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Correlation between isothermal crystallization properties and slow crack growth resistance of polyethylene pipe materials
Polymer Testing 80 (2019) 106128
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Material Behaviour
Correlation between isothermal crystallization properties and slow crack growth resistance of polyethylene pipe materials Farzad Gholami a, Gholamreza Pircheraghi a, *, Reza Rashedi b, Abdulhanan Sepahi b a b
Polymeric Materials Research Group (PMRG), Department of Materials Science and Engineering, Sharif University of Technology, Tehran, Iran Department of Polymer Engineering and Color Technology, Amirkabir University of Technology, Tehran, Iran
A R T I C L E I N F O
A B S T R A C T
Keywords: Isothermal crystallization Avrami index Morphology Slow crack growth SSA-DSC
In this study, three PE100 polyethylene materials with fairly different short chain branch distributions were used to characterize the relation between thermal properties and creep test failure time. The samples were thermally characterized using isothermal and non-isothermal differential scanning calorimetry (DSC). The three resins showed different behavior after fitting on the Avrami equation. N-100J2 sample, in which short chain branches (SCBs) are located on longer molecule chains, has a lower chain mobility and higher crystallization time, while N-100J1 has the opposite crystallization properties. Also, these samples had different Avrami index, n. Micro structural evaluation of these samples revealed that the Avrami index had a direct relationship with the morphology of samples and the higher n correlates to formation of more spherulite structures. The creep test result showed that the sample with a lower crystallization rate and lower Avrami index was more resistant against slow crack growth (SCG) because of its crystalline morphology with the obstacles in crack growth. There is a meaningful relationship between lamellae arrangement and failure time; when a crystalline segment in resin has more spherulite structures, higher Avrami index, the boundary of these spherulites may act as weak regions in front of crack and consequently lead to easier crack growth in such material.
1. Introduction Slow crack growth failure in polyethylene pipes has attracted considerable interest among researchers due to economic and safety reasons. Today, most failure in polyethylene pipes, which are widely in transporting water and gas, is due to slow crack growth (SCG) mecha nism [1–5]. It is quite common that the placement of comonomers on low or high molecular weight chains can significantly affect the me chanical properties of pipes. Selective localization of short chain branches (SCBs) on higher molecular weight fractions or longer chains increases the number of entanglements between molecules, which cau ses the resin to form thinner and fewer crystals with a huge number of tie molecules. As a result, these tie molecules make the resin tougher and increase the resistance of material against slow crack growth [6]. This concept was used in the development of bimodal polyethylene resins with specific short chain branch distribution, which is now widely known as PE100 pipe grade polyethylene [7]. To determine resistance to SCG, several methods and tests are stan dardized; most familiar means of characterizing creep behavior in polyethylene resins are Pennsylvania Edge Notch Test (PENT) [8], Full
Notch Creep Test (FNCT) [9], Notched Pipe Test (NPT) [10], and Crack Round Bar test (CRB) [11]. Yet, all these methods suffer from the same problem of being financially ineffective and time consuming, some of which taking even months to be completed. Few researchers [6,12,13] have addressed the use of other tech niques to determine the lifetime of polyethylene (PE) resins. Considering that crystalline structure in semi-crystalline polymers has a significant influence on mechanical properties, there is a direct relation between the morphology of lamellae and long-time performance of polyethylene resins [14–21]. Thickness and arrangement of lamellae are the two most important parameters that have the most effect on the mechanical properties of PE pipes. Thinner crystalline segments happen to have more tie molecules in between, which increases entanglement density in the amorphous phase and makes the material more resistant against craze propagation during slow crack growth [15,19,21,22]. On the other hand, the crystallinity percentage and arrangement of crystalline seg ments are critical factors that have a direct effect on mechanical prop erties of material, such as stiffness and yield strength [6,16]. The general factor that has a direct effect on thickness and arrangement of lamellae is crystallization kinetics influenced by polymer chain topology [20,
* Corresponding author. E-mail address: [email protected] (G. Pircheraghi). https://doi.org/10.1016/j.polymertesting.2019.106128 Received 3 June 2019; Received in revised form 31 August 2019; Accepted 26 September 2019 Available online 27 September 2019 0142-9418/© 2019 Published by Elsevier Ltd.
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2. Experimental section
Table 1 Coding system and properties of used polyethylenes. Samples Code
Polymer Type
MFI [190� C-5 kg] (g/10min)
Density (g/cm3)
N-100J1 N-100J2 N-100 M
PE100 PE100 PE100
0.22 0.18 0.21
0.956 0.950 0.954
Three different PE100 pipe grade polyethylene granules were sup plied by Jam and Marun petrochemical companies in Iran. All samples were produced through LyondelBasell’s Hostalen process using ZieglerNatta catalyst but with different polymerization conditions which can affect the comonomer distribution in final product. All samples use 1Butene comonomer and their codes, MFI and density are presented in Table 1. For sample coding, letter "N" refers to neat PE100 powder being used, letter "J" refers to samples produced by Jam Petrochemical com pany and "M" for sample produced by Marun Petrochemical company. To characterize the crystalline phase and to evaluate thermal prop erties of the samples, isothermal and non-isothermal crystallization of samples were studied by TA Instruments Q100 differential scanning calorimeter (DSC). Approximately 5 mg of the samples were weighed to carry out the test. In the non-isothermal test, samples were heated and cooled at 10 � C/min rate from 25 � C to 180 � C and this cycle was repeated three times to evaluate the crystallization and melting tem peratures as well as the degree of crystallization. For the isothermal test, samples were heated to 180 � C at 10 � C/min and kept for 3 min to ensure complete melting. Then, they were first cooled to 140 � C at 10 � C/min, held at 140 � C for 2 min, and were then rapidly cooled down (at 40 � C/ min) to and maintained at 121 � C for 2 h. The procedure which was previously used by Krishnaswamy et al. [6], made it possible to evaluate and compare crystallization time and rate of the samples. DSC test was conducted in a pure nitrogen atmosphere. After isothermal crystalliza tion tests, the obtained data were fitted into Avrami equation using a plugin for Origin® software developed by A. T. Lorenzo et al. [23]. Fractionation of molecular chains in the different PE100 compounds were performed using Successive self-nucleation and annealing (SSA) test. Unlike TREF and CRYSTAF tests, SSA is less time consuming and more cost-effective. In this test, different molecules in the polymer are separated from each other to form unique crystals with regards to their ability to crystallize. In this procedure samples were melted first at a temperature higher than the melting point of the resin, then cooled to room temperature in order to eliminate thermal history of samples. Afterward, for fractionation, the samples were heated to a temperature 8� C lower than previous one and maintained in that temperature for 15 min. This procedure was repeated for 4 steps and eventually, the sample was heated to 180� C and a multipeak DSC heating curve could be obtained [29]. In previous studies, SSA-DSC has been used for frac tionation of LLDPE which has a large number of short chain branches. Thus, the separation process of chains is done easily in a short time (5 min) and with small temperature steps(5� C) [29–33]. In this work, the PE100 samples have less than 1% comonomer in their structure and it is necessary to maintain the samples in annealing temperatures for longer times and larger temperature steps to ensure that the fraction ation of the chains occurs properly. After several trial and error tests, the optimum value for annealing time and temperature steps were found to be 15 min and 8 � C, respectively.
23–25]. Several researchers proposed [24–27] theoretical equations for investigation and prediction of crystallization kinetics, among which, the Avrami equation is the most common and practical theory for determination and prediction of polymer’s crystallization kinetics [23, 26]. Sarafpour et al. [20] used the Avrami equation and rheological measurements to predict the behavior of black polyethylene pipe with different masterbatches. They found that by decreasing molecular weight of carbon black carrier polymer, three-dimensional growth of crystal segments could change the crystalline structure and the pipe’s long term performance. Krishnaswamy et al. [6] have studied the in fluence of short chain distribution on mechanical and rheological properties of polyethylene resins. They used two different resins with different location of 1-Butene comonomers on long and short chains. In this investigation, isothermal crystallization is used to survey crystalli zation kinetics and then to study the effect of short-chain branches on mechanical and microstructural characteristics of the samples. They found that by locating SCBs on long chains, molecules were remarkably slower in crystallization, resulting in formation of more tie molecules and subsequently enhanced resistance to slow crack growth. Sardashti et al. [13] investigated the long-time performance of PE pipes with rheological test. They used relaxation time obtained from Carreau–Ya soda equation fitting on rheological measurement data and proposed that by increasing this factor, the chain entanglement would be increased, enhancing resistance to SCG as a result. Deveci et al. [28] used strain hardening test, crack round bar (CRB) test, and notched pipe test to study the relation between molecular weight and molecular weight distribution of PE resins and resistance to SCG phenomenon. They have found that to determine the effect of comonomer type, strain hardening test is better than CRB method. On the other hand, to detect the effect of other parameters like molecular weight, molecular weight distribution and comonomer content, both methods are highly efficient. The aim of this work has been to enhance the current knowledge about the relationship between polymer chain microstructure, forma tion, and morphology of crystalline domains and resistance to slow crack growth in polyethylene pipe materials. In this study, we have used PE100 polyethylene resins with different SCB distribution and almost identical average molecular weight. Slow crack growth has been assessed by means of full notched creep test (FNCT) while the crystalline characteristics has been analyzed using DSC measurements and FE-SEM imaging.
Fig. 1. (a) Schematic and (b) picture of the designed Creep test machine. 2
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Fig. 2. (a) Heating and (b) cooling curve of PE100 samples in non-isothermal DSC test. Table 2 Thermal properties of PE100 samples. sample
Tm ½� C�
Tc ½� C�
Xc ½%�
N-100J1 N-100J2 N-100 M
131.1 129.6 131.6
117.3 115.3 116.7
61.2 53.0 57.9
To evaluate the lifetime of the samples and their creep resistance, a lab scale full notch creep test (FNCT), designed according to ISO 16770 [9] standard, was performed. The instrument and a representative schematic are shown in Fig. 1. Samples with dimensions of 6 � 6 � 9 mm were cut from 6 mm compression molded sheets. The test was carried out at 55 � C with a dead load stress of 9 MPa. Average failure time of three specimens for each sample was measured. The crystalline morphology and lamellar structure of the samples were observed using field emission scanning electron microscope (FESEM). For this purpose, compresseion molded bars were fractured in liquid nitrogen. To remove the amorphous phase of samples, an acid solution (H2SO4:H3PO4: water ¼ 5:2:1 and 1% w/n KMnO4) was used as etchant, followed by washing in H2O2:H2SO4 (dilute) (H2SO4: H2O ¼ 2:7). The etching method was suggested by Eslamian et al. [16]. Fractography tests from creep samples were also carried out by FE-SEM.
Fig. 3. Heat-flow curve vs. time for PE100 samples obtained from isothermal crystallization. Table 3 Obtained parameters from curve fitting on Avrami equation.
3. Results and discussion 3.1. Crystallization kinetics and morphology DSC cooling and heating curves of three samples are shown in Fig. 2 and thermal properties are given in Table 2. It can be deducted from different results in Table 2 that the structure of these resins is not the same. Due to the utilization of identical Ziegler-Natta ethylene poly merization method in the presence of 1-Butene for production of sam ples, one can assume that the difference between samples is originated from different comonomer distribution and content. The lower crystal linity and melting point of N-100J2 can be attributed to its higher commoner content and evenly distributed 1-Butene in long chains [6,16, 33]. Also, the crystallization temperature, TC, of this sample is the lowest among the samples due to difficulty in the folding of its chain in the melt state. Isothermal crystallization curves of the three PE100 samples are shown in Fig. 3. From these curves, it is obvious that in N-100J2 sample, the necessary time for crystallization is higher than other samples and N100J1 crystallization time is the lowest. To quantify the results, these
Samples
t1/2(s)
K (Crystallization rate) (min-n)*
n (Avrami index)
N-100J1 N-100J2 N-100 M
48 158 60
0.340 0.109 0.249
2.88 2.45 2.80
experimental data were fitted to the Avrami equation (Eq. (1)). 1
Vc ðtÞ ¼ expð
Ktn Þ
(1)
in which Vc is relative crystalline fraction, K is crystallization rate con stant and n is Avrami index, which should have a value between 2 and 3 in crystallization of polymer, and t is the time. Obtained parameters are presented in Table 3. From Table 3, it can be inferred that N-100J2 chains are more sluggish than other resins and need more time for crystallization [6,34]. This causes a significant change in crystalline morphology and me chanical properties of materials which can have a direct impact on long-term performance since the rate of crystallization have a correla tion with chain entanglements in the structure of the materials [4,6,20, 21]. Another important parameter in this equation is n, which is related 3
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Fig. 4. FE-SEM images of crystalline structures of the samples at three magnifications (a) N-100J1, (b) N-100J2 and (c) N-100 M.
to the morphology of crystals. This parameter varies between 2 and 3 and the closer n gets to 3, the more number of spherulites can be formed, while with approaching 2, crystalline structure leans towards an unor dered lamellar morphology [20]. To verify the results of isothermal parameters as well as fitting of Avrami equation, etched samples were examined by FE-SEM to observe the morphology of samples. Fig. 4 shows lamella order of three PE100 samples in different magnifications. Spherulite morphology can be observed in N-100J1 and N-100 M (red circles), but in N-100J2 crys talline lamellas have a random placement. These images are in agree ment with the result of Avrami fitting that N-100J1 with higher Avrami index (2.88) has more spherulites in its structure while in N-100J2, which has lower n, there are no spherulites. This is in consistency with Sarafpour et al. [20] results. Observed morphologies are from hot pressed plates which are cooled non-isothermally, thus they may not exactly reveal the morphologies from isothermal crystallization, but they are useful for a qualitative comparison of different morphologies in crystallization of samples. It can be seen that N-100J2 has thinner lamellas and distance be tween these lamellas are higher than two other samples. Herein, one can assume that there were some tiny lamellas in the amorphous phase which could act as tie-crystals and have been removed by the etching
Table 4 Failure time of PE100 samples obtained from creep test. Samples
Failure time (h)
N-100J1
50 � 12
N-100J2
460 � 36
N-100 M
104 � 15
process as reported elsewhere [16]. Thinner lamellas between main lamellas of the N-100J2 sample can cause a significant increase in the number of tie molecules and probably make this resin tougher against SCG phenomenon. 3.2. Resistance against slow crack growth Results of the creep test by FNCT apparatus are presented in Table 4. It is clear that longterm properties of N-100J2 sample are significantly higher compared to other samples, showing excellent creep perfor mance. The main mechanism of the polymer failure in this test is slow crack growth, which is highly sensitive to the craze formation in the 4
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in all samples has the same brittle morphology and there is no noticeable difference between these three samples. In two other zones of the craze, growth is higher, and the surface is relatively rougher than the first zone. In the second zone in which ligament area decreases, some pull out re gions can be seen. These regions are fibrillations that are formed during craze growth and after failure, these signs remain. In the third zone, in which the load bearing area is very small and then stress is higher than other zones, the velocity of craze growth is significantly higher and the fracture surface is not brittle. By comparing the morphology of samples in these two zones, we can investigate the mechanism during creep test. In Fig. 6, it can be seen that zone II and III in N-100J2 are rougher compared to two other samples, indicating that in this sample, crazes could not grow fast because of the resistance of entanglement of polymer chains in amorphous region, and hence the formation of strong fibrils as
Fig. 5. (a) An overview of the fracture surface of creep test samples with three distinct regions with different morphologies. (b) Schematic models for explaining how different regions in the fracture surface are formed. With further progress of the crack in the samples, the stress level is increased and then larger fibrillations formed; as a result, the third zone is rougher than the first zone in the creep test fracture surface.
amorphous phase and fibrillation of polymer in the vicinity of formed crazes. Therefore, the crystalline morphology of the sample has a vital role in the craze growth and resistance against SCG. For further evaluation of creep behavior of these resins, fracture surfaces of the samples have been studied. After failure in the creep test, all samples show a unique arrangement with three different zones where each zone has its own characteristics. These three zones are shown in Fig. 5. As seen, Zone I is the first zone after notch and has a smooth morphology in all samples, since in this region ligament is large and the applied stress is low, so craze growth in this zone is very slow compared to other zones. Because crack growth rate in the zone I is slow, this zone
Fig. 7. The relation of creep failure time of PE100 samples against Avrami index n.
Fig. 6. Microscopic images of zone II and III of the fracture surface in creep test samples. 5
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Fig. 8. Schematic models of crystalline arrangement in (a) N-100J2 and (b) N-100J1 samples.
observed in this figure. These strong and rough fibrils suppress craze growth and cause more failure time in the FNCT test. On the other hand, in N-100J1 sample, which has the lowest creep lifetime, fibrillations are weaker and smoother, indicating an easier crack growth. Based on results from creep and isothermal crystallization tests, we can see a meaningful correlation between the Avrami parameters such as Avrami index n and crystallization rate K in Table 3 and creep lifetime in Table 4 of samples. Based on these results, the lower the value of the Avrami index and crystallization rate, the higher the failure time in the creep test. Fig. 7 shows the relation between the Avrami index, n, and the failure time in the creep test for three samples. It should be mentioned that although both Avrami index, n, and crystallization rate, k, have meaningful relation with failure time in FNCT test, and it is
possible to redraw the failure curve against n and k, but we believe Avrami index can be more important here as it indicates morphological properties of the samples, while K is merely a crystallization rate indi cator and it can be changed with other parameter like nucleating agents without significant changes in morphological properties. The reason for this correlation between Avrami index and creep failure time can be explained based on different crystalline morphology of the samples. As presented before, the Avrami index of N-100J1 sample is 2.88, which is close to 3, which makes this sample prone to formation of three-dimensional spherulites, as shown in Fig. 4. On the other hand, N-100J2 sample has an Avrami index of 2.45, close to 2, indicating tendency for random growth of two-dimensional lamellas. Considering the different crystalline morphologies of the samples, it is reasonable to assume that in the samples with 3D spherulite
Fig. 9. (a) Heat flow vs. temperature curves of PE100 samples after fractionation by SSA-DSC procedure and (b) peaks area portion of SSA curves. 6
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morphology, the spherulite boundaries play as weak zones where crack can grow easily. On the other hand, in the sample with 2D lamellar structure, there is no preferential path for crack growth. It is reported that when approaching spherulites, the boundaries are weakly entan gled and can be a favorable path for crack growth [35,36]. Fig. 8 shows a schematic representing two different crystalline structures that include two and three-dimensional growth of lamellas from the polymer melt. Fig. 8-a is a graphical representation of N-100J2 sample, while Fig. 8-b is a model of N-100J1 that has a preference for spherulite growth. As indicated in Fig. 8-b, there are preferred crack growth paths between spherulite boundaries in the N-100J1 sample which has the lowest creep failure time. In the N-100J2 sample, random distribution of lamellas acts as barriers in front of crack which can suppress its growth. This schematic figure shows the important role of crystalline morphology of the polymeric samples against the slow crack growth in accordance with the experimental results obtained from the FNCT test and presented in Table 4. The difference in crystallization kinetics of the samples which causes the formation of different mor phologies is explained in the next section.
has a direct relationship with the arrangement of lamellas. As results of this study confirm, although isothermal crystallization cannot be a complete substitution for slow crack growth test as each offer their own unique findings, but it can be a much more straightforward and less time-consuming procedure for indirect determination of resistance to creep compared to slow crack growth test, which can make it a useful and fast complementary test method for evaluation of creep perfor mance in semi-crystalline polymers. Acknowledgment The technical support of this project provided by Jam Petrochemical Company, Iran, is highly acknowledged. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.polymertesting.2019.106128. References
3.3. Comonomer distribution
[1] N. Brown, X. Lu, Y.-L. Huang, R. Qian, Slow crack growth in polyethylene - a review, Makromol. Chem. Macromol. Symp. 41 (1991) 55–67, https://doi.org/ 10.1002/masy.19910410107. [2] L.M. Robeson, Environmental stress cracking: a review, Polym. Eng. Sci. 53 (2013) 453–467, https://doi.org/10.1002/pen.23284. [3] C. Li Pi Shan, J.B. Soares, A. Penlidis, HDPE/LLDPE reactor blends with bimodal microstructures—part I: mechanical properties, Polymer 43 (2002) 7345–7365, https://doi.org/10.1016/S0032-3861(02)00703-6. [4] Y. Pan, X. Gao, Z. Wang, J. Lei, Z. Li, K. Shen, Effect of different morphologies on slow crack growth of high-density polyethylene, RSC Adv. 5 (2015) 28191–28202, https://doi.org/10.1039/C5RA00918A. [5] M. Schilling, M. B€ ohning, H. Oehler, I. Alig, U. Niebergall, Environmental stress cracking of polyethylene high density (PE-HD) induced by liquid media - validation and verification of the full-notch creep test (FNCT), Materwiss, Werksttech 48 (2017) 846–854, https://doi.org/10.1002/mawe.201700065. [6] R.K. Krishnaswamy, Q. Yang, L. F.-B, J.A. Kornfield, Effect of the Distribution of Short-Chain Branches on Crystallization Kinetics and Mechanical Properties of High-Density Polyethylene, 2008, https://doi.org/10.1021/MA070454H. [7] F.P. Alt, L.L. B€ ohm, H.-F. Enderle, J. Berthold, Bimodal polyethylene– Interplay of catalyst and process, Macromol. Symp. 163 (2001) 135–144, https://doi.org/ 10.1002/1521-3900(200101)163:1<135::AID-MASY135>3.0.CO;2-7. [8] X. Lu, N. Brown, A test for slow crack growth failure in polyethylene under a constant load, Polym. Test. 11 (1992) 309–319, https://doi.org/10.1016/01429418(92)90025-7. [9] International Organization of Standards, ISO 16770 - Plastics - Determination of Environmental Stress Cracking (ESC) of Polyethylene - Full-Notch Creep Test (FNCT), 2004. https://standards.globalspec.com/std/2587/iso-16770. (Accessed 6 April 2019). [10] T.R.U. Kratochvilla, H. Muschik, H. Dragaun, Experiences with modified test conditions for notch pipe testing, Polym. Test. 27 (2008) 158–160, https://doi.org/ 10.1016/J.POLYMERTESTING.2007.09.006. [11] A. Frank, G. Pinter, Evaluation of the applicability of the cracked round bar test as standardized PE-pipe ranking tool, Polym. Test. 33 (2014) 161–171, https://doi. org/10.1016/J.POLYMERTESTING.2013.11.013. [12] P. Sardashti, A.J. Scott, C. Tzoganakis, M.A. Polak, A. Penlidis, Effect of temperature on environmental stress cracking resistance and crystal structure of polyethylene, J. Macromol. Sci. Part A. 51 (2014) 189–202, https://doi.org/ 10.1080/10601325.2014.871934. [13] P. Sardashti, C. Tzoganakis, M. Zatloukal, M.A. Polak, A. Penlidis, Rheological indicators for environmental stress cracking resistance of polyethylene, Int. Polym. Process. 30 (2015) 70–81, https://doi.org/10.3139/217.2963. [14] E. Nezbedov� a, P. Huta�r, M. Zouhar, Z. Kn�esl, J. Sadílek, L. N� ahlík, The applicability of the Pennsylvania Notch Test for a new generation of PE pipe grades, Polym. Test. 32 (2013) 106–114, https://doi.org/10.1016/J. POLYMERTESTING.2012.09.009. [15] R.A.C. Deblieck, D.J.M. van Beek, K. Remerie, I.M. Ward, Failure mechanisms in polyolefines: the role of crazing, shear yielding and the entanglement network, Polymer 52 (2011) 2979–2990, https://doi.org/10.1016/J. POLYMER.2011.03.055. [16] M. Eslamian, R. Bagheri, G. Pircheraghi, Co-crystallization in ternary polyethylene blends: tie crystal formation and mechanical properties improvement, Polym. Int. 65 (2016) 1405–1416, https://doi.org/10.1002/pi.5191. [17] J. Fawaz, S. Deveci, V. Mittal, Molecular and morphological studies to understand slow crack growth (SCG) of polyethylene, Colloid Polym. Sci. 294 (2016) 1269–1280, https://doi.org/10.1007/s00396-016-3888-5. [18] E. Baer, L. Zhu, 50th anniversary perspective: dielectric phenomena in polymers and multilayered dielectric films, Macromolecules 50 (2017) 2239–2256, https:// doi.org/10.1021/acs.macromol.6b02669.
To further understand the reasons for different crystallization behavior of used polymers, the SSA-DSC method is used to evaluate the chemical composition and comonomer distribution in the samples. The procedure is a successive annealing of samples in different temperatures and subsequent chain fractionation based on their crystallization ability and the comonomer distribution. This test is mostly sensitive to como nomer distribution rather than difference in molecular weight [32]. Fig. 9(a) shows the heat flow vs. temperature curves obtained from SSA-DSC test. In these curves, there are 4 separated peaks and the area of these peaks indicates the fraction of material that is crystallized in the peak temperature. The lower the temperature peaks, the higher the amount of short chain branches on long molecule chains [30]. In Fig. 9 (b), the area portion of each peak in the samples is compared. It is clear that in the N-100J2, the number of comonomers on long molecules in the structure is higher as indicated from the higher portion of the lower temperature peaks shown in Fig. 9-b. Based on the above statement, it can be concluded that in N-100J2, more comonomers are located on long chains and this reduces long chain mobility and ability to crystallize compared to N-100J1. In N100J2, the thinnest lamellas are formed (Fig. 4) as longer chains are less inclined to participate in lamella formation. Another result of this arrangement of SCBs is increasing number of tie molecules which leads to more resistance of the resin against SCG. On the other hand, in N100J1, a reverse arrangement of SCBs exists, thus thick lamellas are formed and the number of tie molecules in this sample is lower because of higher chain mobility in melt state and preferential placement of longer chains in lamellas rather than acting as tie molecules in inter lamellar areas. 4. Conclusion Based on the results obtained from different tests in this study, there is a direct relationship between the location of short chain branches and the performance of resin against slow crack growth. We can conclude that the placement of short chain branches on longer chains has a direct effect on resistance to slow crack growth in two ways: (1) these co monomers make crystals thinner and increase the number of tie mole cules and chain entanglement in amorphous phase and (2) prevention of spherulite formation in favor of a randomly distributed lamellar morphology which decreases the chance of the crack to find a prefer ential path for growth. Another important observation is the relation between Avrami index n and creep test failure time. With increasing n, the resistance to slow crack growth in resins is reduced significantly because this parameter 7
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Polymer Testing 80 (2019) 106128 [28] S. Deveci, D. Fang, Correlation of molecular parameters, strain hardening modulus and cyclic fatigue test performances of polyethylene materials for pressure pipe applications, Polym. Test. 62 (2017) 246–253, https://doi.org/10.1016/J. POLYMERTESTING.2017.07.007. [29] A.J. Müller, Z.H. Hern� andez, M.L. Arnal, J.J. S� anchez, Successive self-nucleation/ annealing (SSA): a novel technique to study molecular segregation during crystallization, Polym. Bull. 39 (1997) 465–472, https://doi.org/10.1007/ s002890050174. [30] K. Pielichowski, K. Flejtuch, J. Pielichowski, Step-scan alternating DSC study of melting and crystallisation in poly(ethylene oxide), Polymer 45 (2004) 1235–1242, https://doi.org/10.1016/J.POLYMER.2003.12.045. [31] M. Razavi-Nouri, Studies of comonomer distributions and molecular segregations in metallocene-prepared polyethylenes by DSC, Polym. Test. 25 (2006) 1052–1058, https://doi.org/10.1016/J.POLYMERTESTING.2006.07.004. [32] M.Y. Keating, E.F. McCord, Evaluation of the comonomer distribution in ethylene copolymers using DSC fractionation, Thermochim. Acta 243 (1994) 129–145, https://doi.org/10.1016/0040-6031(94)85048-8. [33] R.A. Shanks, G. Amarasinghe, Comonomer distribution in polyethylenes analysed by DSC after thermal fractionation, J. Therm. Anal. Calorim. 59 (2000) 471–482, https://doi.org/10.1023/A:1010141508932. [34] T. Wu, L. Yu, Y. Cao, F. Yang, M. Xiang, Effect of molecular weight distribution on rheological, crystallization and mechanical properties of polyethylene-100 pipe resins, J. Polym. Res. 20 (2013) 271, https://doi.org/10.1007/s10965-013-0271-9. [35] A. Lustiger, C.N. Marzinsky, R.R. Mueller, Spherulite boundary strengthening concept for toughening polypropylene, J. Polym. Sci., Part B: Polym. Phys. 36 (1998) 2047–2056, https://doi.org/10.1002/(SICI)1099-0488(19980915)36: 12<2047::AID-POLB4>3.0.CO;2-U. [36] G.-X. Wei, H.-J. Sue, J. Chu, C. Huang, K. Gong, Toughening and strengthening of polypropylene using the rigid–rigid polymer toughening concept Part I. Morphology and mechanical property investigations, Polymer 41 (2000) 2947–2960, https://doi.org/10.1016/S0032-3861(99)00454-1.
[19] B.-H. Choi, J. Weinhold, D. Reuschle, M. Kapur, Modeling of the fracture mechanism of HDPE subjected to environmental stress crack resistance test, Polym. Eng. Sci. 49 (2009) 2085–2091, https://doi.org/10.1002/pen.21458. [20] A. Sarafpour, G. Pircheraghi, R. Rashedi, K. Afzali, Correlation between isothermal crystallization and morphological/rheological properties of bimodal polyethylene/ carbon black systems, Polym. Cryst. 1 (2018), e10014, https://doi.org/10.1002/ pcr2.10014. [21] G. Pircheraghi, R. Rashedi, K. Afzali, Correlation between Rheological and Mechanical Properties of Black PE100 Compounds-Effect of Carbon Black masterbatchThe Substitution of Carbon Black with Optimized Black Masterbatch in the Production of CRP100 Black (PE100 Black) View Project Polyurethane/CNT Nanocomposites View Project, (n.d.). doi:10.3144/expresspolymlett.2017.60. [22] J.J. Cheng, M.A. Polak, A. Penlidis, Influence of micromolecular structure on environmental stress cracking resistance of high density polyethylene, Tunn. Undergr. Space Technol. 26 (2011) 582–593, https://doi.org/10.1016/J. TUST.2011.02.003. [23] A.T. Lorenzo, M.L. Arnal, J. Albuerne, A.J. Müller, DSC isothermal polymer crystallization kinetics measurements and the use of the Avrami equation to fit the data: guidelines to avoid common problems, Polym. Test. 26 (2007) 222–231, https://doi.org/10.1016/J.POLYMERTESTING.2006.10.005. [24] J.N. Hay, Application of the modified avrami equations to polymer crystallisation kinetics, Br. Polym. J. 3 (1971) 74–82, https://doi.org/10.1002/pi.4980030205. [25] L. Mandelkern, F.A. Quinn, P.J. Flory, Crystallization kinetics in high polymers. I. Bulk polymers, J. Appl. Phys. 25 (1954) 830–839, https://doi.org/10.1063/ 1.1721753. [26] E. Piorkowska, A. Galeski, J.-M. Haudin, Critical assessment of overall crystallization kinetics theories and predictions, Prog. Polym. Sci. 31 (2006) 549–575, https://doi.org/10.1016/J.PROGPOLYMSCI.2006.05.001. [27] J. William, R. Mehl, Reaction kinetics in processes of nucleation and growth, Trans. Met. Soc. AIME. 135 (1939) 416–442.
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Material Behaviour
Correlation between isothermal crystallization properties and slow crack growth resistance of polyethylene pipe materials Farzad Gholami a, Gholamreza Pircheraghi a, *, Reza Rashedi b, Abdulhanan Sepahi b a b
Polymeric Materials Research Group (PMRG), Department of Materials Science and Engineering, Sharif University of Technology, Tehran, Iran Department of Polymer Engineering and Color Technology, Amirkabir University of Technology, Tehran, Iran
A R T I C L E I N F O
A B S T R A C T
Keywords: Isothermal crystallization Avrami index Morphology Slow crack growth SSA-DSC
In this study, three PE100 polyethylene materials with fairly different short chain branch distributions were used to characterize the relation between thermal properties and creep test failure time. The samples were thermally characterized using isothermal and non-isothermal differential scanning calorimetry (DSC). The three resins showed different behavior after fitting on the Avrami equation. N-100J2 sample, in which short chain branches (SCBs) are located on longer molecule chains, has a lower chain mobility and higher crystallization time, while N-100J1 has the opposite crystallization properties. Also, these samples had different Avrami index, n. Micro structural evaluation of these samples revealed that the Avrami index had a direct relationship with the morphology of samples and the higher n correlates to formation of more spherulite structures. The creep test result showed that the sample with a lower crystallization rate and lower Avrami index was more resistant against slow crack growth (SCG) because of its crystalline morphology with the obstacles in crack growth. There is a meaningful relationship between lamellae arrangement and failure time; when a crystalline segment in resin has more spherulite structures, higher Avrami index, the boundary of these spherulites may act as weak regions in front of crack and consequently lead to easier crack growth in such material.
1. Introduction Slow crack growth failure in polyethylene pipes has attracted considerable interest among researchers due to economic and safety reasons. Today, most failure in polyethylene pipes, which are widely in transporting water and gas, is due to slow crack growth (SCG) mecha nism [1–5]. It is quite common that the placement of comonomers on low or high molecular weight chains can significantly affect the me chanical properties of pipes. Selective localization of short chain branches (SCBs) on higher molecular weight fractions or longer chains increases the number of entanglements between molecules, which cau ses the resin to form thinner and fewer crystals with a huge number of tie molecules. As a result, these tie molecules make the resin tougher and increase the resistance of material against slow crack growth [6]. This concept was used in the development of bimodal polyethylene resins with specific short chain branch distribution, which is now widely known as PE100 pipe grade polyethylene [7]. To determine resistance to SCG, several methods and tests are stan dardized; most familiar means of characterizing creep behavior in polyethylene resins are Pennsylvania Edge Notch Test (PENT) [8], Full
Notch Creep Test (FNCT) [9], Notched Pipe Test (NPT) [10], and Crack Round Bar test (CRB) [11]. Yet, all these methods suffer from the same problem of being financially ineffective and time consuming, some of which taking even months to be completed. Few researchers [6,12,13] have addressed the use of other tech niques to determine the lifetime of polyethylene (PE) resins. Considering that crystalline structure in semi-crystalline polymers has a significant influence on mechanical properties, there is a direct relation between the morphology of lamellae and long-time performance of polyethylene resins [14–21]. Thickness and arrangement of lamellae are the two most important parameters that have the most effect on the mechanical properties of PE pipes. Thinner crystalline segments happen to have more tie molecules in between, which increases entanglement density in the amorphous phase and makes the material more resistant against craze propagation during slow crack growth [15,19,21,22]. On the other hand, the crystallinity percentage and arrangement of crystalline seg ments are critical factors that have a direct effect on mechanical prop erties of material, such as stiffness and yield strength [6,16]. The general factor that has a direct effect on thickness and arrangement of lamellae is crystallization kinetics influenced by polymer chain topology [20,
* Corresponding author. E-mail address: [email protected] (G. Pircheraghi). https://doi.org/10.1016/j.polymertesting.2019.106128 Received 3 June 2019; Received in revised form 31 August 2019; Accepted 26 September 2019 Available online 27 September 2019 0142-9418/© 2019 Published by Elsevier Ltd.
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2. Experimental section
Table 1 Coding system and properties of used polyethylenes. Samples Code
Polymer Type
MFI [190� C-5 kg] (g/10min)
Density (g/cm3)
N-100J1 N-100J2 N-100 M
PE100 PE100 PE100
0.22 0.18 0.21
0.956 0.950 0.954
Three different PE100 pipe grade polyethylene granules were sup plied by Jam and Marun petrochemical companies in Iran. All samples were produced through LyondelBasell’s Hostalen process using ZieglerNatta catalyst but with different polymerization conditions which can affect the comonomer distribution in final product. All samples use 1Butene comonomer and their codes, MFI and density are presented in Table 1. For sample coding, letter "N" refers to neat PE100 powder being used, letter "J" refers to samples produced by Jam Petrochemical com pany and "M" for sample produced by Marun Petrochemical company. To characterize the crystalline phase and to evaluate thermal prop erties of the samples, isothermal and non-isothermal crystallization of samples were studied by TA Instruments Q100 differential scanning calorimeter (DSC). Approximately 5 mg of the samples were weighed to carry out the test. In the non-isothermal test, samples were heated and cooled at 10 � C/min rate from 25 � C to 180 � C and this cycle was repeated three times to evaluate the crystallization and melting tem peratures as well as the degree of crystallization. For the isothermal test, samples were heated to 180 � C at 10 � C/min and kept for 3 min to ensure complete melting. Then, they were first cooled to 140 � C at 10 � C/min, held at 140 � C for 2 min, and were then rapidly cooled down (at 40 � C/ min) to and maintained at 121 � C for 2 h. The procedure which was previously used by Krishnaswamy et al. [6], made it possible to evaluate and compare crystallization time and rate of the samples. DSC test was conducted in a pure nitrogen atmosphere. After isothermal crystalliza tion tests, the obtained data were fitted into Avrami equation using a plugin for Origin® software developed by A. T. Lorenzo et al. [23]. Fractionation of molecular chains in the different PE100 compounds were performed using Successive self-nucleation and annealing (SSA) test. Unlike TREF and CRYSTAF tests, SSA is less time consuming and more cost-effective. In this test, different molecules in the polymer are separated from each other to form unique crystals with regards to their ability to crystallize. In this procedure samples were melted first at a temperature higher than the melting point of the resin, then cooled to room temperature in order to eliminate thermal history of samples. Afterward, for fractionation, the samples were heated to a temperature 8� C lower than previous one and maintained in that temperature for 15 min. This procedure was repeated for 4 steps and eventually, the sample was heated to 180� C and a multipeak DSC heating curve could be obtained [29]. In previous studies, SSA-DSC has been used for frac tionation of LLDPE which has a large number of short chain branches. Thus, the separation process of chains is done easily in a short time (5 min) and with small temperature steps(5� C) [29–33]. In this work, the PE100 samples have less than 1% comonomer in their structure and it is necessary to maintain the samples in annealing temperatures for longer times and larger temperature steps to ensure that the fraction ation of the chains occurs properly. After several trial and error tests, the optimum value for annealing time and temperature steps were found to be 15 min and 8 � C, respectively.
23–25]. Several researchers proposed [24–27] theoretical equations for investigation and prediction of crystallization kinetics, among which, the Avrami equation is the most common and practical theory for determination and prediction of polymer’s crystallization kinetics [23, 26]. Sarafpour et al. [20] used the Avrami equation and rheological measurements to predict the behavior of black polyethylene pipe with different masterbatches. They found that by decreasing molecular weight of carbon black carrier polymer, three-dimensional growth of crystal segments could change the crystalline structure and the pipe’s long term performance. Krishnaswamy et al. [6] have studied the in fluence of short chain distribution on mechanical and rheological properties of polyethylene resins. They used two different resins with different location of 1-Butene comonomers on long and short chains. In this investigation, isothermal crystallization is used to survey crystalli zation kinetics and then to study the effect of short-chain branches on mechanical and microstructural characteristics of the samples. They found that by locating SCBs on long chains, molecules were remarkably slower in crystallization, resulting in formation of more tie molecules and subsequently enhanced resistance to slow crack growth. Sardashti et al. [13] investigated the long-time performance of PE pipes with rheological test. They used relaxation time obtained from Carreau–Ya soda equation fitting on rheological measurement data and proposed that by increasing this factor, the chain entanglement would be increased, enhancing resistance to SCG as a result. Deveci et al. [28] used strain hardening test, crack round bar (CRB) test, and notched pipe test to study the relation between molecular weight and molecular weight distribution of PE resins and resistance to SCG phenomenon. They have found that to determine the effect of comonomer type, strain hardening test is better than CRB method. On the other hand, to detect the effect of other parameters like molecular weight, molecular weight distribution and comonomer content, both methods are highly efficient. The aim of this work has been to enhance the current knowledge about the relationship between polymer chain microstructure, forma tion, and morphology of crystalline domains and resistance to slow crack growth in polyethylene pipe materials. In this study, we have used PE100 polyethylene resins with different SCB distribution and almost identical average molecular weight. Slow crack growth has been assessed by means of full notched creep test (FNCT) while the crystalline characteristics has been analyzed using DSC measurements and FE-SEM imaging.
Fig. 1. (a) Schematic and (b) picture of the designed Creep test machine. 2
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Fig. 2. (a) Heating and (b) cooling curve of PE100 samples in non-isothermal DSC test. Table 2 Thermal properties of PE100 samples. sample
Tm ½� C�
Tc ½� C�
Xc ½%�
N-100J1 N-100J2 N-100 M
131.1 129.6 131.6
117.3 115.3 116.7
61.2 53.0 57.9
To evaluate the lifetime of the samples and their creep resistance, a lab scale full notch creep test (FNCT), designed according to ISO 16770 [9] standard, was performed. The instrument and a representative schematic are shown in Fig. 1. Samples with dimensions of 6 � 6 � 9 mm were cut from 6 mm compression molded sheets. The test was carried out at 55 � C with a dead load stress of 9 MPa. Average failure time of three specimens for each sample was measured. The crystalline morphology and lamellar structure of the samples were observed using field emission scanning electron microscope (FESEM). For this purpose, compresseion molded bars were fractured in liquid nitrogen. To remove the amorphous phase of samples, an acid solution (H2SO4:H3PO4: water ¼ 5:2:1 and 1% w/n KMnO4) was used as etchant, followed by washing in H2O2:H2SO4 (dilute) (H2SO4: H2O ¼ 2:7). The etching method was suggested by Eslamian et al. [16]. Fractography tests from creep samples were also carried out by FE-SEM.
Fig. 3. Heat-flow curve vs. time for PE100 samples obtained from isothermal crystallization. Table 3 Obtained parameters from curve fitting on Avrami equation.
3. Results and discussion 3.1. Crystallization kinetics and morphology DSC cooling and heating curves of three samples are shown in Fig. 2 and thermal properties are given in Table 2. It can be deducted from different results in Table 2 that the structure of these resins is not the same. Due to the utilization of identical Ziegler-Natta ethylene poly merization method in the presence of 1-Butene for production of sam ples, one can assume that the difference between samples is originated from different comonomer distribution and content. The lower crystal linity and melting point of N-100J2 can be attributed to its higher commoner content and evenly distributed 1-Butene in long chains [6,16, 33]. Also, the crystallization temperature, TC, of this sample is the lowest among the samples due to difficulty in the folding of its chain in the melt state. Isothermal crystallization curves of the three PE100 samples are shown in Fig. 3. From these curves, it is obvious that in N-100J2 sample, the necessary time for crystallization is higher than other samples and N100J1 crystallization time is the lowest. To quantify the results, these
Samples
t1/2(s)
K (Crystallization rate) (min-n)*
n (Avrami index)
N-100J1 N-100J2 N-100 M
48 158 60
0.340 0.109 0.249
2.88 2.45 2.80
experimental data were fitted to the Avrami equation (Eq. (1)). 1
Vc ðtÞ ¼ expð
Ktn Þ
(1)
in which Vc is relative crystalline fraction, K is crystallization rate con stant and n is Avrami index, which should have a value between 2 and 3 in crystallization of polymer, and t is the time. Obtained parameters are presented in Table 3. From Table 3, it can be inferred that N-100J2 chains are more sluggish than other resins and need more time for crystallization [6,34]. This causes a significant change in crystalline morphology and me chanical properties of materials which can have a direct impact on long-term performance since the rate of crystallization have a correla tion with chain entanglements in the structure of the materials [4,6,20, 21]. Another important parameter in this equation is n, which is related 3
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Fig. 4. FE-SEM images of crystalline structures of the samples at three magnifications (a) N-100J1, (b) N-100J2 and (c) N-100 M.
to the morphology of crystals. This parameter varies between 2 and 3 and the closer n gets to 3, the more number of spherulites can be formed, while with approaching 2, crystalline structure leans towards an unor dered lamellar morphology [20]. To verify the results of isothermal parameters as well as fitting of Avrami equation, etched samples were examined by FE-SEM to observe the morphology of samples. Fig. 4 shows lamella order of three PE100 samples in different magnifications. Spherulite morphology can be observed in N-100J1 and N-100 M (red circles), but in N-100J2 crys talline lamellas have a random placement. These images are in agree ment with the result of Avrami fitting that N-100J1 with higher Avrami index (2.88) has more spherulites in its structure while in N-100J2, which has lower n, there are no spherulites. This is in consistency with Sarafpour et al. [20] results. Observed morphologies are from hot pressed plates which are cooled non-isothermally, thus they may not exactly reveal the morphologies from isothermal crystallization, but they are useful for a qualitative comparison of different morphologies in crystallization of samples. It can be seen that N-100J2 has thinner lamellas and distance be tween these lamellas are higher than two other samples. Herein, one can assume that there were some tiny lamellas in the amorphous phase which could act as tie-crystals and have been removed by the etching
Table 4 Failure time of PE100 samples obtained from creep test. Samples
Failure time (h)
N-100J1
50 � 12
N-100J2
460 � 36
N-100 M
104 � 15
process as reported elsewhere [16]. Thinner lamellas between main lamellas of the N-100J2 sample can cause a significant increase in the number of tie molecules and probably make this resin tougher against SCG phenomenon. 3.2. Resistance against slow crack growth Results of the creep test by FNCT apparatus are presented in Table 4. It is clear that longterm properties of N-100J2 sample are significantly higher compared to other samples, showing excellent creep perfor mance. The main mechanism of the polymer failure in this test is slow crack growth, which is highly sensitive to the craze formation in the 4
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in all samples has the same brittle morphology and there is no noticeable difference between these three samples. In two other zones of the craze, growth is higher, and the surface is relatively rougher than the first zone. In the second zone in which ligament area decreases, some pull out re gions can be seen. These regions are fibrillations that are formed during craze growth and after failure, these signs remain. In the third zone, in which the load bearing area is very small and then stress is higher than other zones, the velocity of craze growth is significantly higher and the fracture surface is not brittle. By comparing the morphology of samples in these two zones, we can investigate the mechanism during creep test. In Fig. 6, it can be seen that zone II and III in N-100J2 are rougher compared to two other samples, indicating that in this sample, crazes could not grow fast because of the resistance of entanglement of polymer chains in amorphous region, and hence the formation of strong fibrils as
Fig. 5. (a) An overview of the fracture surface of creep test samples with three distinct regions with different morphologies. (b) Schematic models for explaining how different regions in the fracture surface are formed. With further progress of the crack in the samples, the stress level is increased and then larger fibrillations formed; as a result, the third zone is rougher than the first zone in the creep test fracture surface.
amorphous phase and fibrillation of polymer in the vicinity of formed crazes. Therefore, the crystalline morphology of the sample has a vital role in the craze growth and resistance against SCG. For further evaluation of creep behavior of these resins, fracture surfaces of the samples have been studied. After failure in the creep test, all samples show a unique arrangement with three different zones where each zone has its own characteristics. These three zones are shown in Fig. 5. As seen, Zone I is the first zone after notch and has a smooth morphology in all samples, since in this region ligament is large and the applied stress is low, so craze growth in this zone is very slow compared to other zones. Because crack growth rate in the zone I is slow, this zone
Fig. 7. The relation of creep failure time of PE100 samples against Avrami index n.
Fig. 6. Microscopic images of zone II and III of the fracture surface in creep test samples. 5
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Fig. 8. Schematic models of crystalline arrangement in (a) N-100J2 and (b) N-100J1 samples.
observed in this figure. These strong and rough fibrils suppress craze growth and cause more failure time in the FNCT test. On the other hand, in N-100J1 sample, which has the lowest creep lifetime, fibrillations are weaker and smoother, indicating an easier crack growth. Based on results from creep and isothermal crystallization tests, we can see a meaningful correlation between the Avrami parameters such as Avrami index n and crystallization rate K in Table 3 and creep lifetime in Table 4 of samples. Based on these results, the lower the value of the Avrami index and crystallization rate, the higher the failure time in the creep test. Fig. 7 shows the relation between the Avrami index, n, and the failure time in the creep test for three samples. It should be mentioned that although both Avrami index, n, and crystallization rate, k, have meaningful relation with failure time in FNCT test, and it is
possible to redraw the failure curve against n and k, but we believe Avrami index can be more important here as it indicates morphological properties of the samples, while K is merely a crystallization rate indi cator and it can be changed with other parameter like nucleating agents without significant changes in morphological properties. The reason for this correlation between Avrami index and creep failure time can be explained based on different crystalline morphology of the samples. As presented before, the Avrami index of N-100J1 sample is 2.88, which is close to 3, which makes this sample prone to formation of three-dimensional spherulites, as shown in Fig. 4. On the other hand, N-100J2 sample has an Avrami index of 2.45, close to 2, indicating tendency for random growth of two-dimensional lamellas. Considering the different crystalline morphologies of the samples, it is reasonable to assume that in the samples with 3D spherulite
Fig. 9. (a) Heat flow vs. temperature curves of PE100 samples after fractionation by SSA-DSC procedure and (b) peaks area portion of SSA curves. 6
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morphology, the spherulite boundaries play as weak zones where crack can grow easily. On the other hand, in the sample with 2D lamellar structure, there is no preferential path for crack growth. It is reported that when approaching spherulites, the boundaries are weakly entan gled and can be a favorable path for crack growth [35,36]. Fig. 8 shows a schematic representing two different crystalline structures that include two and three-dimensional growth of lamellas from the polymer melt. Fig. 8-a is a graphical representation of N-100J2 sample, while Fig. 8-b is a model of N-100J1 that has a preference for spherulite growth. As indicated in Fig. 8-b, there are preferred crack growth paths between spherulite boundaries in the N-100J1 sample which has the lowest creep failure time. In the N-100J2 sample, random distribution of lamellas acts as barriers in front of crack which can suppress its growth. This schematic figure shows the important role of crystalline morphology of the polymeric samples against the slow crack growth in accordance with the experimental results obtained from the FNCT test and presented in Table 4. The difference in crystallization kinetics of the samples which causes the formation of different mor phologies is explained in the next section.
has a direct relationship with the arrangement of lamellas. As results of this study confirm, although isothermal crystallization cannot be a complete substitution for slow crack growth test as each offer their own unique findings, but it can be a much more straightforward and less time-consuming procedure for indirect determination of resistance to creep compared to slow crack growth test, which can make it a useful and fast complementary test method for evaluation of creep perfor mance in semi-crystalline polymers. Acknowledgment The technical support of this project provided by Jam Petrochemical Company, Iran, is highly acknowledged. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.polymertesting.2019.106128. References
3.3. Comonomer distribution
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To further understand the reasons for different crystallization behavior of used polymers, the SSA-DSC method is used to evaluate the chemical composition and comonomer distribution in the samples. The procedure is a successive annealing of samples in different temperatures and subsequent chain fractionation based on their crystallization ability and the comonomer distribution. This test is mostly sensitive to como nomer distribution rather than difference in molecular weight [32]. Fig. 9(a) shows the heat flow vs. temperature curves obtained from SSA-DSC test. In these curves, there are 4 separated peaks and the area of these peaks indicates the fraction of material that is crystallized in the peak temperature. The lower the temperature peaks, the higher the amount of short chain branches on long molecule chains [30]. In Fig. 9 (b), the area portion of each peak in the samples is compared. It is clear that in the N-100J2, the number of comonomers on long molecules in the structure is higher as indicated from the higher portion of the lower temperature peaks shown in Fig. 9-b. Based on the above statement, it can be concluded that in N-100J2, more comonomers are located on long chains and this reduces long chain mobility and ability to crystallize compared to N-100J1. In N100J2, the thinnest lamellas are formed (Fig. 4) as longer chains are less inclined to participate in lamella formation. Another result of this arrangement of SCBs is increasing number of tie molecules which leads to more resistance of the resin against SCG. On the other hand, in N100J1, a reverse arrangement of SCBs exists, thus thick lamellas are formed and the number of tie molecules in this sample is lower because of higher chain mobility in melt state and preferential placement of longer chains in lamellas rather than acting as tie molecules in inter lamellar areas. 4. Conclusion Based on the results obtained from different tests in this study, there is a direct relationship between the location of short chain branches and the performance of resin against slow crack growth. We can conclude that the placement of short chain branches on longer chains has a direct effect on resistance to slow crack growth in two ways: (1) these co monomers make crystals thinner and increase the number of tie mole cules and chain entanglement in amorphous phase and (2) prevention of spherulite formation in favor of a randomly distributed lamellar morphology which decreases the chance of the crack to find a prefer ential path for growth. Another important observation is the relation between Avrami index n and creep test failure time. With increasing n, the resistance to slow crack growth in resins is reduced significantly because this parameter 7
F. Gholami et al.
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