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Morphology of high-density polyethylene pipes stored under hydrostatic pressure at elevated temperature
Polymer xxx (2014) 1e9
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Morphology of high-density polyethylene pipes stored under hydrostatic pressure at elevated temperature Ning Sun a, Mirko Wenzel b, Alina Adams a, * a b
Institut für Technische und Makromolekulare Chemie, RWTH Aachen University, Templergraben 55, 52056 Aachen, Germany SKZ-German Plastic Center, Friedrich-Bergius-Ring 22, 97076 Würzburg, Germany
a r t i c l e i n f o
a b s t r a c t
Article history: Received 28 February 2014 Received in revised form 19 May 2014 Accepted 20 May 2014 Available online xxx
The morphology of unimodal and bimodal high-density polyethylene (HDPE) pipes during a hydrostatic pressure test was studied in detail using 1H solid-state NMR. Characterizing the changes of the molecular network during such a test is of key importance for understanding the long-term properties of different HDPE pipe grades. The changes in amount, thickness, and molecular mobility of the crystalline phase, the interface, and the amorphous phase of the two pipe grades with the storage time have been quantified for the first time. The most sensitive microscopic parameter to storage is the molecular mobility of the amorphous phase, with the strongest changes shown by the unimodal HDPE. The density of the tiemolecules is not the main factor controlling the very different behavior of the two pipe grades, but rather it is the density of the entanglements. The NMR results offer unprecedented insights into the changes in the molecular network and support existing deformation models. © 2014 Elsevier Ltd. All rights reserved.
Keywords: High-density polyethylene pipes Hydrostatic pressure test Solid-state NMR
1. Introduction The use of high-density polyethylene (HDPE) in applications requiring a long service lifetime such as for pressurized pipes for water and gas, has dramatically increased in the last years. For the new generation of pipe materials, a lifetime as long as 100 years is expected under normal operation conditions. To ensure proper performance over such a long time, precisely predicting the behavior of the HDPE pipes under the respective storage conditions is very important. Such information is usually acquired using accelerated laboratory aging procedures, one of which is the hydrostatic pressure test, the most frequently and widely used method. During this test, one can differentiate between three different failure stages according to the relation between the hoop stress (s) and the failure time (ts) [1,2]. In stage I, the fractures are ductile. Here, a yield followed by a large deformation of the material can be observed. In stage II, the fractures are brittle being characterized by the formation of slit leakage caused by slow crack growth (SCG) [3e5]. This kind of failure has great practical importance, because it occurs at low stress and can affect the long-term properties of the pipe material. An excellent improvement in
* Corresponding author. Tel.: þ49 241 80 26428; fax: þ49 241 80 22185. E-mail addresses: [email protected], [email protected] (A. Adams).
resistance to SCG was achieved with the development of new grades of pipe material such as bimodal PEs composed of linear short chains and long chains with short side branches. Yet, despite their superior performance, a molecular understanding of the much better resistance is still missing [6e8]. Stage III shows brittle fractures as well but with simultaneous appearance of numerous slit leakages [1,2,9,10]. During this stage, the failure happens at much lower stresses and it requires a certain degree of oxidative deterioration [10,11]. This chemical deterioration is caused by the migration and the gradual loss of antioxidants from the bulk material and, consequently, by the oxidation of the molecular chains [11]. As the failure mechanisms attributed to each stage are different, it is important to obtain a deeper knowledge of the responsible changes in the molecular network for a better understanding of the structureeproperty relationships [8]. These insights, in turn, can be used to design cost-saving alternative testing methods which could replace the extremely time-consuming hydrostatic pressure tests for predicting the long-term performance of a polymer material. Polyethylene (PE) is a semi-crystalline material with a morphology consisting of stacks of crystal lamellae with noncrystalline regions sandwiched in between them. During the last decades, different analytical methods such as differential scanning calorimetry (DSC) [12], infrared spectroscopy (IR) [13], and solidstate Nuclear Magnetic Resonance (NMR) [14e16] have been applied to morphologically characterize PE. Yet, most of these cited
http://dx.doi.org/10.1016/j.polymer.2014.05.056 0032-3861/© 2014 Elsevier Ltd. All rights reserved.
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methods use a two-phase crystalline-non-crystalline model to describe the complex morphology of PE. Nowadays, however, it is generally accepted that a more precise description of the morphology of PE is based on a three-phase model that accounts for the existence of an interface between the crystalline and the amorphous phase. It is evident that a clearer understanding of the very different behaviors of HDPE pipe grades during the hydrostatic pressure test requires analytical methods that can directly detect all phases and monitor their changes with the storage time. In particular, 1H solid-state NMR is one of the most powerful analytical tools, next to 13C solid-state NMR, capable of differentiating the three phases of PE based on their different molecular mobilities which, in turn, also enables acquiring information about the fraction of protons in each phase (the so-called “phase composition”) [14,17]. This method, which can directly probe the interface and the amorphous phase, has already been successfully applied to study, for example, the morphological changes of PE under annealing and in the presence of fillers [17,18]. Besides, it was shown that 1H spinespin relaxation measurements are sensitive to entanglements, as demonstrated for ultra-high molecular weight PE samples in the melt state [19,20]. The presence of entanglements restricts the molecular mobility of the polymer chains as they act like cross-links. Thus, the spinespin relaxation time T2 of entangled polymer chains is smaller compared with the value describing the molecular mobility of disentangled polymer chains [19,20]. Furthermore, low-field proton NMR was employed to investigate the thermo-oxidative aging of HDPE [21] and cross-linked PE [22]. Both studies showed that the phase composition and the molecular mobility of each phase change with the aging time and that the thermo-oxidative aging of PE is induced by chain scissions in the amorphous phase. Litvinov and Soliman used 1H solid-state NMR together with dynamic mechanical thermal analysis (DMTA), DSC, and X-Ray Diffraction (XRD) to investigate the changes in pipes made of polypropylene copolymers mainly during stage I of a hydrostatic pressure test at different testing temperatures [23]. It was found that the molecular mobility measured by 1H solid-state NMR relaxation is the most sensitive microscopic parameter with regard to the storage conditions of the polypropylene copolymers and that 1 H solid-state NMR can even probe morphological changes at longer storage times where the stabilizers were already consumed. They showed that the failure is related to the embrittlement of the material due to the perfection of existing crystals and to the formation of new crystals as well as to the chain elongation in the noncrystalline regions. The aim of this current work is to present a detailed investigation of the morphological changes in PE pipes, made from bimodal and unimodal resins, under hydrostatic pressure and at a temperature of 80 C, by applying various 1H solid-state NMR techniques. The bimodal pipe remains in stage I of the hydrostatic pressure test while the unimodal one reaches stage II after a short storage time. To the best of our knowledge, never before have morphological changes in PE pipes been directly observed experimentally and reported for all three phases in terms of amounts, chain mobility, and domain thickness with the storage time and the results discussed in relation to the proposed mechanisms for ductile and brittle failure.
the HDPE pipe group. This resin has a weight-average molecular weight Mw of about 250 kg/mol. According to ISO 12162, this material should have at least a lifetime of 50 years at 20 C and under a minimum hoop stress of 10 MPa [24]. The other pipe class, named here PEHD, is a unimodal PE. It was chosen to gain a better understanding of the molecular mechanisms behind the aging process, since this pipe material reaches the stage II of the hydrostatic pressure test after a short storage time. This resin has a Mw of about 160 kg/mol. For comparison reasons, PE100 and PEHD pipes without exposure to the hydrostatic pressure test were investigated as well. These samples will be named as ‘non-exposed’. 2.2. Hydrostatic pressure test The hydrostatic pressure tests of the HDPE pipes were carried out using a standard setup. The ends of pipes were closed by end caps, and controlled hydrostatic pressure was applied to the inner part of the pipe. The whole pipe segment was immersed in a water bath at a constant temperature of 80 C. In order to prevent chemical reactions, the water in and outside of the pipe was deionized and the hydrostatic pressure inside was provided by nitrogen gas. Different hydrostatic pressures resulting in hoop stresses s within the range of 2e7 MPa were applied to observe different failures with respect to the storage time ts. After various storage times ts, the pipe segments were removed from the hydrostatic test and then investigated after carefully cutting small pieces of pipe material. Each of these pipe segments belongs to a new pipe sample which was aged and measured independently of the others. None of the investigated pipe pieces in this work had macroscopically visible deformations. 2.3.
1
H NMR measurements
The 1H NMR measurements were performed using a 200 MHz Bruker DSX spectrometer working at a proton frequency of 200.12 MHz and a 4 mm solid-state probe-body without background signal. All measurements were conducted under static conditions using a radio-frequency pulse of 3 ms and a dwell time of 0.5 ms. Preliminary tests showed that using a measurement temperature of 90 C helps in better distinguishing among the three phases of PE. The phase composition and the transverse relaxation times T2 of the PE samples were extracted from simple and fast 1H Free Induction Decay (FID) measurements (Fig. 1). The best fit results of the recorded FIDs could be obtained by combining an Abragam function [25] to model the decay of the crystalline phase, a Weilbull function for the interface, and an exponential function for the amorphous phase [16,17] as shown in Eq. (1):
"
" # 1 t 2 sinð2pntÞ i þ Ið0Þ exp IðtÞ ¼ Ið0Þ exp 2 T2c 2pnt t þ Ið0Þa exp a T2 c
t T2i
!n #
(1) 2. Experimental 2.1. Materials The pipes investigated here were produced by SKZ with a throughput of 40 kg/h and a barrel temperature of 190 C. Two different HDPE resins were used for this study. The first one is PE100 which is a bimodal PE and it is currently the highest class in
where t denotes the time, I the signal intensity, and T2 the transverse relaxation time. The superscripts c, i, and a stand for the crystalline phase, the interface, and the amorphous phase. The term n denotes the sinusoidal oscillation constant and n is the exponent of the Weilbull function. In order to reduce the number of fitting parameters from Eq. (1), the values of the relaxation times of the crystalline and amorphous phase and the values of n and n were
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and at different times during the NMR experiments showed no changes with the time. 3. Results 3.1. Hoop stress
Fig. 1. Typical 1H FID of PE and the fitting result. The contributions of three phases are presented with different line styles. For clarity, only some of the experimental points are presented.
obtained for each sample from filtering experiments. These values were then kept constant during fitting the corresponding FIDs. For all investigated samples n was 1.5 while 2pn was 0.14 for the PE100 samples and 0.13 for the PHDE samples. The phase composition can be deduced from the signal intensity of each phase at zero acquisition time. For instance, the degree of crystallinity can be calculated by:
wc ¼ Ið0Þc
.h i Ið0Þc þ Ið0Þi þ Ið0Þa :
(2)
The thickness of each phase was measured with the help of 1H spin-diffusion NMR using a double-quantum (DQ) filter (Fig. 2). A detailed description of such measurements can be found, for example, in Refs. [26] and [27]. The values of t and tDQ were fixed to 5 ms to ensure that the signal from the crystalline phase can be properly selected. The values of the spin-diffusion time tm were varied from 20 ms to 300 ms. In order to minimize the effects of T1 in the spin-diffusion measurements values of tm not exceeding 300 ms were used. The three-component fitting model described by the Eq. (1) was used to obtain the amount of each phase at different spin-diffusion times. In order to check the reproducibility of the results, the NMR measurements were performed three times for each sample. All reported results are related to the exposure to the hydrostatic pressure test and not to changes induced by the storage of the samples at room temperature after their removal from the hydrostatic pressure test or by annealing during the NMR experiments. Proton FID measurements performed on various samples at different times after the removal from the hydrostatic pressure test
Fig. 2. Proton spin-diffusion NMR pulse-sequence with a double-quantum filter to measure domain thickness.
The results of the hydrostatic pressure tests are illustrated in Fig. 3. For the PE100 with storage times ts of up to 12 000 h, the absolute gradient of jd (lg s)/d (lg ts)j was around 0.03. The applied hydrostatic pressures were higher than 5.5 MPa and the detected fractures were ductile. All these characteristics clearly indicate a typical stage I failure. However, the PEHD pipes behave differently. A stage II failure characterized by jd (lg s)/d (lg ts)j of around 0.3 and hydrostatic pressures of less than 5.5 MPa can be detected for this pipe material. The transition from stage I to stage II takes place within a region as depicted in Fig. 3. The initial s during stage I failure of both pipe classes looks similar which indicates that they have about the same tensile strengths. No stage III failure was observed for either HDPE grades. This implies that the chemical deterioration is largely absent as confirmed by Fourier Transform Infrared Spectroscopy (FT-IR) measurements (data not shown). 3.2. Phase composition and molecular mobility The ability of 1H FID measurements to determine the phase composition of PE samples relies on the large difference in the molecular mobility of the three phases at a measurement temperature that is significantly higher than the corresponding glass transition temperature Tg. In the crystalline regions the molecular mobility is restricted and this is characterized by a short spinespin relaxation time T2. Even though the polymer chains in the interface and amorphous phase are more mobile than those in the crystalline phase, they still experience the constraints imposed by the crystalline phase and by other topological constraints such as entanglements. The chains in the amorphous phase, which are further away from the surface of the crystals than the chains in the interface, experience the constraint of the crystalline phase to a lesser extent, thereby leading to longer T2 values. As the molecular mobility in the interface regions is more restricted than that in the amorphous phase, it is thus characterized by an intermediate T2
Fig. 3. Variation of the hoop stress with the storage time. The lines serve as a visual guide. The marked area represents the transition region from stage I to stage II for the PEHD pipe grade.
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value. By analyzing the FID decays with appropriate analytical fitting models such as described by Eq. (1), information about the T2 values and the amount of each phase can be obtained. Yet, noteworthy here is that the crystallinity values acquired by NMR may not be identical to those found by other methods due to various factors: 1. The model used to describe the morphology of PE (twophase model versus three-phase model); 2. Depending on the NMR measurement temperature, a part of the interface may behave either more like a crystalline phase or more like an amorphous one. Nevertheless, the observed trends are generally the same for various methods. Fig. 4 depicts the changes in the phase composition as a function of storage time by using 1H FID measurements. The presented data show that the degree of crystallinity of non-exposed PE100 is higher than that of non-exposed PEHD. The amounts of the interface and of the amorphous phase of non-exposed PEHD are respectively higher and slightly higher than those of non-exposed PE100. This sequential arrangement in the amount of the three phases of the two samples is conserved with the storage time. The phase composition of both samples shows the same trend with the most significant change for the crystalline phase. The degree of crystallinity of the exposed samples is generally higher than that of the non-exposed samples and increases with the storage time reaching a kind of saturation at longer storage times. For PE100, the amounts of the interface and amorphous phase decrease steadily with ts. For PEHD, the amount of the interface decreases during stage I and then increases back to its initial value at ts ¼ 0 during stage II. The amount of the amorphous phase of PEHD decreases slightly with ts. As shown in Fig. 5, the relaxation times of all phases for both PE100 and PEHD decrease with the storage time ts but with overall higher values for the interface and amorphous phase for PE100 than for PEHD. The decline in T2 is more significant in PEHD than in PE100, especially for the interface and the amorphous phase. For instance, T2 of the amorphous phase of PEHD falls about 30% after a storage time of 3300 h compared with the value of the nonexposed sample. The reduction is only about 15% for PE100 even after 12000 h storage time. The values of the relaxation times at this long storage time are almost at the same level as those for the PEHD pipe, when the transition from stage I-failure to stage II-failure occurs. 3.3. Thickness of domains by 1H spin-diffusion NMR Spin-diffusion NMR is an important analytical tool used to acquire information about the thickness of domains in various
multiphase polymers including semi-crystalline polymers at a temperature well above their Tg [17,18,26,27]. In a first step, this experiment is based on the selection of the nuclear magnetization of one of the phases of the semi-crystalline polymer with the help of various kinds of dipolar filters [26,27]. In a second step, the equilibration of the magnetization within all phases is then allowed and monitored by means of a dipolar-mediated flip-flop process and, later on, is analyzed by appropriate fitting models [26,27]. Fig. 6 depicts typical 1H FIDs of the HDPE sample recorded with the pulse-sequence depicted in Fig. 2 and at relevant spin-diffusion times. The analysis of the proton FID for each spin diffusion time in terms of three components leads to the spin diffusion build-up and decay curves (Fig. 7). After the application of the DQ-filter using the parameters reported in the experimental section and at short spin-diffusion time tm, the magnetization is concentrated only in the crystalline phase. This is reflected in the high-intensity signal of this phase and no signal from the other two phases (Figs. 6a and 7). Upon increasing the spin-diffusion time, the magnetization flows from the crystalline phase into the interface and then further to the amorphous phase as it can clearly be observed from the increase in the intensity of the signal from the interface at shorter values for tm and later on also from that of the amorphous phase (Fig. 6bed and Fig. 7). The intensity ratio of the phases reaches constant values for tm higher than 100 ms (Fig. 7). This indicates that, at these spindiffusion times, the magnetization is already homogeneously distributed in all three phases. The obtained spin-diffusion curves were analyzed based on analytical solutions developed to describe the spin-diffusion process in lamellar morphologies composed of three phases with different diffusivities for the whole range of the spin-diffusion times [27]. Such a one-dimensional model can be applied for PE for which a lamellar morphology is largely accepted [17,18]. Moreover, we assume that it can also be used for the samples investigated in this study in both stages of the hydrostatic pressure test. Particularly, during stage II, the morphology of the samples is mainly controlled by the morphology of the bulk material, because the amount of material in the crazed regions (if they at all exist) is far less than that of the non-crazed material. The input parameters needed to obtain the theoretical spin-diffusion curves are the proton densities (r) and the spin-diffusion coefficients (D). For HDPE, the density of the crystalline phase is 0.99 g/cm3 and that of the amorphous phase is 0.87 g/cm3 [28]. Based on these values, proton densities of rc ¼ 0.14 g/cm3 and ra ¼ 0.12 g/cm3 were calculated. The values of the spin-diffusion coefficients of the crystalline phase Dc and the amorphous phase Da were calculated
Fig. 4. Change in the phase composition as a function of storage time ts for: (a) PE100 and (b) PEHD. The marked area represents the transition region from stage I to stage II. The lines serve as a visual guide.
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Fig. 5. Decrease in the proton transverse relaxation times T2 as a function of storage time ts for: (a) PE100 and (b) PEHD. The marked area represents the transition region from stage I to stage II. The lines serve as a visual guide.
according to the equations proposed in Ref. [29]. The spin-diffusion coefficient of the interface Di was calculated as an arithmetic average between Dc and Da according to the strategy proposed in Refs. [17,27]. The values for the spin-diffusion coefficients obtained as a function of ts are depicted in Fig. 8. As expected, the spin-diffusion coefficients of the crystalline phase are much higher than those of the amorphous phase due to the different chain mobilities and they show similar behavior with the storage time as for the T2 relaxation
times. Furthermore, their values are in the same range as those reported in Ref. [17]. There it had been clearly shown that such values for the spin-diffusion coefficients lead to values of domain thickness for HDPE which concur with the values obtained by other established methods. Typical fitting results of spin-diffusion data using the procedure outlined above are presented in Fig. 7. A good quality of the fitted experimental data can be clearly observed. Fig. 9 illustrates the changes in the domain thickness as a function of ts. The difference between the two pipe grades is also
Fig. 6. Normalized FIDs of HDPE recorded with the pulse-sequence from Fig. 2 at different spin-diffusion times tm: (a) 20 ms, (b) 1 ms, (c) 22 ms, and (d) 300 ms and the corresponding fitting results. To better observe the quality of the fitting, only some of the experimental points are presented.
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Besides the calculation of the long period, the interlamellar spacing (the distance between two adjacent lamella) Li ¼ 2di þ da and its changes with the storage time can be determined. The interlamellar spacing is slightly higher for the non-exposed PE100 sample than for the non-exposed PEHD sample and it shows only a small decrease with the storage time. Even after 12,000 h of hydrostatic pressure testing, the interlamellar spacing of PE100 has a higher value than that of non-exposed PEHD. In the case of PEHD, the interlamellar spacing strongly decreases with the storage time. 4. Discussions 4.1. Fracture behavior of the PE pipes exposed to hydrostatic pressure
Fig. 7. Typical 1H spin-diffusion NMR curves for PE. The symbols are the experimental data and the solid lines are the fitting results using analytical solutions for a onedimensional spin-diffusion process.
obvious here: the domain thickness of the crystalline phase (dc) of the non-exposed PE100 is about 30% higher than that of the nonexposed PEHD in overall and the thickness of the amorphous phase (da) is about 5% higher. The thicknesses of the interface (di) have a similar value. The obtained values for the thickness of the crystalline phase of the non-exposed samples are in the same range as those reported in Ref. [17] for a HDPE sample. Under the effect of hydrostatic pressure and temperature, the thickness of the crystalline phase increases with the storage time for both samples. The thickness of the interface decreases and that of the amorphous phase increases with the storage time for PE100 and for PEHD during stage I. These two phases show a reversed trend for the PEHD samples in stage II: Whereas the interface becomes thicker, the amorphous phase becomes thinner. The changes in the long period Lp ¼ dc þ 2di þ da as a function of the storage time are depicted in Fig. 10. For both non-exposed pipe materials, the acquired values are in the same range as those reported for an HDPE material [17]. The value for Lp slightly rises during stage I of the hydrostatic pressure test for both pipe materials and it seems to reach a plateau in stage II of the aging test for the PEHD pipe. Our data show that the changes in the long period are mainly controlled by the changes in the thickness of the crystalline phase. Thus, the smaller changes in the thickness of the interface and mobile amorphous phase, which may finally have a greater impact on the failure as one may expect, will not be reflected in the variation of Lp. Therefore, monitoring the changes of the long period during the hydrostatic pressure test will give only limited insight into the morphological changes that are occurring.
In the time frame of our investigation, the bimodal pipe material PE100 exhibits only ductile failure. By contrast, PEHD, a unimodal pipe material, already reaches the brittle regime after a short storage time. This indicates that this resin is not suited for pipe applications requiring long-term stability. The observed failure behaviors can be explained based on different mechanisms. Ductile failure is mainly attributed to material damage caused by mechanical load. At longer storage times and lower pressures, the material breaks due to brittle failure. The damage usually initiates at a defect or a stress concentration position inside the pipe wall due to the presence of carbon black, pigments, additive, or catalytic residues. Yet, according to Ref. [30], catalytic residues are an initiating factor of creep damage, but their presence is not sufficient to generate an SCG in PE, the nature of the PE matrix being a predominant factor. Thus it has still not been elucidated which properties of the polymer matrix actually are responsible for initiating such damage. 4.2. Phase composition, thickness of the domains, and molecular mobility The combined exposure of PE pipes to hydrostatic pressure and temperature leads to morphological changes with the storage time and the largest changes are shown by unimodal PEHD. Both samples display the same trend for all NMR parameters during stage I. This suggests that the observed microscopic properties are independent of the sample type but are strongly dependent on the storage conditions. The degree of crystallinity, the thickness of the crystalline phase and of the amorphous phase, and the long period all increase during this stage. Meanwhile, the amount of interface and mobile amorphous phase decreases, while the interlamellar spacing and the molecular mobility in all three phases diminish. In particular, the thickness of the amorphous phase and the long
Fig. 8. 1H spin-diffusion coefficients versus the storage time for: (a) PE100 and (b) PEHD. The marked area represents the transition region from stage I to stage II.
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Fig. 9. Variation in the domain sizes with the storage time ts for: (a) PE100 and (b) PEHD. The marked area represents the transition region from stage I to stage II. The lines serve as a visual guide.
period reach maximum values during the transition region from stage I to stage II. This transition is accompanied by a slight change in the crystallinitydin agreement with the results from the Ref. [31]dand by a decrease in the interlamellar spacing. In particular, a critical interlamellar spacing of about 6e7 nm was proposed by some authors as a possible embrittlement criterion to describe the ductile-to-brittle transition for PE samples aged under other conditions than in the current study [32,33]. In the case of PEHD sample studied here, the ductile-to-brittle transition occurs when the interlamellar spacing is about 6.4 nm. This result indicates that the interlamellar spacing could indeed be defined as a directly measurable embrittlement criterion. During stage II, the amount and thickness of the crystalline phase and the long period seems to be largely independent of the storage time. On the contrary, the amount and the thickness of the interface increase at the expense of the amount and the thickness of the amorphous phase. The interlamellar spacing decreases further on with the storage time, and the molecular mobility in all phases becomes more restricted. The observed morphological changes in both stages are usually attributed to the combined effect of annealing and hydrostatic pressure [23]. Yet based on our NMR data, qualitative information about the competition between these two processes could be obtained. Elevated temperatures provide an ideal condition to initialize annealing. The increased mobility with rising temperature, also partially due to the melting of small crystals, together with chain
diffusion into and out of the crystals [34] and chain rearrangement upon prolonged annealing lead to morphological changes towards a more thermodynamically stable crystalline structure with a higher amount of crystalline phase and a thicker lamella. Annealing of an HDPE sample at 121 C, which is only 15 C below the reported melting temperature, for about 100 h results in a 1%-increase in the amount of the crystalline phase at the expense of the interface and the amorphous phase [17]. This is accompanied by a perfection of the crystalline phase, an increase in its thickness and a decrease of the thickness of the interface and a decrease of about 1e2% of the T2 relaxation times. Similar trends were observed for our data but with much stronger changes within the same time frame as in Ref. [17] even for a storage temperature of 80 C which is more than 50 C below the melting temperatures of the two pipe materials. These findings clearly imply that annealing has a reduced effect on the detected morphological changesdmostly on the perfection of the crystalline structure. More, compared to Ref. [17], a different behavior for the thickness of the amorphous phase was observed, during stage I and for the interface during stage II which indicates that these changes have to be related to the effect of the hydrostatic pressure. Hydrostatic pressure subjects the pipes to constant stress over time. NMR studies of polypropylene showed that the polymer chains orient themselves and lose mobility under the application of a uniaxial stress [35,36]. A substantial reduction in the chain mobility, being most pronounced in the amorphous phase, was clearly observed also for our samples during the whole range of
Fig. 10. The long period Lp (a) and the interlamellar spacing Li (b) as a function of the storage time ts, respectively. The marked area represents the transition region from stage I to stage II for the PEHD pipe grade. The lines serve as a visual guide.
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storage times. Like annealing, drawing contributes to an increase in the fraction of the rigid phase at the expense of both semi-rigid and mobile amorphous fractions as well as to an increase in the thickness of the rigid phase and in the long period [35]. With an increased amount and thickness of the crystalline phase and a continuously diminishing interlamellar spacing, additional constraints are imposed on the chain dynamics of both interface and amorphous phase. Therefore, continuously shorter T2 relaxation times are found with increased storage time. To understand the obtained results for the thicknesses of the amorphous phase and the interface during stage I, one can make a small analogy with the effect of stretching a cross-linked rubber. Under stretching, the rubber chains start to orient themselves and the distance between two cross-link points becomes greater upon stretching compared to that of the non-stretched material. This analogy also applies to the chains in the amorphous phase where the crystal lamellae could be regarded as physical cross-link points [37]. Upon stretching, the chains in the amorphous phase start to orient themselves, thereby resulting in an increase in the thickness of this domain and a decrease in chain mobilitydwhich agrees with proposed deformation models [38]. However, the amount of this phase decreases, because a part of it is transferred to the interface. The same increased thickness would be expected for the interface. Here, probably due to the restriction of the chain mobility upon stretching and additional constrains imposed by the crystals, a part of the interface close to the surface of the crystal contribute to it and the increased thickness is no longer seen. Consequently, a reduced amount and thickness of the interface is observed. At a certain storage time, the thickness of the amorphous phase peaks as the chains cannot stretch further. At this point, the lamellae may break up into smaller units, according to the “mosaic block” model proposed by Peterlin [38,39]. Reaching the maximum thickness of the amorphous phase seems to be the point where the PEHD sample reaches stage II. From now on, the amount and the thickness of the amorphous phase decrease. Since the amount and the thickness of the crystalline phase are largely independent of the storage time, the only observed changes take place at the interphase between the amorphous phase and the interface. The chains in the amorphous phase become increasingly restricted and, thus, a larger part of the amorphous phase is transformed into the interface region. At a certain point, almost no chains will exist anymore that would count for the amorphous phase. Moreover, the interlamellar spacing decreases with the storage time and probably reaches a minimum limit under which the chains can no longer sustain the applied stress; this ought to be the point where the material starts to deform [31]. Since a distribution of domain thicknesses exists in PE, in some regions this decrease in the interlamellar spacing will be faster than in others. Probably here are the “low stability” points at which crazes are initiated. The presence of impurities such as catalyst residues generates additional stress locally that will compound the effect of the hydrostatic pressure. This will lead to an additional decrease in the interlamellar spacing. Therefore, if this is also a region where the interlamellar spacing is smaller, precisely this will be the spot where the crazes will be initiated.
Upon considering our results which clearly show that the amount of crystallinity and the crystal thickness are higher for PE100 than for PEHD, the use of the Brown model leads to the same density of the tie-molecules of 0.333 for both non-exposed resins (see SI for the details of calculation). Moreover, this value is independent of the storage time for both resins under the assumption that the molecular weight remains constant. This result clearly shows that the theoretical density of tie-molecules fails to explain the very different behavior of the bimodal resin compared with that of the unimodal one under storage conditions. The density of tiemolecules does play a role in the long-term properties of HDPE pipes but it is not the main controlling factor. Ultimately, this is not surprising as most of the load in the amorphous phase is borne not only by the tie-molecules but also by entanglements [6,42]. Even though the role of the entangled amorphous network during the deformation of semi-crystalline polymers has been recognized [6,8,43], there is very little data regarding a correlation between the entanglement density and the fracture behavior of PE pipes and about a higher density of trapped entanglements compared with the density of tie-molecules [44,45]. The smaller changes in the T2 values of the chains in the interface and especially in the amorphous phase can be related to a smaller degree of deformation of the chains in these phases for PE100 than for PEHD with respect to the storage time [20]. This is also reflected by the slower increase in the thickness of the amorphous phase of PE100 compared to PEHD during stage I although both pipe grades are exposed to the same stress. Thus, this result proves that the density of entanglements between the crystallites is higher in PE100 compared to that of PEHD. This conclusion is supported by Raman results on drawn samples of various PE pipes where it was shown that the mechanically active chain portions are less strained in bimodal copolymers than in equivalent unimodal copolymers [46]. This conclusion is further supported by the higher T2 value of the non-exposed PE100 that of the non-exposed PEHD. A higher amount of topological constraints, such as the amount and the thickness of the crystalline phase and the amount of entanglements, in PE100 than in PEHD would normally lead to a stronger reduction in T2 of the two phases in PE100 compared to PEHD. The interlamellar spacing for both non-exposed samples has values in the same range and therefore, a similar constraint effect on the chains. Yet, the unexpected observation of higher molecular mobility in PE100 than in PEHD could be understood if one takes into account that the chains orient themselves during the extrusion of the pipes. Due to the lower density of entanglements, the chains in PEHD become oriented more easily than those in PE100 which, in turn, leads to lower relaxation times for the unimodal pipe grade. This explanation needs to be further verified by investigating, for example, PEHD grades produced at different extrusion speeds, which were unfortunately not available for the current study. Should this interpretation be confirmed, extreme care will have to be taken during the extrusion of polymer pipes; the induced orientation will lead to even faster morphological changes and, thus, to faster embrittlement. 5. Conclusions
4.3. Density of tie-molecules and entanglements The different behavior of bimodal and unimodal pipe materials observed during hydrostatic pressure tests is often attributed to the difference in the density of tie-molecules as one of the main factors controlling the long-term properties of PE pipes [40,41]. Computations based on the theoretical model proposed by Brown predict a drastic rise in the tie-molecule probability as crystal thickness diminishes with decreasing crystallinity [40,41].
Based on simple 1H NMR measurements and for the first time, this study has provided detailed information about the morphological changes in HDPE pipes, i.e. in a unimodal PEHD pipe and a bimodal PE100 pipe, during a hydrostatic pressure test at a temperature of 80 C. The dependence of the amount, molecular mobility, and thickness of the crystalline phase, interface, and amorphous phase on the storage time was quantified using a single experimental setup, rather than a combination of various analytical
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methods. The unimodal pipe resin PEHD displays the most pronounced morphological changes, reaching stage II of the hydrostatic pressure test after a short storage time. By contrast, the bimodal pipe grade PE100 shows only small morphological changes, and it stays in stage I during the investigated time frame. Our results clearly demonstrate that the detected morphological changes are caused by combined annealing and creep with the highest effects induced by creep. Such information can be acquired, only because NMR garners direct information about changes in all phases of HDPE, and especially of the amorphous phase and of the molecular mobility. The largest changes with respect to the storage time are observed for the amount of the crystalline phase, the interlamellar spacing, and for the molecular mobility of the amorphous phase. In particular, the value of the interlamellar spacing at the ductileebrittle transition is in the same range as those values reported for polyethylene samples aged under other conditions; thus, this microscopic parameter may further help to define a directly measurable embrittlement criterion. The changes at the microscopic level during stage I, as detected by NMR, not only support the existing deformation models but they give a deeper insight into the changes in the molecular network during this stage. Furthermore, the arguments of various researchers for the initiation of the crazes during stage II are supported based on the changes in the molecular network. It was also found that the very different behavior of the two pipe grades under the hydrostatic pressure test cannot be explained by the density of tie-molecules but by the density of entanglements. The changes in the values of the spinespin relaxation time, especially of the amorphous phase, which is an indirect measure of this density with the highest value for the bimodal pipe gradedcould also be exploited to define an additional embrittlement criterion. The great advantage of this approach is that this NMR relaxation parameter is very sensitive and is accessible fast and easily. Moreover, even for a polymer pipe currently in use and made of resin other than PE, this NMR relaxation parameter can be obtained in a non-destructive way by employing inexpensive single-sided NMR devices [47]. This means that the status of a polymer pipe in use could be monitored by simply measuring the changes in the values of T2 of the amorphous phase in comparison with those of the unexposed pipe. Further work along this line is in progress. Acknowledgments The authors acknowledge financial support by AiF, project IGFFV 16019 N, Prof. B. Blümich, ITMC, RWTH Aachen University for his support during this project, and the reviewers for their useful comments. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.polymer.2014.05.056.
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Morphology of high-density polyethylene pipes stored under hydrostatic pressure at elevated temperature Ning Sun a, Mirko Wenzel b, Alina Adams a, * a b
Institut für Technische und Makromolekulare Chemie, RWTH Aachen University, Templergraben 55, 52056 Aachen, Germany SKZ-German Plastic Center, Friedrich-Bergius-Ring 22, 97076 Würzburg, Germany
a r t i c l e i n f o
a b s t r a c t
Article history: Received 28 February 2014 Received in revised form 19 May 2014 Accepted 20 May 2014 Available online xxx
The morphology of unimodal and bimodal high-density polyethylene (HDPE) pipes during a hydrostatic pressure test was studied in detail using 1H solid-state NMR. Characterizing the changes of the molecular network during such a test is of key importance for understanding the long-term properties of different HDPE pipe grades. The changes in amount, thickness, and molecular mobility of the crystalline phase, the interface, and the amorphous phase of the two pipe grades with the storage time have been quantified for the first time. The most sensitive microscopic parameter to storage is the molecular mobility of the amorphous phase, with the strongest changes shown by the unimodal HDPE. The density of the tiemolecules is not the main factor controlling the very different behavior of the two pipe grades, but rather it is the density of the entanglements. The NMR results offer unprecedented insights into the changes in the molecular network and support existing deformation models. © 2014 Elsevier Ltd. All rights reserved.
Keywords: High-density polyethylene pipes Hydrostatic pressure test Solid-state NMR
1. Introduction The use of high-density polyethylene (HDPE) in applications requiring a long service lifetime such as for pressurized pipes for water and gas, has dramatically increased in the last years. For the new generation of pipe materials, a lifetime as long as 100 years is expected under normal operation conditions. To ensure proper performance over such a long time, precisely predicting the behavior of the HDPE pipes under the respective storage conditions is very important. Such information is usually acquired using accelerated laboratory aging procedures, one of which is the hydrostatic pressure test, the most frequently and widely used method. During this test, one can differentiate between three different failure stages according to the relation between the hoop stress (s) and the failure time (ts) [1,2]. In stage I, the fractures are ductile. Here, a yield followed by a large deformation of the material can be observed. In stage II, the fractures are brittle being characterized by the formation of slit leakage caused by slow crack growth (SCG) [3e5]. This kind of failure has great practical importance, because it occurs at low stress and can affect the long-term properties of the pipe material. An excellent improvement in
* Corresponding author. Tel.: þ49 241 80 26428; fax: þ49 241 80 22185. E-mail addresses: [email protected], [email protected] (A. Adams).
resistance to SCG was achieved with the development of new grades of pipe material such as bimodal PEs composed of linear short chains and long chains with short side branches. Yet, despite their superior performance, a molecular understanding of the much better resistance is still missing [6e8]. Stage III shows brittle fractures as well but with simultaneous appearance of numerous slit leakages [1,2,9,10]. During this stage, the failure happens at much lower stresses and it requires a certain degree of oxidative deterioration [10,11]. This chemical deterioration is caused by the migration and the gradual loss of antioxidants from the bulk material and, consequently, by the oxidation of the molecular chains [11]. As the failure mechanisms attributed to each stage are different, it is important to obtain a deeper knowledge of the responsible changes in the molecular network for a better understanding of the structureeproperty relationships [8]. These insights, in turn, can be used to design cost-saving alternative testing methods which could replace the extremely time-consuming hydrostatic pressure tests for predicting the long-term performance of a polymer material. Polyethylene (PE) is a semi-crystalline material with a morphology consisting of stacks of crystal lamellae with noncrystalline regions sandwiched in between them. During the last decades, different analytical methods such as differential scanning calorimetry (DSC) [12], infrared spectroscopy (IR) [13], and solidstate Nuclear Magnetic Resonance (NMR) [14e16] have been applied to morphologically characterize PE. Yet, most of these cited
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methods use a two-phase crystalline-non-crystalline model to describe the complex morphology of PE. Nowadays, however, it is generally accepted that a more precise description of the morphology of PE is based on a three-phase model that accounts for the existence of an interface between the crystalline and the amorphous phase. It is evident that a clearer understanding of the very different behaviors of HDPE pipe grades during the hydrostatic pressure test requires analytical methods that can directly detect all phases and monitor their changes with the storage time. In particular, 1H solid-state NMR is one of the most powerful analytical tools, next to 13C solid-state NMR, capable of differentiating the three phases of PE based on their different molecular mobilities which, in turn, also enables acquiring information about the fraction of protons in each phase (the so-called “phase composition”) [14,17]. This method, which can directly probe the interface and the amorphous phase, has already been successfully applied to study, for example, the morphological changes of PE under annealing and in the presence of fillers [17,18]. Besides, it was shown that 1H spinespin relaxation measurements are sensitive to entanglements, as demonstrated for ultra-high molecular weight PE samples in the melt state [19,20]. The presence of entanglements restricts the molecular mobility of the polymer chains as they act like cross-links. Thus, the spinespin relaxation time T2 of entangled polymer chains is smaller compared with the value describing the molecular mobility of disentangled polymer chains [19,20]. Furthermore, low-field proton NMR was employed to investigate the thermo-oxidative aging of HDPE [21] and cross-linked PE [22]. Both studies showed that the phase composition and the molecular mobility of each phase change with the aging time and that the thermo-oxidative aging of PE is induced by chain scissions in the amorphous phase. Litvinov and Soliman used 1H solid-state NMR together with dynamic mechanical thermal analysis (DMTA), DSC, and X-Ray Diffraction (XRD) to investigate the changes in pipes made of polypropylene copolymers mainly during stage I of a hydrostatic pressure test at different testing temperatures [23]. It was found that the molecular mobility measured by 1H solid-state NMR relaxation is the most sensitive microscopic parameter with regard to the storage conditions of the polypropylene copolymers and that 1 H solid-state NMR can even probe morphological changes at longer storage times where the stabilizers were already consumed. They showed that the failure is related to the embrittlement of the material due to the perfection of existing crystals and to the formation of new crystals as well as to the chain elongation in the noncrystalline regions. The aim of this current work is to present a detailed investigation of the morphological changes in PE pipes, made from bimodal and unimodal resins, under hydrostatic pressure and at a temperature of 80 C, by applying various 1H solid-state NMR techniques. The bimodal pipe remains in stage I of the hydrostatic pressure test while the unimodal one reaches stage II after a short storage time. To the best of our knowledge, never before have morphological changes in PE pipes been directly observed experimentally and reported for all three phases in terms of amounts, chain mobility, and domain thickness with the storage time and the results discussed in relation to the proposed mechanisms for ductile and brittle failure.
the HDPE pipe group. This resin has a weight-average molecular weight Mw of about 250 kg/mol. According to ISO 12162, this material should have at least a lifetime of 50 years at 20 C and under a minimum hoop stress of 10 MPa [24]. The other pipe class, named here PEHD, is a unimodal PE. It was chosen to gain a better understanding of the molecular mechanisms behind the aging process, since this pipe material reaches the stage II of the hydrostatic pressure test after a short storage time. This resin has a Mw of about 160 kg/mol. For comparison reasons, PE100 and PEHD pipes without exposure to the hydrostatic pressure test were investigated as well. These samples will be named as ‘non-exposed’. 2.2. Hydrostatic pressure test The hydrostatic pressure tests of the HDPE pipes were carried out using a standard setup. The ends of pipes were closed by end caps, and controlled hydrostatic pressure was applied to the inner part of the pipe. The whole pipe segment was immersed in a water bath at a constant temperature of 80 C. In order to prevent chemical reactions, the water in and outside of the pipe was deionized and the hydrostatic pressure inside was provided by nitrogen gas. Different hydrostatic pressures resulting in hoop stresses s within the range of 2e7 MPa were applied to observe different failures with respect to the storage time ts. After various storage times ts, the pipe segments were removed from the hydrostatic test and then investigated after carefully cutting small pieces of pipe material. Each of these pipe segments belongs to a new pipe sample which was aged and measured independently of the others. None of the investigated pipe pieces in this work had macroscopically visible deformations. 2.3.
1
H NMR measurements
The 1H NMR measurements were performed using a 200 MHz Bruker DSX spectrometer working at a proton frequency of 200.12 MHz and a 4 mm solid-state probe-body without background signal. All measurements were conducted under static conditions using a radio-frequency pulse of 3 ms and a dwell time of 0.5 ms. Preliminary tests showed that using a measurement temperature of 90 C helps in better distinguishing among the three phases of PE. The phase composition and the transverse relaxation times T2 of the PE samples were extracted from simple and fast 1H Free Induction Decay (FID) measurements (Fig. 1). The best fit results of the recorded FIDs could be obtained by combining an Abragam function [25] to model the decay of the crystalline phase, a Weilbull function for the interface, and an exponential function for the amorphous phase [16,17] as shown in Eq. (1):
"
" # 1 t 2 sinð2pntÞ i þ Ið0Þ exp IðtÞ ¼ Ið0Þ exp 2 T2c 2pnt t þ Ið0Þa exp a T2 c
t T2i
!n #
(1) 2. Experimental 2.1. Materials The pipes investigated here were produced by SKZ with a throughput of 40 kg/h and a barrel temperature of 190 C. Two different HDPE resins were used for this study. The first one is PE100 which is a bimodal PE and it is currently the highest class in
where t denotes the time, I the signal intensity, and T2 the transverse relaxation time. The superscripts c, i, and a stand for the crystalline phase, the interface, and the amorphous phase. The term n denotes the sinusoidal oscillation constant and n is the exponent of the Weilbull function. In order to reduce the number of fitting parameters from Eq. (1), the values of the relaxation times of the crystalline and amorphous phase and the values of n and n were
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and at different times during the NMR experiments showed no changes with the time. 3. Results 3.1. Hoop stress
Fig. 1. Typical 1H FID of PE and the fitting result. The contributions of three phases are presented with different line styles. For clarity, only some of the experimental points are presented.
obtained for each sample from filtering experiments. These values were then kept constant during fitting the corresponding FIDs. For all investigated samples n was 1.5 while 2pn was 0.14 for the PE100 samples and 0.13 for the PHDE samples. The phase composition can be deduced from the signal intensity of each phase at zero acquisition time. For instance, the degree of crystallinity can be calculated by:
wc ¼ Ið0Þc
.h i Ið0Þc þ Ið0Þi þ Ið0Þa :
(2)
The thickness of each phase was measured with the help of 1H spin-diffusion NMR using a double-quantum (DQ) filter (Fig. 2). A detailed description of such measurements can be found, for example, in Refs. [26] and [27]. The values of t and tDQ were fixed to 5 ms to ensure that the signal from the crystalline phase can be properly selected. The values of the spin-diffusion time tm were varied from 20 ms to 300 ms. In order to minimize the effects of T1 in the spin-diffusion measurements values of tm not exceeding 300 ms were used. The three-component fitting model described by the Eq. (1) was used to obtain the amount of each phase at different spin-diffusion times. In order to check the reproducibility of the results, the NMR measurements were performed three times for each sample. All reported results are related to the exposure to the hydrostatic pressure test and not to changes induced by the storage of the samples at room temperature after their removal from the hydrostatic pressure test or by annealing during the NMR experiments. Proton FID measurements performed on various samples at different times after the removal from the hydrostatic pressure test
Fig. 2. Proton spin-diffusion NMR pulse-sequence with a double-quantum filter to measure domain thickness.
The results of the hydrostatic pressure tests are illustrated in Fig. 3. For the PE100 with storage times ts of up to 12 000 h, the absolute gradient of jd (lg s)/d (lg ts)j was around 0.03. The applied hydrostatic pressures were higher than 5.5 MPa and the detected fractures were ductile. All these characteristics clearly indicate a typical stage I failure. However, the PEHD pipes behave differently. A stage II failure characterized by jd (lg s)/d (lg ts)j of around 0.3 and hydrostatic pressures of less than 5.5 MPa can be detected for this pipe material. The transition from stage I to stage II takes place within a region as depicted in Fig. 3. The initial s during stage I failure of both pipe classes looks similar which indicates that they have about the same tensile strengths. No stage III failure was observed for either HDPE grades. This implies that the chemical deterioration is largely absent as confirmed by Fourier Transform Infrared Spectroscopy (FT-IR) measurements (data not shown). 3.2. Phase composition and molecular mobility The ability of 1H FID measurements to determine the phase composition of PE samples relies on the large difference in the molecular mobility of the three phases at a measurement temperature that is significantly higher than the corresponding glass transition temperature Tg. In the crystalline regions the molecular mobility is restricted and this is characterized by a short spinespin relaxation time T2. Even though the polymer chains in the interface and amorphous phase are more mobile than those in the crystalline phase, they still experience the constraints imposed by the crystalline phase and by other topological constraints such as entanglements. The chains in the amorphous phase, which are further away from the surface of the crystals than the chains in the interface, experience the constraint of the crystalline phase to a lesser extent, thereby leading to longer T2 values. As the molecular mobility in the interface regions is more restricted than that in the amorphous phase, it is thus characterized by an intermediate T2
Fig. 3. Variation of the hoop stress with the storage time. The lines serve as a visual guide. The marked area represents the transition region from stage I to stage II for the PEHD pipe grade.
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value. By analyzing the FID decays with appropriate analytical fitting models such as described by Eq. (1), information about the T2 values and the amount of each phase can be obtained. Yet, noteworthy here is that the crystallinity values acquired by NMR may not be identical to those found by other methods due to various factors: 1. The model used to describe the morphology of PE (twophase model versus three-phase model); 2. Depending on the NMR measurement temperature, a part of the interface may behave either more like a crystalline phase or more like an amorphous one. Nevertheless, the observed trends are generally the same for various methods. Fig. 4 depicts the changes in the phase composition as a function of storage time by using 1H FID measurements. The presented data show that the degree of crystallinity of non-exposed PE100 is higher than that of non-exposed PEHD. The amounts of the interface and of the amorphous phase of non-exposed PEHD are respectively higher and slightly higher than those of non-exposed PE100. This sequential arrangement in the amount of the three phases of the two samples is conserved with the storage time. The phase composition of both samples shows the same trend with the most significant change for the crystalline phase. The degree of crystallinity of the exposed samples is generally higher than that of the non-exposed samples and increases with the storage time reaching a kind of saturation at longer storage times. For PE100, the amounts of the interface and amorphous phase decrease steadily with ts. For PEHD, the amount of the interface decreases during stage I and then increases back to its initial value at ts ¼ 0 during stage II. The amount of the amorphous phase of PEHD decreases slightly with ts. As shown in Fig. 5, the relaxation times of all phases for both PE100 and PEHD decrease with the storage time ts but with overall higher values for the interface and amorphous phase for PE100 than for PEHD. The decline in T2 is more significant in PEHD than in PE100, especially for the interface and the amorphous phase. For instance, T2 of the amorphous phase of PEHD falls about 30% after a storage time of 3300 h compared with the value of the nonexposed sample. The reduction is only about 15% for PE100 even after 12000 h storage time. The values of the relaxation times at this long storage time are almost at the same level as those for the PEHD pipe, when the transition from stage I-failure to stage II-failure occurs. 3.3. Thickness of domains by 1H spin-diffusion NMR Spin-diffusion NMR is an important analytical tool used to acquire information about the thickness of domains in various
multiphase polymers including semi-crystalline polymers at a temperature well above their Tg [17,18,26,27]. In a first step, this experiment is based on the selection of the nuclear magnetization of one of the phases of the semi-crystalline polymer with the help of various kinds of dipolar filters [26,27]. In a second step, the equilibration of the magnetization within all phases is then allowed and monitored by means of a dipolar-mediated flip-flop process and, later on, is analyzed by appropriate fitting models [26,27]. Fig. 6 depicts typical 1H FIDs of the HDPE sample recorded with the pulse-sequence depicted in Fig. 2 and at relevant spin-diffusion times. The analysis of the proton FID for each spin diffusion time in terms of three components leads to the spin diffusion build-up and decay curves (Fig. 7). After the application of the DQ-filter using the parameters reported in the experimental section and at short spin-diffusion time tm, the magnetization is concentrated only in the crystalline phase. This is reflected in the high-intensity signal of this phase and no signal from the other two phases (Figs. 6a and 7). Upon increasing the spin-diffusion time, the magnetization flows from the crystalline phase into the interface and then further to the amorphous phase as it can clearly be observed from the increase in the intensity of the signal from the interface at shorter values for tm and later on also from that of the amorphous phase (Fig. 6bed and Fig. 7). The intensity ratio of the phases reaches constant values for tm higher than 100 ms (Fig. 7). This indicates that, at these spindiffusion times, the magnetization is already homogeneously distributed in all three phases. The obtained spin-diffusion curves were analyzed based on analytical solutions developed to describe the spin-diffusion process in lamellar morphologies composed of three phases with different diffusivities for the whole range of the spin-diffusion times [27]. Such a one-dimensional model can be applied for PE for which a lamellar morphology is largely accepted [17,18]. Moreover, we assume that it can also be used for the samples investigated in this study in both stages of the hydrostatic pressure test. Particularly, during stage II, the morphology of the samples is mainly controlled by the morphology of the bulk material, because the amount of material in the crazed regions (if they at all exist) is far less than that of the non-crazed material. The input parameters needed to obtain the theoretical spin-diffusion curves are the proton densities (r) and the spin-diffusion coefficients (D). For HDPE, the density of the crystalline phase is 0.99 g/cm3 and that of the amorphous phase is 0.87 g/cm3 [28]. Based on these values, proton densities of rc ¼ 0.14 g/cm3 and ra ¼ 0.12 g/cm3 were calculated. The values of the spin-diffusion coefficients of the crystalline phase Dc and the amorphous phase Da were calculated
Fig. 4. Change in the phase composition as a function of storage time ts for: (a) PE100 and (b) PEHD. The marked area represents the transition region from stage I to stage II. The lines serve as a visual guide.
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Fig. 5. Decrease in the proton transverse relaxation times T2 as a function of storage time ts for: (a) PE100 and (b) PEHD. The marked area represents the transition region from stage I to stage II. The lines serve as a visual guide.
according to the equations proposed in Ref. [29]. The spin-diffusion coefficient of the interface Di was calculated as an arithmetic average between Dc and Da according to the strategy proposed in Refs. [17,27]. The values for the spin-diffusion coefficients obtained as a function of ts are depicted in Fig. 8. As expected, the spin-diffusion coefficients of the crystalline phase are much higher than those of the amorphous phase due to the different chain mobilities and they show similar behavior with the storage time as for the T2 relaxation
times. Furthermore, their values are in the same range as those reported in Ref. [17]. There it had been clearly shown that such values for the spin-diffusion coefficients lead to values of domain thickness for HDPE which concur with the values obtained by other established methods. Typical fitting results of spin-diffusion data using the procedure outlined above are presented in Fig. 7. A good quality of the fitted experimental data can be clearly observed. Fig. 9 illustrates the changes in the domain thickness as a function of ts. The difference between the two pipe grades is also
Fig. 6. Normalized FIDs of HDPE recorded with the pulse-sequence from Fig. 2 at different spin-diffusion times tm: (a) 20 ms, (b) 1 ms, (c) 22 ms, and (d) 300 ms and the corresponding fitting results. To better observe the quality of the fitting, only some of the experimental points are presented.
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Besides the calculation of the long period, the interlamellar spacing (the distance between two adjacent lamella) Li ¼ 2di þ da and its changes with the storage time can be determined. The interlamellar spacing is slightly higher for the non-exposed PE100 sample than for the non-exposed PEHD sample and it shows only a small decrease with the storage time. Even after 12,000 h of hydrostatic pressure testing, the interlamellar spacing of PE100 has a higher value than that of non-exposed PEHD. In the case of PEHD, the interlamellar spacing strongly decreases with the storage time. 4. Discussions 4.1. Fracture behavior of the PE pipes exposed to hydrostatic pressure
Fig. 7. Typical 1H spin-diffusion NMR curves for PE. The symbols are the experimental data and the solid lines are the fitting results using analytical solutions for a onedimensional spin-diffusion process.
obvious here: the domain thickness of the crystalline phase (dc) of the non-exposed PE100 is about 30% higher than that of the nonexposed PEHD in overall and the thickness of the amorphous phase (da) is about 5% higher. The thicknesses of the interface (di) have a similar value. The obtained values for the thickness of the crystalline phase of the non-exposed samples are in the same range as those reported in Ref. [17] for a HDPE sample. Under the effect of hydrostatic pressure and temperature, the thickness of the crystalline phase increases with the storage time for both samples. The thickness of the interface decreases and that of the amorphous phase increases with the storage time for PE100 and for PEHD during stage I. These two phases show a reversed trend for the PEHD samples in stage II: Whereas the interface becomes thicker, the amorphous phase becomes thinner. The changes in the long period Lp ¼ dc þ 2di þ da as a function of the storage time are depicted in Fig. 10. For both non-exposed pipe materials, the acquired values are in the same range as those reported for an HDPE material [17]. The value for Lp slightly rises during stage I of the hydrostatic pressure test for both pipe materials and it seems to reach a plateau in stage II of the aging test for the PEHD pipe. Our data show that the changes in the long period are mainly controlled by the changes in the thickness of the crystalline phase. Thus, the smaller changes in the thickness of the interface and mobile amorphous phase, which may finally have a greater impact on the failure as one may expect, will not be reflected in the variation of Lp. Therefore, monitoring the changes of the long period during the hydrostatic pressure test will give only limited insight into the morphological changes that are occurring.
In the time frame of our investigation, the bimodal pipe material PE100 exhibits only ductile failure. By contrast, PEHD, a unimodal pipe material, already reaches the brittle regime after a short storage time. This indicates that this resin is not suited for pipe applications requiring long-term stability. The observed failure behaviors can be explained based on different mechanisms. Ductile failure is mainly attributed to material damage caused by mechanical load. At longer storage times and lower pressures, the material breaks due to brittle failure. The damage usually initiates at a defect or a stress concentration position inside the pipe wall due to the presence of carbon black, pigments, additive, or catalytic residues. Yet, according to Ref. [30], catalytic residues are an initiating factor of creep damage, but their presence is not sufficient to generate an SCG in PE, the nature of the PE matrix being a predominant factor. Thus it has still not been elucidated which properties of the polymer matrix actually are responsible for initiating such damage. 4.2. Phase composition, thickness of the domains, and molecular mobility The combined exposure of PE pipes to hydrostatic pressure and temperature leads to morphological changes with the storage time and the largest changes are shown by unimodal PEHD. Both samples display the same trend for all NMR parameters during stage I. This suggests that the observed microscopic properties are independent of the sample type but are strongly dependent on the storage conditions. The degree of crystallinity, the thickness of the crystalline phase and of the amorphous phase, and the long period all increase during this stage. Meanwhile, the amount of interface and mobile amorphous phase decreases, while the interlamellar spacing and the molecular mobility in all three phases diminish. In particular, the thickness of the amorphous phase and the long
Fig. 8. 1H spin-diffusion coefficients versus the storage time for: (a) PE100 and (b) PEHD. The marked area represents the transition region from stage I to stage II.
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Fig. 9. Variation in the domain sizes with the storage time ts for: (a) PE100 and (b) PEHD. The marked area represents the transition region from stage I to stage II. The lines serve as a visual guide.
period reach maximum values during the transition region from stage I to stage II. This transition is accompanied by a slight change in the crystallinitydin agreement with the results from the Ref. [31]dand by a decrease in the interlamellar spacing. In particular, a critical interlamellar spacing of about 6e7 nm was proposed by some authors as a possible embrittlement criterion to describe the ductile-to-brittle transition for PE samples aged under other conditions than in the current study [32,33]. In the case of PEHD sample studied here, the ductile-to-brittle transition occurs when the interlamellar spacing is about 6.4 nm. This result indicates that the interlamellar spacing could indeed be defined as a directly measurable embrittlement criterion. During stage II, the amount and thickness of the crystalline phase and the long period seems to be largely independent of the storage time. On the contrary, the amount and the thickness of the interface increase at the expense of the amount and the thickness of the amorphous phase. The interlamellar spacing decreases further on with the storage time, and the molecular mobility in all phases becomes more restricted. The observed morphological changes in both stages are usually attributed to the combined effect of annealing and hydrostatic pressure [23]. Yet based on our NMR data, qualitative information about the competition between these two processes could be obtained. Elevated temperatures provide an ideal condition to initialize annealing. The increased mobility with rising temperature, also partially due to the melting of small crystals, together with chain
diffusion into and out of the crystals [34] and chain rearrangement upon prolonged annealing lead to morphological changes towards a more thermodynamically stable crystalline structure with a higher amount of crystalline phase and a thicker lamella. Annealing of an HDPE sample at 121 C, which is only 15 C below the reported melting temperature, for about 100 h results in a 1%-increase in the amount of the crystalline phase at the expense of the interface and the amorphous phase [17]. This is accompanied by a perfection of the crystalline phase, an increase in its thickness and a decrease of the thickness of the interface and a decrease of about 1e2% of the T2 relaxation times. Similar trends were observed for our data but with much stronger changes within the same time frame as in Ref. [17] even for a storage temperature of 80 C which is more than 50 C below the melting temperatures of the two pipe materials. These findings clearly imply that annealing has a reduced effect on the detected morphological changesdmostly on the perfection of the crystalline structure. More, compared to Ref. [17], a different behavior for the thickness of the amorphous phase was observed, during stage I and for the interface during stage II which indicates that these changes have to be related to the effect of the hydrostatic pressure. Hydrostatic pressure subjects the pipes to constant stress over time. NMR studies of polypropylene showed that the polymer chains orient themselves and lose mobility under the application of a uniaxial stress [35,36]. A substantial reduction in the chain mobility, being most pronounced in the amorphous phase, was clearly observed also for our samples during the whole range of
Fig. 10. The long period Lp (a) and the interlamellar spacing Li (b) as a function of the storage time ts, respectively. The marked area represents the transition region from stage I to stage II for the PEHD pipe grade. The lines serve as a visual guide.
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storage times. Like annealing, drawing contributes to an increase in the fraction of the rigid phase at the expense of both semi-rigid and mobile amorphous fractions as well as to an increase in the thickness of the rigid phase and in the long period [35]. With an increased amount and thickness of the crystalline phase and a continuously diminishing interlamellar spacing, additional constraints are imposed on the chain dynamics of both interface and amorphous phase. Therefore, continuously shorter T2 relaxation times are found with increased storage time. To understand the obtained results for the thicknesses of the amorphous phase and the interface during stage I, one can make a small analogy with the effect of stretching a cross-linked rubber. Under stretching, the rubber chains start to orient themselves and the distance between two cross-link points becomes greater upon stretching compared to that of the non-stretched material. This analogy also applies to the chains in the amorphous phase where the crystal lamellae could be regarded as physical cross-link points [37]. Upon stretching, the chains in the amorphous phase start to orient themselves, thereby resulting in an increase in the thickness of this domain and a decrease in chain mobilitydwhich agrees with proposed deformation models [38]. However, the amount of this phase decreases, because a part of it is transferred to the interface. The same increased thickness would be expected for the interface. Here, probably due to the restriction of the chain mobility upon stretching and additional constrains imposed by the crystals, a part of the interface close to the surface of the crystal contribute to it and the increased thickness is no longer seen. Consequently, a reduced amount and thickness of the interface is observed. At a certain storage time, the thickness of the amorphous phase peaks as the chains cannot stretch further. At this point, the lamellae may break up into smaller units, according to the “mosaic block” model proposed by Peterlin [38,39]. Reaching the maximum thickness of the amorphous phase seems to be the point where the PEHD sample reaches stage II. From now on, the amount and the thickness of the amorphous phase decrease. Since the amount and the thickness of the crystalline phase are largely independent of the storage time, the only observed changes take place at the interphase between the amorphous phase and the interface. The chains in the amorphous phase become increasingly restricted and, thus, a larger part of the amorphous phase is transformed into the interface region. At a certain point, almost no chains will exist anymore that would count for the amorphous phase. Moreover, the interlamellar spacing decreases with the storage time and probably reaches a minimum limit under which the chains can no longer sustain the applied stress; this ought to be the point where the material starts to deform [31]. Since a distribution of domain thicknesses exists in PE, in some regions this decrease in the interlamellar spacing will be faster than in others. Probably here are the “low stability” points at which crazes are initiated. The presence of impurities such as catalyst residues generates additional stress locally that will compound the effect of the hydrostatic pressure. This will lead to an additional decrease in the interlamellar spacing. Therefore, if this is also a region where the interlamellar spacing is smaller, precisely this will be the spot where the crazes will be initiated.
Upon considering our results which clearly show that the amount of crystallinity and the crystal thickness are higher for PE100 than for PEHD, the use of the Brown model leads to the same density of the tie-molecules of 0.333 for both non-exposed resins (see SI for the details of calculation). Moreover, this value is independent of the storage time for both resins under the assumption that the molecular weight remains constant. This result clearly shows that the theoretical density of tie-molecules fails to explain the very different behavior of the bimodal resin compared with that of the unimodal one under storage conditions. The density of tiemolecules does play a role in the long-term properties of HDPE pipes but it is not the main controlling factor. Ultimately, this is not surprising as most of the load in the amorphous phase is borne not only by the tie-molecules but also by entanglements [6,42]. Even though the role of the entangled amorphous network during the deformation of semi-crystalline polymers has been recognized [6,8,43], there is very little data regarding a correlation between the entanglement density and the fracture behavior of PE pipes and about a higher density of trapped entanglements compared with the density of tie-molecules [44,45]. The smaller changes in the T2 values of the chains in the interface and especially in the amorphous phase can be related to a smaller degree of deformation of the chains in these phases for PE100 than for PEHD with respect to the storage time [20]. This is also reflected by the slower increase in the thickness of the amorphous phase of PE100 compared to PEHD during stage I although both pipe grades are exposed to the same stress. Thus, this result proves that the density of entanglements between the crystallites is higher in PE100 compared to that of PEHD. This conclusion is supported by Raman results on drawn samples of various PE pipes where it was shown that the mechanically active chain portions are less strained in bimodal copolymers than in equivalent unimodal copolymers [46]. This conclusion is further supported by the higher T2 value of the non-exposed PE100 that of the non-exposed PEHD. A higher amount of topological constraints, such as the amount and the thickness of the crystalline phase and the amount of entanglements, in PE100 than in PEHD would normally lead to a stronger reduction in T2 of the two phases in PE100 compared to PEHD. The interlamellar spacing for both non-exposed samples has values in the same range and therefore, a similar constraint effect on the chains. Yet, the unexpected observation of higher molecular mobility in PE100 than in PEHD could be understood if one takes into account that the chains orient themselves during the extrusion of the pipes. Due to the lower density of entanglements, the chains in PEHD become oriented more easily than those in PE100 which, in turn, leads to lower relaxation times for the unimodal pipe grade. This explanation needs to be further verified by investigating, for example, PEHD grades produced at different extrusion speeds, which were unfortunately not available for the current study. Should this interpretation be confirmed, extreme care will have to be taken during the extrusion of polymer pipes; the induced orientation will lead to even faster morphological changes and, thus, to faster embrittlement. 5. Conclusions
4.3. Density of tie-molecules and entanglements The different behavior of bimodal and unimodal pipe materials observed during hydrostatic pressure tests is often attributed to the difference in the density of tie-molecules as one of the main factors controlling the long-term properties of PE pipes [40,41]. Computations based on the theoretical model proposed by Brown predict a drastic rise in the tie-molecule probability as crystal thickness diminishes with decreasing crystallinity [40,41].
Based on simple 1H NMR measurements and for the first time, this study has provided detailed information about the morphological changes in HDPE pipes, i.e. in a unimodal PEHD pipe and a bimodal PE100 pipe, during a hydrostatic pressure test at a temperature of 80 C. The dependence of the amount, molecular mobility, and thickness of the crystalline phase, interface, and amorphous phase on the storage time was quantified using a single experimental setup, rather than a combination of various analytical
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methods. The unimodal pipe resin PEHD displays the most pronounced morphological changes, reaching stage II of the hydrostatic pressure test after a short storage time. By contrast, the bimodal pipe grade PE100 shows only small morphological changes, and it stays in stage I during the investigated time frame. Our results clearly demonstrate that the detected morphological changes are caused by combined annealing and creep with the highest effects induced by creep. Such information can be acquired, only because NMR garners direct information about changes in all phases of HDPE, and especially of the amorphous phase and of the molecular mobility. The largest changes with respect to the storage time are observed for the amount of the crystalline phase, the interlamellar spacing, and for the molecular mobility of the amorphous phase. In particular, the value of the interlamellar spacing at the ductileebrittle transition is in the same range as those values reported for polyethylene samples aged under other conditions; thus, this microscopic parameter may further help to define a directly measurable embrittlement criterion. The changes at the microscopic level during stage I, as detected by NMR, not only support the existing deformation models but they give a deeper insight into the changes in the molecular network during this stage. Furthermore, the arguments of various researchers for the initiation of the crazes during stage II are supported based on the changes in the molecular network. It was also found that the very different behavior of the two pipe grades under the hydrostatic pressure test cannot be explained by the density of tie-molecules but by the density of entanglements. The changes in the values of the spinespin relaxation time, especially of the amorphous phase, which is an indirect measure of this density with the highest value for the bimodal pipe gradedcould also be exploited to define an additional embrittlement criterion. The great advantage of this approach is that this NMR relaxation parameter is very sensitive and is accessible fast and easily. Moreover, even for a polymer pipe currently in use and made of resin other than PE, this NMR relaxation parameter can be obtained in a non-destructive way by employing inexpensive single-sided NMR devices [47]. This means that the status of a polymer pipe in use could be monitored by simply measuring the changes in the values of T2 of the amorphous phase in comparison with those of the unexposed pipe. Further work along this line is in progress. Acknowledgments The authors acknowledge financial support by AiF, project IGFFV 16019 N, Prof. B. Blümich, ITMC, RWTH Aachen University for his support during this project, and the reviewers for their useful comments. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.polymer.2014.05.056.
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