Introduction

Isotactic poly(1-butene) (PB-1) is a thermoplastic polyolefin with remarkable properties, distinguishing it from ordinary polyolefins, such as isotactic polypropylene or polyethylene. Compared to other polyolefins, PB-1 exhibits low stiffness, high creep resistance, good elastic recovery, and exceptional chemical resistance, e.g., to oils, fats, acids, alcohols, and ketones. Equally important is its high heat distortion temperature [1]. The physical, chemical, and mechanical properties are influenced by the given crystalline phase.

Despite its excellent properties, the application of PB-1 compared to other polyolefins is limited by conformational polymorphism and dimensional instability, reaching up to 4%, which occurs during the transition between the initial metaphase II and stable phase I. The individual phases exhibit different densities [2,3,4]. Four crystalline phases of PB-1 have been characterized thus far. Crystallization from the melt follows Ostwald's rules of stages [5]. Initially, the kinetically favored unstable tetragonal mesophase II is formed from the melt, transforming into the hexagonal phase I within a few days. During the transition between phases II and I, phase I', which has the same crystal structure as phase I but differs in lower melting temperature, can be formed, probably caused by more conformational defects present. The unstable phase III has an orthorhombic structure and is obtained by solvent evaporation or precipitation, followed by transformation to phase I under certain conditions [6].

Important factors influencing the aging process of polymers include aging time, applied radiation (solar or fluorescent light sources), the presence of oxygen, ozone (O3), applied heat, humidity, rain, impurities in the air, and others. The chemical structure of the polymer chains and morphology determine its resistance to chain degradation. Stable structures with higher density and crystalline regions showed an enhanced resistance to the diffusion of penetrating substances and radiation [7].

However, the course of degradation of PB-1 still needs to be fully understood. This process is assumed to be similar to the degradation of isotactic polypropylene due to identical presence of tertiary carbon in the main chain [8,9,10]. Craig et al. [11] investigated the oxidation of isotactic polypropylene at room temperature; the polymer was in the solid state and well below its crystallization temperature. Changes occurred primarily in the amorphous region; oxygen can diffuse freely through this area. In contrast, it hardly passes through the crystalline regions. The main changes were (i) chain scission, (ii) cross-linking, and (iii) formation of molecular defects, e.g., formation of carbonyl groups.

During chain scission, the chain segments that were previously entangled were released, and crystallization might occur. If the polymer was formed from many free segments, the formation of new crystals could be assumed, but it is more likely that the segments formed in this way would join the existing growth surfaces. When the crystallinity of the material increases, secondary crystallization occurs, known as the “chemi-crystallization” process [12, 13].

Cross-linking inhibits further crystallization of chain segments connected by networks. Molecular defects such as formed carbonyl groups do not fit into the crystal lattice, and parts of the molecular segments containing such defects cannot participate in secondary crystallization [11, 14, 15].

During the so-called “chemi-crystallization”, the amount of low-density amorphous phase increases compared to the crystalline phase; shrinkage of the material occurs. Chain scission (and subsequently secondary crystallization and shrinkage) tends to be highly dependent on the distance from the UV source resulting in residual stresses on the material's surface [16]. Residual and “frozen stresses” arising during processing can lead to material cracking. Fine cracks that are not deep and do not lead to total failure of the material can, among other things, impair the appearance of the material, for example, its gloss.

A material affected by molecular degradation and subsequent loss of mechanical integrity provided by the molecular network exhibits reduced flexibility resulting from increased crystallinity and additional cross-linking [17, 18]. The degradation conditions also influence the degradation behavior [19]. The general factors that can influence or even limit the degradation of polymers are divided into physicochemical conditions and material properties.

Material properties play a significant role, among them thickness of material [20], polymer crystallinity [21], molecular weight [22, 23], or polymer composition [24, 25]. The formation of peroxide substances in the main chain of a polymer is encountered as the primary step of oxidative degradation. Hydroperoxide groups are essential for inducing thermal oxidation and photo-oxidation of polymers, but they are also important intermediates in overall oxidation reactions. Critical physicochemical conditions determined by the environment, include relative humidity [26, 27], pH value [28, 29], temperature [30, 31], amount of UV and visible irradiation [32, 33], or availability of oxygen [34, 35]; therefore, different results in this work would be expected.

Understanding the degradation mechanisms and the impact of degradation conditions on the properties of PB-1 is crucial due to its widespread application in various fields such as geothermal piping, easy-peel films, hot-water systems, heating systems, cables, and wires. The presented work focused on a multi-stage study of the degraded PB-1 extrudates in two degradation modes: natural and artificial aging. The chemistry of the degradation and, subsequently, the degradation rate of the extrudates were studied. The acceleration time equivalent was used to determine the degradation rate between naturally and artificially aged extrudates. The effects of the degradation on the mechanical properties, namely tensile strength, tensile modulus, and elongation-at-break values were studied and compared to the development of the surface cracks.

Experimental

Materials and methods

Isotactic poly(1-butene) (PB-1) DP 0401 M was obtained from LyondellBasell (Louvain la Neuve, Belgium). According to information from the manufacturer, it was a semi-crystalline homopolymer with a balance between creep resistance and formation of cracks in an environment with increased temperature. The basic properties were a density of 0.915 g.cm−3 (ISO 1183) and a melt flow index of 15 g/10 min (ISO 1133).

The material was extruded on a Brabender Plasti-Corder PL2000 single-screw extruder, and the process conditions were set as follows: screw compression ratio 1:4; screw speed 20 rpm; barrel temperatures of 135 °C, 140 °C, 145 °C, and 150 °C; extrusion slit die with profile 2 × 20 mm2 with temperature 150 °C.

Characterization

For tensile testing the dog-bone-shaped plates (ISO 527–3 type 5A) with a total length of 110 mm, weight of 3 g and the length of working distance 30 mm, width 6 mm and thickness 2 mm, naturally colored (no colorants were used) were cut from the extruded samples.

Artificial aging was performed in a SEPAP 12/24 MPC device (ATLAS company) equipped with four 400 W medium-pressure mercury arc lamps in borosilicate envelops that emit discrete radiation at 290, 313, 365, 405, 436, 547, and 579 nm, at a temperature of 60 °C, and at different irradiation time intervals (0, 10, 20, 40, and 60 h). A single sample, with a length of 50 mm and a weight of 1.5 g, along with six dog-bone-shaped samples for each specific exposure duration, were subjected to both artificial aging and natural weathering processes. One sample length of 50 mm, weight of 1.5 g, and six dog-bone-shaped samples for each individual exposition time were exhibited to artificial aging resp. natural weathering. Testing specimens were in the middle of the test chamber mounted on carousel.

The study of natural weathering was conducted in Nivnice, Czech Republic, which is situated at an altitude of 250 m with geographical coordinates 48°58′42″N latitude and 17°38′47″E longitude. During the period extending from July to September, specimens in a dog-bone shape were oriented toward the southern direction at an inclination of 45°. This setup facilitated the exposure of the samples to UV light exposure for a cumulative duration reaching up to 102 days. Meteorological data affecting the degradation process of naturally aged samples are given in Table 1.

Table 1 Meteorological data affecting the naturally aged samples
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A Nicolet 380 FTIR spectrometer, equipped with a SPECAC ATR Diamond-Germanium crystal, with a resolution of 4 cm−1 and measuring 32 scans, was used for surface analysis of the degraded samples. Molecular degradation was characterized by the carbonyl index \(CI\) as follows:

$$CI ={ A}_{C}/{ A}_{R},$$

where \({ A}_{C}\) is the area of carbonyl absorption peaks (in the range of 1504–1816 cm−1), and \({A}_{R}\) is the area of reference peaks unaffected by photo-oxidation and different degrees of crystallinity (in the range from 1400 to 1500 cm−1) [17, 36].

The yellow index YI E313 was measured using an UltraScan Pro spectrophotometer from Hunterlab.

For wide-angle X-ray scattering in reflection mode, a Philips X’pert (PANalytical, Netherlands) was used with an interval of 2θ = 5–30°, a shift angle of 0.05°, and a duration 0.5 s. A monochromatic CuKα anode beam with a wavelength of λ = 0.154 nm and a Ni filter were used. The tetragonal phase of PB-1 (phase II) shows the following diffraction peaks at 11.9°, 16.9°, and 18.5° 2θ, corresponding to the (200), (220), and (213) planes, respectively. Hexagonal phase I is characterized by four distinct diffraction peaks of 10.0°, 17.5°, and 20.4° 2θ, originating from the (110), (300), (220 + 211) planes, respectively. It should be noted that due to the phase transition with increasing time, the intensity of the peaks typical of hexagonal phase I (i.e., (110) at 10.0° 2θ) increased, while the intensity of the peaks related to tetragonal phase II (i.e., (200) at 11.9° 2θ) decreased [37].

PEAKFIT v4 was used to evaluate crystallinity. From the point of view of phase transformation, it was more appropriate to observe the disappearance of the peak 2θ = 11.9° of phase II (reflection plane 200) rather than the increase of the peak in the region of 2θ = 10.0° of form I (reflection plane 110). A post-crystallization phenomenon occurred with the phase transformation, during which, part of the amorphous polymer crystallized into phase I [38,39,40]. The development of the peaks of the individual phases was observed (2θ angles 10.0° and 11.9°), and the ratio of the peak heights and the content of the given phase to the sum of both heights were subsequently calculated [41].

A Zwick 145,665 multi-purpose apparatus was used for the tensile test according to standard EN ISO 527–3. The specimens were stretched at room temperature until rupture at a test speed of 50 mm/min. For each measurement set, 5 specimens were used. The stress and strain values derived strength-at-break and elongation-at-break. In addition, the elastic modulus was evaluated using a Zwick external extensometer (gauge length of 20 mm) at a test speed of 1 mm/min (ISO).

The surface morphology of the degraded samples was examined using a scanning electron microscope (SEM) with an accelerating voltage of 10 kV (Phenom Pro desktop scanning electron microscope with the BSE detector, Phenom-World B.V., The Netherlands). No treatment of degraded samples has been applied.

Results and discussion

Changes in chemical composition

FTIR–ATR can identify the specific functional groups present in the degraded polymer, providing information about the degradation mechanism and the types of reactions. In this work, the carbonyl groups in the degraded polymer sample can indicate that the sample underwent oxidative degradation. The number of the presented carbonyl groups can be used to quantify the changes in the chemical structure of the polymer due to the degradation [17].

A comparison of the evolution of the carbonyl index and, thus, the chemical changes evoked by the degradation is shown in Fig. 1. In Fig. 1a, a, visual representation of the evolution of the carbonyl groups for natural weathering can be seen. In Fig. 1b, the evolution of the carbonyl groups for artificially aged samples is shown. Each curve represents a characteristic aging time 0, 24, 37, 58, 72, 88, and 102 days of natural weathering and 0, 10, 20, 40, and 60 h of artificial aging. In the artificially aged and naturally weathered samples, a gradual increase in the characteristic peak can be observed in the absorption region of the carbonyl groups at 1700–1750 cm−1. An increase can also be observed in the region at 1645 cm−1, which is attributed to vinyl groups. In the naturally weathered samples in the region of carbonyl groups, the most significant increase was observed after only 24 days of exposure. In all the naturally weathered samples, the significant formation of carboxylic acids was observed (1712 cm−1).

Fig. 1
figure 1

FTIR–ATR spectra of: a naturally weathered and b artificially aged PB-1 samples

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A comparison of the properties of PB-1 in this work was based on the interpolation of the development of the carbonyl index values during artificial and natural weathering and the subsequent recalculation of the accelerated aging time according to the corresponding carbonyl index values to the equivalent time during natural weathering. This approach was based only on the development of the carbonyl index, which was primarily affected by photo-oxidation [10, 42]. Therefore, the contribution of the influence of humidity and temperature fluctuations was neglected, however, these fluctuations influenced the development of the material's mechanical properties. Figure 2 shows the evolution of the carbonyl index as a function of the exposure time concerning the accelerated time equivalent.

Fig. 2
figure 2

Comparison of the evolution of the carbonyl index of the artificially aged and naturally weathered PB-1 samples with the application of accelerated time equivalent

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The increase in carbonyl products was slower due to washing out of the sample surface due to weather conditions. On the other hand, in artificially irradiated samples, the content of chemical groups essentially increased throughout the exposure period, which may also lead to the erosion of the crystallites and possible breakdown of the material.

In Fig. 3, the development of yellow coloring during degradation can be observed. For both series, the trend was similar to the measured carbonyl index (Fig. 2). Based on this comparison, the yellow index (in the case of undyed polymers) indicated the degree of degradation [43]. The yellow index was higher in naturally weathered samples due to the presence of external influences not present during artificial aging, e.g., a higher amount of oxygen, rain, and changeable surrounding temperature, which resulted in different chemical degradation processes and formation of different chemical compounds; these conditions cannot be precisely defined due to the impossibility of deconvoluting the peaks in the range of carbonyls (Fig. 1).

Fig. 3
figure 3

Yellow index of artificially aged and naturally weathered PB-1 samples with the application of accelerated time equivalent

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Changes in supermolecular structure

The development of crystallinity during the degradation of both sets is recorded in Fig. 4. It must be emphasized that both series of samples were prepared from the exact extrusion string. WAXS measurements were performed in reflection mode; although an increase in crystallinity in the sample's center might not be observable, higher crystallinity was expected in the surface layers, which was confirmed by the performed measurements. In both cases, the expected increase in crystallinity occurred due to the process of chemi-crystallization, as mentioned above. After 20 h of artificial aging (37 days of natural weathering), differences between the development of crystallinity were visible. When the samples were artificially aged, the crystallinity slightly raised. Increasing the time of natural weathering (more than 37 days) led to slight decrease in the crystallinity due to chain shortening.

Fig. 4
figure 4

Development of crystallinity during the artificial aging and natural weathering of PB-1 samples with the application of accelerated time equivalent

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Surface observation of the artificially aged samples is depicted in Fig. 5, from which it is evident that the visible surface cracks started to appear after 40 h of UV irradiation. The crack orientation was parallel to the extrusion direction. These parallel cracks were then perpendicularly connected, thus creating a cracked net after 60 h of irradiation; further UV irradiation led to wider cracks.

Fig. 5
figure 5

Stereomicroscopic and SEM micrographs of PB-1 samples subjected to a progressive degree of UV irradiation for: a 20, b 40, and c 60 h compared with those of naturally weathered samples for: d 37, e 72, and f 102 days related on the basis of the accelerated time equivalent

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On the other hand, natural weathering led to a different crack initiation mechanism, leading to a higher quantity of more condensed initial cracks. The surface structure of the sample with most prolonged exposure to natural weathering (102 days) was consisted of high amounts of individual small cracks, which were only partially connected into wider and bigger cracks.

Mechanical properties

Figures 6 − 8 show the evolution of the mechanical properties of the samples during aging. Compared to artificial aging, natural weathering led to different results in all three measured parameters (tensile modulus, tensile strength and elongation-at-break). In the case of the artificially aged samples, the tensile modulus gradually decreased from approximately 500 to 450 MPa, except for the initial 20 h of the measurement (Fig. 6). Conversely, in the case of the natural weathering, there was a slight increase in the modulus; the modulus values were partly corresponded to the samples artificially irradiated for 20 h. This phenomenon could be attributed to chemi-crystallization, where chain scission and rearrangement into crystalline domains occurred, explaining the slight increase in tensile modulus in the case of the accelerated UV aging. In natural weathering due to chemi-crystallization, the crystallinity increased in the first exposure phase (37 days). Then, it slowly decreased while remaining higher than it was at the beginning of exposure (Fig. 4).

Fig. 6
figure 6

Development of the tensile modulus of the PB-1 samples during artificial aging and natural weathering with an accelerated time equivalent

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Figure 7 shows the dependence of the tensile strength-at-break on the exposure time. As with the tensile modulus, there was again a slight increase in the tensile strength within the first 10 h of exposure for the artificially aged samples; again, this could be attributed to a slight increase in crystallinity. With an increasing exposure time, the decrease from 21 to 15 MPa occurred.

Fig. 7
figure 7

Development of tensile strength-at-break of PB-1 samples during artificial aging and natural weathering with an accelerated time equivalent

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The first cracks were visible after 40 h of exposure, but at 20 h of exposure, the seeds of cracks must have been already formed in the samples (or shortened chains, caused a decrease in strength), while the subsequent aging caused a decrease in strength to 14.5 MPa. On the other hand, the tensile strength of the naturally aged samples slightly decreased after 24 days of exposure. However, after 37 days of exposure, the tensile strength drops to half of the original value; on this day, the cracks were not directly observable, but the chains of the sample were so shortened, that the tensile strength was dramatically reduced from 20 MPa to 10.5 MPa. In the following days, there was a slight decrease to approximate values of 8 to 9 MPa.

The development of elongation-at-break in the first 60 h of irradiation can be observed in Fig. 8. The radical downfall of elongation-at-break value was observed during the first 20 h of the UV irradiation, from 38% to merely 3% in the case of the naturally weathered sample and approximately 7% in the case of the artificially aged samples. Further irradiation has an almost negligible effect. Elongation hovered around 4% for naturally weathered samples and around 10% for artificially aged samples; it is necessary to consider that those values were shallow and could therefore be assumed as a minimum for irradiated samples. A faster decrease in elongation for naturally weathered samples could be resulted from temperature fluctuations and rain washing of the sample surface and, thus, accelerated the degradation process [21]. The fact that chemi-crystallization and chain shortening generally led to decreased ductility was consistent with findings in this work [44].

Fig. 8
figure 8

Development of elongation-at-break of PB-1 samples during artificial aging and natural weathering with an accelerated time equivalent

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After comparing both aging methods, it was evident that individual aging conditions impacted the resulting properties of the degraded material. The most significant differences were on the surface of the examined samples, and those differences were projected into the mechanical properties in different ways. The naturally weathered samples were both more brittle than the artificially aged samples while having a higher tensile modulus [45]. Although there were almost no differences in the polymorphic structure of PB-1 and the total crystallinity was almost constant during the experiment, the surface changes induced by both degradation methods seemed to be the driving mechanism of the changes in the mechanical properties.

For the sake of the chemical product development during aging, the carbonyl index and yellow index provided an acceptable correlation and, thus, can be used to roughly estimate the material state of degradation in the initial period.

Comparison of the morphologies of the degraded surfaces via stereomicroscopy and SEM contributed to understanding the different trends in the mechanical behavior of the degraded samples. In natural weathering, the large number of minor cracks caused a significant decrease in strength and ductility. Concurrently, the lower energy of UV irradiation, cyclical day and night shifting, and rain/wind washing all contributed to the development of chemi-crystallization, thus maintaining the tensile modulus still relatively high. In contrast, the number of cracks created during artificial aging was lower; however, the cracks reached much larger dimensions, contributing to drop in the tensile modulus. The reasoning behind the lower number of initial cracks could be the surface that was oversaturated with chemical products of the degradation (aldehydes and ketones), which were not washed away, leading to the formation of a protective layer, and degradation was more pronounced in the deep and wide cracks [46].

Conclusion

This work is the first to compare the degradation behavior of isotactic poly(1-butene) (PB-1) under artificial aging and natural weathering conditions, using accelerated time equivalent to correlate the final properties of the aged PB-1. It also investigated the effects of degradation on the mechanical properties, surface morphology, and crystalline structure of PB-1. The tailoring of the accelerated aging experiment led to slightly different mechanical behavior in the same material, mainly attributed to differences in crack development. During natural weathering, the chemi-crystallization phenomenon was observed, which led to an increase in the tensile modulus; on the other hand, during artificial aging, the material did not have enough time to supply different contaminants, and although its crystallinity slightly increased, the modulus decreased due to the size of the surface cracks. For both methods, the elongation-at-break and tensile strength values decreased, which was the expected outcome of any photo-oxidative process. Artificial aging predicted the behavior of naturally weathered material quite well and within an acceptable range of values. The only exception was the tensile strength-at-break value which was higher for the accelerated samples. However, the decrease was outside the acceptable range and general usability was still predicted correctly.