Nanomechanical Mapping of a Deformed Elastomer: Visualizing a Self

Jun 28, 2016 - Mapping the structure evolution and mechanical properties of elastic polymers or biomaterials during bulk deformation has been difficul...
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Nanomechanical Mapping of a Deformed Elastomer: Visualizing a Self-Reinforcement Mechanism Shuquan Sun,† Dong Wang,*,† Thomas P. Russell,∥,‡,§ and Liqun Zhang*,† †

State Key Laboratory of Organic−Inorganic Composites, College of Materials Science and Engineering, Beijing University of Chemical Technology, Beijing 100029, China ‡ Polymer Science and Engineering Department, University of Massachusetts, Amherst, Massachusetts 01003, United States § Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States ∥ Beijing Advanced Innovation Center for Soft Matter Science and Engineering, Beijing University of Chemical Technology, Beijing 100029, China S Supporting Information *

ABSTRACT: Mapping the structure evolution and mechanical properties of elastic polymers or biomaterials during bulk deformation has been difficult, yet this information has long been thought to be key for understanding the structure− mechanical property relationship necessary to guide the design of new materials. Here we use a nanomechanical mapping to assess the structural evolution and mechanical properties of a deformed isoprene rubber (IR) to elucidate a self-reinforcement mechanism in this material. A hierarchical nanofibrillar structure, ranging from several to a hundred nanometers in size, comprised of fibers oriented parallel to the stretching direction was found. The nanofibers, connected by oriented amorphous tie chains, form a network structure that is responsible for significantly enhanced stress, a key factor giving rise to the self-reinforcement of IR and, more than likely, most elastomers that undergo strained-induced crystallization.

been used to characterize SIC where further structural properties have been determined. However, a direct visualization of the evolution of the microstructure and the change in the mechanical properties of SIC during bulk deformation have not been assessed, which is important in understanding and developing self-reinforced materials. Atomic force microscopy (AFM) is a versatile technique for probing the structure and properties of materials at the surface with high spatial resolution and with the capacity to operate under a range of different conditions.28 Recently, AFM nanomechanical mapping measurements, including force volume (FV),29−31 harmonic force,32 band excitation,33 contact resonance,34 peak force,35 and multifrequency force microscopies,36 have emerged and are increasingly being used for the simultaneous determination of the microstructure and mechanical properties of a material. As of yet, though, measurements during the application of an external force have been lacking, with the exception of tapping mode AFM.37,38 Watabe et al. observed line structures several micrometers in width along the strain direction and found a strain localization and stress concentration in natural rubber

Elastomers such as natural and synthetic rubbers play a very important role in automotive, aerospace, construction, semiconductor manufacturing, consumer products, packaging, healthcare, and medical industries.1−4 The immense use of elastomers, especially natural rubber (NR)-based elastomers, can be attributed to their outstanding properties, including extraordinarily large elongations at break, sharply increasing modulus at high strains, low hysteresis, and high strengths combined with fatigue and tear resistance.5,6 The superior mechanical properties of NR occur by a self-reinforcement in the absence of fillers.7,8 Consequently, understanding the mechanism of this self-reinforcement, in rubbers like NR, isoprene rubber (IR), and butyl rubber (IIR), is important for designing and preparing advanced materials. It is widely accepted that strain-induced crystallization (SIC) of NR or other elastomers is a key factor contributing to its superior mechanical properties, in particular crack growth resistance and fatigue behavior.5−11 The crystallization process dissipates strain energy, and the induced crystallites reinforce, significantly increasing fracture resistance. SIC, first discovered in 1925 by Katz using X-ray diffraction,12 has since been studied extensively using static and in situ synchrotron-based X-ray diffraction where the size, orientation, volume fraction, and unit cell parameters of the crystals have been determined.13−19 Dilatometry,9 electron microscopy,20−22 optical birefringence,23,24 infrared absorption,25 and NMR26,27 have also © XXXX American Chemical Society

Received: April 8, 2016 Accepted: June 27, 2016

839

DOI: 10.1021/acsmacrolett.6b00278 ACS Macro Lett. 2016, 5, 839−843

Letter

ACS Macro Letters vulcanizates at large strains by height and phase AFM tappingmode images and concluded the breakdown of the affine deformation assumption.37 To understand SIC and the resultant self-reinforcement, though, it is essential to perform studies as the sample is strained, so that the evolution of the morphology and an understanding of the structure−mechanical property relationship can be determined to provide a guide for the design of advanced materials. In the present work, AFM nanomechanical mapping was used to map the structure evolution and mechanical properties of dicumyl peroxide (DCP) cross-linked polyisoprene rubber (IR) under deformation. IR was chosen since its molecular structure is similar to that of NR but without impurities, like proteins, fatty acids, conjugated phospholipids, and gelators.1 AFM nanomechanical mapping results show the formation of a multilength scale fibrillar network structure that is responsible for the tough mechanical properties of the deformed IR. A typical stress−strain curve of IR is shown in Figure 1. The stress slowly increases during the initial stages of deformation

Figure 2. AFM stress images of IR as a function of strain: (a) 0%, (b) 100%, (c) 200%, (d) 300%, (e) 400%, (f) 500%, (g) 600%, and (h) 700%. The strain is applied in the vertical direction of these images. (i) Typical force-deformation curves (blue) of the samples and the curve fitting against JKR contact (pink). The upper and lower curves represent the IR without deformation and IR at strain of 500%, respectively (blue dots in (a) and (f)). (j) Section analysis on the AFM stress image of (f). Figure 1. Stress−strain curves of the vulcanized IR and SBR.

deformed IR at strains of 500% is responsible for the sharp increase in the stress−strain curve in Figure 1. Further increasing the strain to 600 and 700% (it should to be noted that the 700% strain sample (Figure 2g) was obtained at a slow stretching speed, and the strain was ∼640% at a stretching speed of 500 mm/min during tensile testing) results in innumerable fibrils that cover 20−32% of the area (Figure 2h). Moreover, these fibrils aggregate parallel to the stretching direction to form bundles having diameters ranging from several to one hundred nanometers, which causes the significant increase in the stress. Figure 2a−2d shows the structural evolution from a homogeneous single phase to one with an oriented band-like morphology, indicating the polymer chains undergo a transition from a random orientation to a preferential orientation along the stretching direction upon deformation. The size of the band-like structures, oriented amorphous regions, ranges from 50 to 120 nm normal to the direction of strain. These oriented amorphous regions increased with stretching and formed a coherent framework in the stretching direction by connecting with each other. Subsequently, crystallization occurs in these oriented amorphous regions, as evidenced by the appearance of fibers (Figure 2e). The emergence of the fibrils was found to correspond to the onset of SIC. Crystallites grow from the oriented network, and some of the crystallites connect with each other along the direction of stretching, resulting in a fibrillar rather than a lamellar structure. The increase in the number of fibrils after the onset of SIC corresponds to the

and, then, sharply increases with further deformation until the sheet fractures. A noticeable upturn in stress is observed at a strain of ca. 500%, corresponding to the onset strain of SIC, where crystallites act as a reinforcing filler. The stress−strain curve of SBR, on the other hand (Figure 1), increases slowly, and no upturn is observed until fracture. The marked difference between IR and SBR indicates a substantial change in the microstructure of IR during deformation. Figure 2 shows a series of AFM stress maps of the deformed IR at different strains. The color scale bar (from red to yellow) corresponds to an AFM stress increase from small to high values, respectively. The image shown in Figure 2a is the AFM stress map of the IR before deformation that is featureless, as would be expected. As the strain is increased, linear and bandlike regions of high AFM stress in Figure 2b,c, and d can be attributed to a localized deformation-induced orientation hardening since the cross-linking is not perfectly homogeneous. As the strain is further increased to 400% (Figure 2e), the linear and band-like structures disappeared, and a nanofibrillar structure is observed. The most apparent change in the morphology is seen when the strain is increased to 500% (Figure 2f), where a large amount of fibrils, having diameters ranging from several to several tens of nanometers, were observed (Figure 2f). Some fibrils are oriented parallel to the stretching direction in an amorphous matrix, just like the homogeneous dispersion of nanofibers used for polymer reinforcement. The formation of this nanofibrillar structure in 840

DOI: 10.1021/acsmacrolett.6b00278 ACS Macro Lett. 2016, 5, 839−843

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ACS Macro Letters

Figure 3. AFM stress distribution of IR at different strain: (a) 0%, (b) 300%, (c) 500%, (d) 600%, and (e) 700%. (f) Peak AFM stress−strain curve of deformed IR.

increase in the volume fraction of crystals. The diameters of the fibrillar structures, about 10−50 nm, do not change during extension, which agrees with previous WAXD results.16,43 During stretching, the stress is initially carried by the vulcanized (cross-linked) chains, whereas the fibrillar crystallites form a new network that can bear the majority of the applied stress at high strain. Figure 3 shows the histograms of the AFM stress of the deformed IR at different strains. As shown in Figure 3a, the statistical distribution of the AFM stress of the undeformed IR can be described by a Gaussian function with a mean value of 1.7 MPa, consistent with that of bulk values of the IR.1 It should be noted that, as shown in Figure 2i, JKR contact mechanics fits both undeformed and deformed IR at strains of 500%, as evidenced by the superpositioning of the withdrawing force curves. Consequently, PeakForce QNM and the JKR contact mechanics can describe the force curves, and as such, the nanomechanical mapping technique can accurately describe the mechanical properties of the undeformed elastomer and provides valuable mechanical data of the deformed elastomers. With deformation the AFM stress distribution broadens and deviates from the Gaussian function (Figure 3b−e) with the appearance of a shoulder, indicating the formation of some high AFM stress regions. These observations are consistent with the formation of crystallites that increase in number with increasing strain. Figure 3f shows the peak AFM stress as a function of strain for IR, showing an initial increase at the early stages of deformation, followed by a sharp increase in strain over 400%. The trend of the AFM stress increase at the initial stages of deformation (600%), the fraction of the SIC of IR is about 20 wt %, where the fraction of the oriented amorphous phase is about 5 wt %.11,42 The majority of molecules (up to 75 wt %) in the amorphous phase remain unoriented.11,42 Therefore, the application of JKR contact mechanics on the current system can still provide valuable mechanical data, the AFM stress, and therefore help to identify the key structural features of the deformed elastomer. Tensile Tests. The rubber samples were cut into tensile specimens using a punching machine. The cutting die punched the samples into dumbbell-shaped according to ASTM D412. Testing was carried out on tensile testing machine (SANS CMT 4104) at crosshead speed of 500 mm/min at 25 °C. At least five specimens were measured for each sample.

points can be high enough to form crystallites. The number of crystallites increases abruptly after the onset strain of SIC. These crystallites are connected by oriented amorphous tie chains, which form a new network structure that is primarily responsible for the tough mechanical properties. This new network structure is immersed in a large volume of unoriented amorphous chains. The proposed model is analogous to the fiber-reinforced composites, where the amorphous phase acts as the “matrix” and fibrillar crystals act as reinforcing fibers to which the load is transferred during stretching. The aspect ratio of these reinforcing fibers is the main factor that affects the strength of the composites, and the strength increases with increasing aspect ratio. The fibrillar crystals have lager aspect ratios than the band-like oriented amorphous regions. The oriented amorphous regions lead to a slight increase in the stress at low strain, and the formation of a crystalline network leads to a sharp increase in the stress after the onset strain of SIC. In conclusion, we have used a nanomechanical mapping to probe the structural evolution and mechanical properties of deformed IR. We directly visualized the microstructure evolution and demonstrated that the highly deformed IR consists of a hierarchical network of nanofibrils ranging from several to one hundred nanometers in size, which is responsible for the significant enhancement in the stress, leading to the superior mechanical properties. Given the similarity in the molecular structure of natural rubber (NR) to IR, such a network is, more than likely, the origin of the enhanced properties of NR at large strains. Our findings have identified key structural features of the highly deformed IR and provide new insights into the structure−mechanical property relationship that should aid in the rational design of more enhanced mechanical properties of self-reinforcement elastomers by SIC.

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EXPERIMENTAL SECTION

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Materials and Sample Preparation. IR (Mw ≈ 1356 kg/mol, trade name SKI-3S) was obtained from Sterlitamak, Russia. The preparation and cure conditions for the preparation of the rubber sample are shown in Table 1. First, the unvulcanized IR sample was

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsmacrolett.6b00278.

Table 1. Recipes and Cure Conditions of Vulcanized Rubber Samples sample

rubber (part)

DCP (part)

curing time (min)

IR SBR

100 100

1 0.4

60 60

ASSOCIATED CONTENT

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Additional experimental details and figures (PDF)

AUTHOR INFORMATION

Corresponding Authors

masticated by a two-roll mill at 50 °C for 2 min. Dicumyl peroxide (DCP) was then added to the IR sample and thoroughly mixed for 3 min. The rotors were operated at a speed ratio of 1:1.4. The rubber compound was then cured at 150 °C for 60 min in a 1 mm depth mold under 15 MPa pressure. For a comparative study, styrene butadiene rubber (SBR) (Mw ≈ 325 kg/mol, trade name 1502#, Jilin Petrochemical Industrial Corporation) was also prepared as given in Table 1. For AFM experiments, the samples were precut into small strips with a length and a width of 10 and 1.0 mm, respectively. A specially designed sample holder, shown in Scheme 1 in the Supporting Information, was used to hold the IR strips under a symmetric stretched condition. The holder has a diameter of 10 mm and a height of 5 mm. The sample was first stretched to a preset strain value and then was fixed by the two clamps. A steel rod (diameter of 0.3 mm) was used to hold the center of the samples and form a platform for ultramicrotomy. Before stretching, two lines normal to the stretching direction were marked on the surface of the samples. The strain can be

*E-mail: [email protected]. *E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We would like to thank Prof. Ken Nakajima, Dr. So Fujinami, and Dr. Xiaobin Liang for valuable discussions on the JKR analysis and assistance with the experiments. We gratefully acknowledge financial support from the National 973 Basic Research Program of China 2015CB654700(2015CB654704), the Foundation for Innovative Research Groups of the NSF of China (51221002), and the Major International Cooperation(51320105012) of the National Nature Science Foundation of China. 842

DOI: 10.1021/acsmacrolett.6b00278 ACS Macro Lett. 2016, 5, 839−843

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(38) Oderkerk, J.; de Schaetzen, G.; Goderis, B.; Hellemans, L.; Groeninckx, G. Macromolecules 2002, 35, 6623−6629. (39) Hutter, J. L.; Bechhoefer, J. Rev. Sci. Instrum. 1993, 64, 1868− 1873. (40) Wang, D.; Liang, X. B.; Russell, T. P.; Nakajima, K. Macromolecules 2014, 47, 3761−3765. (41) Johnson, K. L.; Kendall, K.; Roberts, A. D. Proc. R. Soc. London, Ser. A 1971, 324, 301−313. (42) Toki, S.; Sics, I.; Ran, S. F.; Liu, L. Z.; Hsiao, B. S.; et al. Macromolecules 2002, 35, 6578−6584. (43) Trabelsi, S.; Albouy, P. A.; Rault, J. Macromolecules 2003, 36, 9093−9099. (44) Martinez, J. R. S.; Cam, J. B. L.; Balandraud, X.; Toussaint, E.; Caillard, J. Eur. Polym. J. 2014, 55, 98−107. (45) Le Cam, J. B.; Toussaint, E. Mech. Mater. 2009, 41, 898−901. (46) 1# No stretching and no cooling process; 2# No stretching. Step 1: the sample was first put into the chamber with −140 °C for 30 min (since we usually finish the microtome process at −140 °C in 20 min); step 2: the sample was brought out from the chamber and kept at 25 °C for 8 h and then was measured by SAXS; 3# Step 1: the sample was first stretched to 500% at 25 °C and then was kept under stretched condition at 25 °C for 30 min; step 2: the sample was put into the chamber with −140 °C for 30 min; step 3: the sample was brought out from the chamber and released from the holder; step 4: the sample was kept at 25 °C for 8 h and then was measured by SAXS; 4# Step 1: the sample was first stretched to 500% at 25 °C and then was kept under stretched condition at 25 °C for 30 min; step 2: the sample was released from the holder and kept at 25 °C for 8 h, and then was measured by SAXS; 5# Step 1: the sample was first stretched to 500% at 25 °C and then was kept under stretched condition at 25 °C for 30 min; step 2: the sample was released from the holder and kept at 25 °C for 8 h; step 3: the sample was put into the chamber with −140 °C for 30 min; step 4: the sample was brought out from the chamber and kept at 25 °C for 8 h and then was measured by SAXS; 6# Step 1: the sample was first stretched to 500% at 25 °C and then was kept under stretched condition at 25 °C for 30 min; step 2: the sample was put into the chamber with −140 °C for 30 min; step 3: the sample was released from the holder and kept at −140 °C for another 30 min; step 4: the sample was brought out from the chamber and kept at 25 °C for 8 h and then was measured by SAXS.

REFERENCES

(1) Gent, A. N. Engineering with Rubber; Hanser Publishers, Oxford University Press: Oxford, England, 1992. (2) Jeong, G. S.; Baek, D. H.; Jung, H. C.; Lee, S. H.; Moon, J. H.; Hong, S. W.; Kim, I. Y. Nat. Commun. 2012, 3, 1−8. (3) Wang, Y.; Ameer, G. A.; Sheppard, B. J.; Langer, R. Nat. Biotechnol. 2002, 20, 602−606. (4) Stoyanov, H.; Kollosche, M.; Risse, S.; Waché, R.; Kofod, G. Adv. Mater. 2013, 25, 578−583. (5) Treloar, L. R. G.; Montgomery, D. J. The Physics of Rubber Elasticity. The physics of rubber elasticity; Clarendon Press, Oxford University Press, 2005. (6) Hamed, G. R. Rubber Chem. Technol. 1991, 64 (3), 493−500. (7) Kmetty, A.; Barany, T.; Karger-Kocsis, J. Prog. Polym. Sci. 2010, 35, 1288−1310. (8) Fukahori, Y. Polymer 2010, 51, 1621−1631. (9) Flory, P. J. J. Chem. Phys. 1947, 15, 397−408. (10) Zhang, H. P.; Niemczura, J.; Dennis, G.; Ravi-Chandar, K.; Marder, M. Phys. Rev. Lett. 2009, 102, 245503. (11) Toki, S.; Hsiao, B. S. Macromolecules 2003, 36, 5915−5917. (12) Katz, J. R. Naturwissenschaften 1925, 13, 410−416. (13) Huneau, B. Rubber Chem. Technol. 2011, 84, 425−452. (14) Toki, S.; Sics, I.; Ran, S. F.; Liu, L. Z.; Hsiao, B. S.; Murakami, S.; Tosaka, M.; Kohjiya, S.; Poompradub, S.; Ikeda, Y.; Tsou, A. H. Rubber Chem. Technol. 2004, 77, 317−335. (15) Candau, N.; Laghmach, R.; Chazeau, L.; Chenal, J. M.; Gauthier, C.; Biben, T.; Munch, E. Macromolecules 2014, 47, 5815−5824. (16) Che, J.; Burger, C.; Toki, S.; Rong, L.; Hsiao, B. S.; Amnuaypornsri, S.; Sakdapipanich, J. Macromolecules 2013, 46, 4520−4528. (17) Ikeda, Y.; Yasuda, Y.; Hijikata, K.; Tosaka, M.; Kohjiya, S. Macromolecules 2008, 41, 5876−5884. (18) Shao, C. G.; An, H. N.; Wang, X.; Jia, R.; Zhao, B. J.; Ma, Z.; Li, X. Y.; Pan, G. Q.; Li, L. B.; Hong, S. M. Macromolecules 2007, 40, 9475−9481. (19) Tosaka, M.; Murakami, S.; Poompradub, S.; Kohjiya, S.; Ikeda, Y.; Toki, S.; Sics, I.; Hsiao, B. S. Macromolecules 2004, 37, 3299−3309. (20) Andrews, E. H. Proc. R. Soc. London, Ser. A 1964, 277, 562−570. (21) Luch, D.; Yeh, G. S. Y. J. Appl. Phys. 1972, 43, 4326−4338. (22) Edwards, B. C.; Phillips, P. J. J. Polym. Sci., Polym. Phys. Ed. 1975, 13, 2117−2127. (23) Mukherjee, D. P. Rubber Chem. Technol. 1974, 47, 1234−1240. (24) Hashiyama, M.; Gaylord, R.; Stein, R. S. Makromol. Chem. 1975, 1, 579−597. (25) Siesler, H. W. Appl. Spectrosc. 1985, 39 (5), 761−765. (26) Nishi, T.; Chikaraishia, T. J. Macromol. Sci., Part B: Phys. 1981, 19 (3), 445−457. (27) Rault, J.; Marchal, J.; Judeinstein, P.; Albouy, P. A. Eur. Phys. J. E: Soft Matter Biol. Phys. 2006, 21 (3), 243−261. (28) Garcia, R.; Magerle, R.; Perez, R. Nat. Mater. 2007, 6, 405−411. (29) Tsukruk, V. V.; Sidorenko, A.; Gorbunov, V. V.; Chizhik, S. A. Langmuir 2001, 17, 6715−6719. (30) Cappella, B.; Kaliappan, S. K. Macromolecules 2006, 39, 9243− 9252. (31) Wang, D.; Fujinami, S.; Nakajima, K.; Nishi, T. Macromolecules 2010, 43, 3169−3172. (32) Sahin, O.; Magonov, S.; Su, C. M.; Quate, C. F.; Solgaard, O. Nat. Nanotechnol. 2007, 2, 507−514. (33) Jesse, S.; Kalinin, S. V.; Proksch, R.; Baddorf, A. P.; Rodriguez, B. J. Nanotechnology 2007, 18, 435503. (34) Stan, G.; Cook, R. F. Nanotechnology 2008, 19, 235701. (35) https://www.bruker.com/fileadmin/user_upload/8-PDF-Docs/ SurfaceAnalysis/AFM/Brochures/PeakForce_Quantitative_ Nanomechanical_Property_Mapping_b.pdf. (36) Herruzo, E. T.; Perrino, A. P.; Garcia, R. Nat. Commun. 2014, 5, 3126. (37) Watabe, H.; Komura, M.; Nakajima, K.; Nishi, T. Jpn. J. Appl. Phys. 2005, 44, 5393−5396. 843

DOI: 10.1021/acsmacrolett.6b00278 ACS Macro Lett. 2016, 5, 839−843