 FTIR spectrometer with 63 scans, a resolution of 2 cm⁻¹, and a spectral range of 600–4000 cm⁻¹. The six as-prepared formulations were analyzed, together with samples of F3 and F6 taken after the two DMA heating runs (denoted F3DMA and F6DMA). The areas of the bands centered at 1525, 1710, 2270, and 1410 cm⁻¹ were normalized to the phenyl absorption band (1590 cm⁻¹). Representative spectra are shown in Figure 3, and the corresponding band areas are reported in Table 3.
DSC. DSC measurements were performed on samples of approximately 3 mg using a Shimadzu DSC-60, with a heating rate of 10°C/min up to a maximum temperature of 300°C. Figure 4 shows the thermal evolution of the six materials over two consecutive heating runs, separated by a cooling cycle; the labels XX-P1 and XX-P2 denote the first and second runs, respectively, for material XX.
TGA. TGA measurements were performed on a Shimadzu TGA-60 (Japan) under an air flow of 100 mL/min and a nitrogen purge flow of 30 mL/min, at a heating rate of 10°C/min over a temperature range of 20–700°C. Sample weights were approximately 4.5 mg in all cases. Results were normalized to the initial sample weight and are presented together with the derivative thermogravimetric (DTG) profiles in Figure 5.

Results and Discussion

Apparent Density

The apparent density of the formulations ranged between 100.2 kg/m³ and 327.3 kg/m³. The increase in apparent density was achieved by reducing the amount of blowing agent added. Notably, for the biobased materials, the amount of blowing agent required was approximately twice that used for the petrochemical formulations. This result is noteworthy because, in low-density foams, high polyol functionality is generally associated with an increase in foam density, an effect attributed to the predominance of the gelation reaction over the blowing reaction [26]. In the case of the high-density materials studied here, the opposite trend was observed, which may be associated with the lower hydroxyl number of the biopolyol and the consequently lower amount of reaction heat potentially available during foaming. Another possible contributing factor is the higher molecular weight and molecular architecture of the biopolyol, which may reduce the accessibility of hydroxyl groups and consequently influence the reaction kinetics.
To interpret these results, it is useful to consider the foaming process, which is governed by two exothermic reactions: gelation and blowing. The gelation reaction occurs between isocyanate and polyol, producing urethane. The blowing reaction occurs between isocyanate and water, first forming carbamic acid, which subsequently decomposes into an amine and carbon dioxide (CO₂) [4]. For formulations F1 and F4, the blowing reaction does not occur, since no water is present; in these cases, the physical blowing agent evaporates as a result of the heat released by the gelation reaction alone. Consequently, when the system has a low hydroxyl content, less heat is available to evaporate the physical blowing agent, reducing its efficiency. For formulations F2, F5, and F6, only a small amount of water was intentionally added. However, the longer gelation times observed for the biobased systems may have allowed a fraction of the generated CO₂ to escape before sufficient network development occurred to retain it within the cellular structure. This behavior suggests that the balance between gelation and blowing kinetics was not optimal in the biobased formulations.
Finally, no physical or chemical blowing agent was intentionally added to F3; nevertheless, cellular expansion was observed. One possible explanation is that residual moisture provided a source of CO₂, while entrapped air may have acted as nucleation sites for cell formation. This hypothesis is further evaluated based on the cellular morphology observed by SEM and discussed in the corresponding section.

Cell Geometry and Morphology Using SEM

Figure 1 shows the cellular microstructure of the materials, and Table 3 reports the corresponding cell dimensions in the longitudinal and transverse directions. The image illustrates the cellular morphology of all materials. For the petrochemical materials, longitudinal cell size ranged from approximately 202 µm to 130 µm with increasing density. For the biobased materials, longitudinal cell size was 14–33% larger than for the corresponding petrochemical formulations. In the transverse direction, differences between the two material families were smaller (12–18%); as a result, the anisotropy factor was higher for the biobased materials. A similar increase in anisotropy has been reported for foams produced with low-functionality polyols [5,14]. In all cases, cell size decreased with increasing foam density. The decrease in cell size with increasing density is consistent with the greater amount of solid material per unit volume in the denser foams.
For formulations with densities of 200 and 300 kg/m³, the biobased materials show a marked change in cell morphology, with the cells no longer appearing hollow. To the best of our knowledge, no previous studies have characterized biobased polyurethane foams from mixtures of soybean oil and glycerin within this density range. A qualitatively similar heterogeneous morphology has been reported in polyurethane foams containing MnO₂ fibers, where the authors associated the observed features with filler agglomeration [29]. Although no filler was present in the present formulations, the morphological similarity suggests that local differences in polymer accumulation and/or reaction–foaming behavior may contribute to the observed structure. These observations indicate that increasing density had a more pronounced effect on the cellular morphology of the biobased foams. The larger cells observed in the biobased foams may be related to their higher blowing-agent requirements, although differences in reaction kinetics and network development may also contribute.

Compressive Mechanical Tests

For the petrochemical foams, the longitudinal compressive strength increased from 1.03 MPa to 8.92 MPa across the density range studied, a trend consistent with values reported for PVC panels of similar density [27]. For the biobased materials, the same parameter increased from 0.60 MPa to 4.07 MPa. When comparing materials of similar density, the biobased formulations showed reductions in compressive strength of 58.2%, 37.0%, and 45.6% relative to the petrochemical materials at approximately 100, 200, and 300 kg/m³, respectively. A similar trend was observed in the transverse direction, with values ranging from 0.68 MPa to 7.38 MPa for the petrochemical materials and from 0.43 MPa to 2.57 MPa for the biobased materials. Reported literature values for compressive strength include 0.2–0.7 MPa for cork cores (density 140–250 kg/m³), 2.4 MPa for polylactic acid (PLA) foam cores (320 kg/m³), and 5.5–22 MPa for balsa wood (density 100–280 kg/m³) [1]. The biobased materials characterized in this work showed compressive strengths above the reported values for cork and PLA foam cores at comparable densities, although they remained below the values reported for balsa wood.
The elastic modulus followed a similar trend: in the longitudinal direction, values ranged from 23.65 MPa to 164.02 MPa for the petrochemical materials (100–300 kg/m³) and from 11.18 MPa to 77.12 MPa for the biobased materials over the same density range. In all cases, both compressive strength and elastic modulus increased substantially as density tripled, by approximately 7–9 times depending on formulation and loading direction. Simplified models predict that longitudinal wave speed scales with the square root of the ratio between elastic modulus and density [2]; the results obtained here follow this trend, with both foam families falling within the expected range predicted by the model. The differences in density-dependent behavior between the two polyol systems may be associated with differences in the functionality, hydroxyl number, and molecular architecture of the polyols and changes in the morphology. Lim et al. reported changes of only about 20% in mechanical properties for foams of approximately 100 kg/m³ formulated with polyols having functionalities between 3 and 6 [28]. The substantially larger differences observed between the two foam families in the present study therefore suggest that functionality alone cannot account for the observed mechanical behavior. Other differences in polyol characteristics, particularly molecular weight and molecular architecture, may also contribute.
Within the density range investigated, changes in the blowing-agent system were not associated with systematic changes in compressive strength or elastic modulus. This suggests that, once density was controlled, polyol type had a greater influence on the mechanical response than the specific blowing-agent conditions used to achieve the target densities. The lower compressive strength and elastic modulus of the biobased foams may also be related to the differences in cellular morphology observed by SEM. In particular, the larger cell size, greater anisotropy, and heterogeneous structures observed in the biobased materials may reduce the efficiency of the cellular network in supporting and transferring applied loads.

Dynamic Mechanical Analysis (DMA)

Dynamic mechanical analysis was performed on materials without post-curing treatment. The evolution of the storage modulus (E′) and damping factor (tan δ) with temperature is shown in Figure 2. The initial value of E′ increased with density, following the same trend observed in the compression tests. For petrochemical formulations, the initial storage modulus increased to 2.56x108 Pa for F1, to 3.00x108 Pa for F2 and 6.58x108 Pa for F3. Similarly, for biobased foams, E’ was 8.68x107 Pa for F4 to 1.72x108 Pa for F5 and 2.43x108 Pa for F6, respectively. Consequently, increasing density from F1 to F3 resulted in a 157% increase in the initial E′ for the petrochemical formulations, while an increase of 180% was observed for the biobased foams from F4 to F6. At comparable densities, the initial E′ of the biobased foams was 66%, 43%, and 63% lower than that of the petrochemical foams for the F1/F4, F2/F5, and F3/F6 pairs, respectively. As temperature increased, E’ decreased markedly in both material groups, although the magnitude of the decrease differed between them. At 170ºC, the storage modulus of F1, F2 and F3 decreased to 6%, 8% and 8% of their respective initial values. For the biobased materials, E’ decreased to 32%, 39% and 34% of the initial values for F4, F5 and F6. At the beginning of the second heating run, E′ showed no consistent change relative to the first run at atmospheric temperature. For the petrochemical formulations, the initial E’ decreased by 15% for F1, and 17 % for F2, while it increased by 3% for F3. For the biobased formulations, E’ increased by 22% for F4 and decreased by 5% and 1 %, for F5 and F6, respectively. At 170ºC during the second heating run, E’ decreased to 5%, 4% and 3% of the respective initial values for F1, F2, and F3. In contrast, the biobased formulations retained approximately 50% of their respective initial E’ values.
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Figure 2. First and second heating run of DMA.
Figure 2. First and second heating run of DMA.
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Regarding the damping factor, its initial value increased with density, from 0.0313 for F1 to 0.0467 for F3, corresponding to an approximately 50% increase. On the other hand, for the biobased materials, tan δ decreased to 0.0381 for F4, to 0.0312 for F5 and 0.0318 for F6. The petrochemical-based materials exhibited a maximum tan δ during the first heating run at approximately 170 °C. For F1, the maximum tan δ value was 0.27, whereas for F2 and F3 it was 0.34. For the biobased formulations, tan δ reached approximately 0.25 around 220ºC. In contrast, the tan δ curves of F4, F5, and F6 remained relatively stable up to approximately 170 °C.
The lower initial E′ of the biobased foams is consistent with their lower compressive modulus observed in the mechanical tests, confirming the lower stiffness of these materials under both quasi-static and dynamic loading conditions. Despite the lower initial E′ of the biobased foams, these materials exhibited greater retention of the storage modulus at elevated temperatures. At comparable densities, the biobased foams exhibited 43–66% lower initial E′ than the petrochemical foams. This difference may be associated with differences in polymer network structure arising from the different polyol characteristics. On the other hand, the petrochemical based foams show a strong effect of temperature to the storage modulus, for the same range of temperature biobased foams shown more stable values of E’. At 170ºC, E’ decreased by approximately 93% for the petrochemical formulations and 65% for the biobased formulations. Additionally, in the second heating run, the tan δ response of the petrochemical formulations shifted to a lower temperature, whereas the biobased formulations retained a higher fraction of their initial E′ over the temperature range up to approximately 170ºC. These results indicate that the biobased foams exhibited greater retention of stiffness upon heating, suggesting a lower susceptibility to thermally induced softening over the temperature range investigated. Another interesting result was the change in the range of growth of Tan δ for F4, F5 and F6, the graph showed a decrease of this range in the second heating run, this change suggests the occurrence of irreversible structural changes during the first heating cycle. In contrast, for F1, F2 and F3 there was a decrease in the temperatures associated with the pick. This shift may be related to irreversible changes occurring during the first heating cycle, such as further curing or the onset of thermally induced structural changes. The contrasting behavior of the initial E′ and its temperature dependence suggests that different structural factors may govern the initial stiffness and the resistance to thermally induced softening of the polymer network.

Infrared Absorption Analysis (FTIR)

FTIR spectroscopy was used to investigate the chemical origin of the DMA results. Representative spectra for all materials are shown in Figure 3, and the corresponding band areas, normalized to the phenyl absorption band, are reported in Table 3. As noted previously, F3 and F6 were also analyzed after undergoing the two DMA heating runs. A general challenge in using FTIR to characterize polyurethane chemical structure is spectral overlap; a notable exception is the isocyanate group (–N=C=O), whose absorption band near 2270 cm⁻¹ has no significant overlapping contributions.
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Figure 3. Characteristic FTIR spectra for the materials.
Figure 3. Characteristic FTIR spectra for the materials.
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A further point of interest is the increase in absorbance observed near 2920 cm⁻¹ and 2850 cm⁻¹, bands associated with aliphatic CH₂ groups characteristic of triglycerides [22]. Formulations F4, F5, and F6 showed significantly stronger absorbance in this region compared with F1, F2, and F3, consistent with the triglyceride-based structure of the soy polyol; by contrast, the petrochemical polyol has a lower aliphatic content. Comparison of F3, F3DMA, F6, and F6DMA shows that the area of this band did not change significantly for F3, but increased noticeably for F6. For F6, the residual isocyanate content decreased after the thermal cycle, indicating that additional reactions involving isocyanate groups occurred during heating. In parallel, the area of the aliphatic CH₂ band increased, although this change cannot be directly assigned to a specific reaction mechanism.
Because isocyanurate structures can enhance the thermal stability of polyurethane/polyisocyanurate networks, their possible contribution to the thermomechanical differences was also examined. The band near 1410 cm⁻¹, commonly associated with isocyanurate structures, showed no significant changes among the formulations or after the thermal cycle. Therefore, additional isocyanurate formation does not appear to be a major factor governing the observed differences in thermal behavior.
The carbonyl group (C=O) region of the FTIR spectrum contains multiple overlapping contributions, including free urethane (1715–1730 cm⁻¹), hydrogen-bonded (associated) urethane (1695–1712 cm⁻¹), urea and biuret (1630–1710 cm⁻¹), and ester groups (1735–1750 cm⁻¹) [32]. Table 3 reports the total area of this composite band, including all contributions. However, inspection of the spectra shows that the peak center is shifted toward higher wavenumbers for the biobased materials, which may partly reflect the contribution of ester carbonyl groups present in the soy-based polyol. Following the thermal cycle, the carbonyl band area increased for F3 but decreased for F6.
Finally, the amide II band, centered near 1525 cm⁻¹, includes contributions from urethane, hydrogen bonding, urea, and biuret groups [4]. For the petrochemical materials, this band area ranged from 2,74 to 3,22; the biobased materials showed lower values, between 2,87 and 2,92. The thermal cycle had a notable effect on this band, with an 18% increase observed specifically for the biobased materials. Considering the concurrent decrease in the carbonyl band area for the same materials and the greater retention of E′ observed by DMA, these changes may indicate further reactions involving residual isocyanate groups during heating. One possible mechanism is the reaction of residual isocyanate groups with trace moisture, generating amines that can subsequently react with additional isocyanate groups to form urea structures. Such additional network development could contribute to the greater retention of E′ observed for the biobased formulations during the second heating run.
To sum up this section, it is important to set the significant contribution of aliphatic groups in the biobased foams. On the other hand, for F6, the heating run decreases the area associated with the isocyanate and carbonyl groups and increases the area of amide II. Although the amide II band contains contributions from several chemical groups, its increase, together with the previous results in the experiments of the greater retention of storage modulus observed during the second DMA run. Taken together, these changes are consistent with further network development during the first heating cycle, suggesting that a post-curing process may occur in the soy-based foams. For the petrochemical foam, the carbonyl band area decreased after the thermal cycle; however, because of the substantial overlap of the contributing carbonyl species, the chemical origin of this change cannot be determined from the FTIR results alone.

Differential Scanning Calorimetry (DSC)

Differential scanning calorimetry (DSC) was used to identify thermal events associated with residual reactions and post-curing in the polyurethane materials. The thermal evolution of all formulations is shown in Figure 4. No appreciable differences were observed in the temperature range or overall shape of the exothermic events as a function of density, despite the threefold variation in this parameter across the formulations studied. However, two distinct behaviors were observed depending on the origin of the polyol, as detailed below. For the petrochemical foams, two exothermic events were observed: the first between 60°C and 170°C, and the second between 180°C and 230°C. For the biobased materials, two events were likewise observed, but at higher temperatures: the first between 60°C and 245°C, and the second between 255°C and 290°C.
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Figure 4. Results for two consecutive DSC experiments.
Figure 4. Results for two consecutive DSC experiments.
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The broad first event observed in the biobased materials may reflect the presence of chemically heterogeneous reactive environments associated with the combination of the higher-molecular-weight soy-oil polyol and glycerin, which differ substantially in molecular size and functionality. For the petrochemical foams, the presence of two distinct thermal events may indicate the occurrence of reactions involving chemical environments with different thermal reactivities. This behavior may be related to differences in the reactivity and network development associated with the high-functionality petrochemical polyol.
In the second heating run, the exothermic events were strongly reduced in all formulations, indicating that the first run induced irreversible changes in the samples. This is consistent with a post-curing process, particularly for the biobased materials, for which the upper-temperature region of the first exothermic event overlaps with the tan δ transition observed by DMA. This interpretation is further supported by the decrease in residual isocyanate content measured by FTIR. For the petrochemical materials, the first event likewise overlaps with the DMA tan δ transition. However, because no detectable residual isocyanate was observed and the tan δ peak shifted to a lower temperature during the second DMA run, the observed changes are more likely associated with irreversible structural rearrangements than with residual curing. Finally, the second thermal event observed for the biobased materials may be associated with secondary reactions, such as further urea, biuret, or allophanate formation, or with the onset of thermal degradation; these hypotheses are examined further using TGA results, presented in the following section.

TGA in Air Atmosphere

Figure 5 illustrates the mass-loss and derivative thermogravimetric (DTG) profiles of the materials analyzed. TGA was performed under an air atmosphere to facilitate comparison with the DMA results. Formulations F1–F3 exhibited two degradation stages, whereas F4–F6 exhibited three, with no significant effect of density on the overall degradation behavior. For the biobased materials, the first two degradation stages occurred within the same temperature range as the single first stage observed for the petrochemical materials. Onset degradation temperatures, calculated using the tangent method, together with the temperatures corresponding to 5%, 10%, and 50% weight loss, are reported in Table 4. In general, no significant differences were observed among formulations in the temperatures corresponding to the same weight loss, with the exception of the temperature at 50% weight loss. This parameter may reflect differences in the relative contribution of the degradation events occurring at lower temperatures, particularly in the biobased materials, as discussed below. As observed by DSC, density had no significant effect on thermal degradation behavior, whereas polyol origin had a pronounced influence. An additional effect of water content was observed, as the temperatures corresponding to a given weight loss generally increased with increasing water addition. This trend was particularly consistent for the petrochemical formulations, whereas it was observed only for some of the biobased materials. The increase in degradation temperatures with water content may be associated with the formation of urea structures through the reaction of water with isocyanate groups, which can increase the contribution of hard segments and consequently enhance the thermal resistance of the polyurethane network. The less consistent behavior observed for the biobased formulations may be related to differences in network development and isocyanate conversion, as indicated by the residual isocyanate detected by FTIR.
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Figure 5. Thermogravimetric profiles and DTG.
Figure 5. Thermogravimetric profiles and DTG.
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The DTG profiles show a single degradation peak between 280°C and 400°C for the petrochemical materials, with a shoulder near 320°C. Over the same temperature range, the biobased materials show two distinct peaks, centered at approximately 315°C and 365°C. At higher temperatures, all materials show a peak beginning at 485°C and ending at 650°C. Similar degradation stages have been reported for soy- and palm-oil-based foams by Mohammed et al.31 The presence of two distinct peaks in the intermediate temperature range for the biobased materials may reflect the degradation of chemically distinct domains or structures within the polyurethane network. This behavior is consistent with the broader tan δ response observed by DMA and may indicate greater molecular heterogeneity in the biobased materials. In contrast, the single broad event observed for the petrochemical materials may represent overlapping degradation processes that are more clearly separated in the biobased materials. The final, high-temperature degradation stage may be associated with the degradation of MDI-derived structures, consistent with the similar behavior observed for both material families in this temperature range.
The onset degradation temperature for F4, F5 and F6 provide further support for the interpretation of the second event in DSC experiments for biobased materials. In particular, the range of the first stage of degradation was an overlap with the values of the second event for the experiments of DSC; these results upgrade the hypothesis of the occurring degradation process in the first run of DSC experiments.

Conclusions

Rigid polyurethane foams with densities ranging from approximately 100 to 300 kg/m³ were produced with mixtures of soybean oil and glycerin were formulated and compared with a commercial petrochemical polyether polyol. The biobased materials exhibited compressive strength and elastic modulus values substantially lower than those of the petrochemical materials, with compressive strength reductions of approximately 37–58% at comparable densities. The difference was associated with the combined effects of polyol functionality, molecular architecture, and the distinct cellular morphology observed by SEM.
In contrast to their lower mechanical performance at room temperature, the biobased foams exhibited greater retention of the storage modulus at elevated temperatures. The combined DSC, FTIR, and TGA results indicate that this behavior is associated with differences in chemical structure and thermal evolution between the two material families. The biobased foams showed aliphatic structures associated with the triglyceride-based structure of the soybean-oil polyol, as well as evidence of residual reactive groups and further network development during the first heating cycle. The reduction in residual isocyanate content and the changes observed in the carbonyl and amide II regions after heating are consistent with additional reactions during the thermal cycle, supporting the occurrence of a post-curing process in the biobased materials. The higher tan δ peak temperature and broader damping response further suggest a more heterogeneous distribution of polymer environments, which may be related to the molecular architecture of the polyol mixture.
The thermogravimetric analysis also revealed that water addition influenced the degradation behavior of the formulations. In particular, the temperatures corresponding to a given weight loss generally increased with increasing water content, with this trend being more consistent in the petrochemical materials. This behavior may be related to increased formation of urea-containing hard domains through the isocyanate–water reaction and the subsequent reaction of the generated amines with isocyanate groups. The less consistent response observed in the biobased materials, together with their higher residual isocyanate content, suggests that water addition did not result in the same extent of network development under the processing conditions employed.
Overall, the results demonstrate that the replacement of the petrochemical polyol with a mixture of soy-oil and glycerine to polyol produces a distinct balance between mechanical and thermomechanical performance. Although the biobased foams exhibited lower compressive properties, their greater retention of storage modulus at elevated temperatures and their distinct thermal response indicate potential advantages for applications where high-temperature dimensional or mechanical stability is important. The results also demonstrate that polyol molecular architecture, reaction kinetics, and cellular morphology should be considered together when designing high-density biobased polyurethane foams.

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Figure 1. Cell structure for the different materials.
Figure 1. Cell structure for the different materials.
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Table 1. Composition of the formulations tested in this work.
Table 1. Composition of the formulations tested in this work.
Formulation pMDI (pbw) SD441 (pbw) AK8805 (pbW) DMCHA (pbW) DBTDL (pbW) OHE290P(pbW) Glycerin H2O (pbW) Opteon 1100 (pbW)
F1 111.6 100 2 0.5 0.5 0 0 0 2.5
F2 115.5 100 2 0.5 0.5 0 0 0.25 0
F3 111.6 100 2 0.5 0.5 0 0 0 0
F4 126 0 2 0.5 0.5 100 20 0 5
F5 133.9 0 2 0.5 0.5 100 20 0.5 0
F6 130 0 2 0.5 0.5 100 20 0.25 0
Table 2. Physical properties of the foams. Part 1.
Table 2. Physical properties of the foams. Part 1.
Formulation Apparent density (kg/m3) RcL (MPa) RcT (MPa) RcspL (MPa.( Kg/m3)-1) RcspT (MPa.( Kg/m3)-1) AF ML (MPa) MT (MPa)
F1 100.3 +/- 1.22 1.03 +/- 0.20 0.68 +/- 0.09 10.28.10-3 +/- 1.77.10-3 6.87.10-3 +/- 0.95.10-3 1.51 +/- 8.32.10-2 23.65 +/- 6.66 11.73 +/- 2.08
F2 212.9 +/- 1.16 3.84 +/- 0.10 2.93 +/- 0.06 18.04.10-3 +/- 0.43.10-3 13.9.10-3 +/- 0.38.10-3 1.31 +/- 0.62.10-2 65.30 +/- 15.92 42.50 +/- 2.56
F3 327.3 +/- 13.18 8.92 +/- 0.49 7.38 +/- 0.90 27.68.10-3 +/- 1.31.10-3 22.2.10-3 +/- 1.63.10-3 1.21 +/- 7.99.10-2 164.02 +/- 10.77 128.14 +/- 31.65
F4 116.4 +/- 3.70 0.60 +/- 0.04 0.43 +/- 0.08 5.31.10-3 +/- 0.42.10-3 3.68.10-3 +/- 0.58.10-3 1.39 +/- 15.10-2 11.18 +/- 2.14 6.87 +/- 1.29
F5 190.8 +/- 3.99 1.42 +/- 0.07 0.91 +/- 0.05 7.5.10-3 +/- 0.18.10-3 4.73.10-3 +/- 0.13.10-3 1.56 +/- 0.88.10-2 27.77 +/- 2.18 14.25 +/- 0.96
F6 314.3 +/- 6.99 4.07 +/- 016 2.57 +/- 0.06 13.03.10-3 +/- 0.28.10-3 8.12.10-3 +/- 0.13.10-3 1.58 +/- 2.77.10-2 77.12 +/- 3.22 37.63 +/- 2.88
Table 3. Physical properties of the foams. Part 2.
Table 3. Physical properties of the foams. Part 2.
Formulation C=O normalized area C-H normalized area Amida II normalized area -N=C=O normalized area Cell size L (um) Cell size T (um)
F1 4.14 +/- 0.18 0.36 +/- 0.04 2.74 +/- 0.22 0.33 +/- 0.14 202.25+/-15.95 113.99+/-5.25
F2 4.17 +/- 0.26 0.37 +/- 0.02 3.03 +/- 0.14 0.24 +/- 0.12 171.29+/-8.03 117.39+/-3.97
F3 4.48 +/- 0.15 0.40 +/- 0.01 3.22 +/- 0.26 0.16 +/- 0.03 129.91+/-35.12 95.28+/-10.36
F4 5.10 +/- 0.07 0.36 +/- 0.02 2.92 +/- 0.15 0.82 +/- 0.20 229.95+/-10.41 134.41+/-3.18
F5 4.93 +/- 0.42 0.37 +/- 0.05 2.89 +/- 0.20 0.88 +/- 0.41 190.15+/-19.38 118.13+/-8.23
F6 4.78 +/- 0.66 0.35 +/- 0.04 2.87 +/- 0.25 0.67 +/- 0.11 172.24+/-9.46 107.08+/-8.39
F3DMA 5.42 +/- 0.52 0.45 +/- 0.08 3.39 +/- 0.13 0 - -
F6DMA 3.69 +/- 0.11 0.38 +/- 0.05 3.41 +/- 0.27 0 - -
Table 4. Characteristic temperatures in the termogravimetric test.
Table 4. Characteristic temperatures in the termogravimetric test.
Formulation TOnset T (5%) T (10%) T (50%)
F1 285.54 281.64 302.78 380.60
F2 296.09 299.93 315.01 421.71
F3 293.82 294.26 311.41 413.29
F4 282.22 294.06 309.17 509.26
F5 287.17 301.04 314.90 488.27
F6 284.41 300.38 314.41 498.54
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