Quantitative Nanomechanical Investigation on Deformation of Poly

Article pubs.acs.org/Macromolecules

Quantitative Nanomechanical Investigation on Deformation of Poly(lactic acid) Hao Liu, Na Chen, So Fujinami, Dmitri Louzguine-Luzgin, Ken Nakajima,* and Toshio Nishi WPI-Advanced Institute for Materials Research (WPI-AIMR), Tohoku University, Sendai 980-8577, Japan ABSTRACT: This work deals with the tensile drawing behavior of poly(lactic acid) (PLA) in the amorphous state at room temperature. Nanomechanical mapping method is applied to obtain the submicrometer morphology together with mechanical properties using atomic force microscopy (AFM). By carefully collecting the data of PLA specimens deformed with different elongation, we managed to monitor the microscopic structural/mechanical evolution simultaneously, and thus, the microscopic deformation mechanism was studied phenomenologically. The PLA matrix shows a highly heterogeneous response to the external stress. The stress is concentrated in some preferential places, thereby leading to high local Young’s modulus and giving rise to crazes perpendicular to the tensile direction. However, in the position away from the crazes, the matrix undergoes no stress and shows similar Young’s modulus as neat PLA specimens. The inhomogeneous stress distribution results in the catastrophic fracture of the matrix at an early state, which could explain the brittle nature of PLA at room temperature.

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INTRODUCTION Poly(lactic acid) (PLA) is a kind of aliphatic polyester with promising capabilities for the environmental concerns. It is renewable in nature and can be biodegraded into nontoxic lactic acid, which has been widely studied for the substitution of oil-based plastics.1−3 Nevertheless, the application of PLA had been limited to biomedical products due to the relatively high production cost compared with polyolefin. Owing to the recent developments in production techniques, PLA is becoming economically viable for commodity application.4,5 However, PLA cannot be used as the substitute of commodity plastics for its relatively low strength and resilience.6,7 The mechanical properties and chemical stability of a polymer, in general, strongly depend on its morphology and crystal structure.8 PLA is a semicrystalline polymer with slow crystallization rate at room temperature. Under normal processing conditions, it can be quenched into the glassy state. For the plastic deformation of amorphous polymers, the main mechanisms observable on a microscopic scale are crazing and shear yielding. While further crystallization might take place during the drawing process, making the mechanism more complicated. Accordingly, differential scanning calorimeter (DSC), wide-angle X-ray diffraction (WAXD), wide-angle Xray scattering (WAXS), polarized Raman spectroscopy, etc., are employed to study PLA under uniaxial drawing at temperatures between the glass transition temperature and melt temperature.9−13 The morphology and crystal structure evolution under different drawing ratios have been well studied by these characterization methods. However, they are incapable of detecting the local mechanical properties. The role of local stress on cavitations and the response of individual chains during plastic deformation process have not been completely © 2012 American Chemical Society

understood. Therefore, it is of great interest to investigate the evolution of submicrometer-scale local mechanical properties for the better understanding of deformation mechanisms. Atomic force microscopy (AFM) might be a good solution to the problem. It is not only a tool to image the topography of solid surfaces, but also offers the potential for imaging materials’ mechanical properties with nanometer-scale resolution.14−17 The idea of probing surface mechanical properties by AFM was proposed decades ago.18,19 Since samples are compressed by the indenting AFM sharp tip hence the elastic response of the sample under loading is collected and analyzed. With the subsequent development on both equipments and analysis methods,20−23 AFM is proved capable of detecting the local surface topography, elastic modulus, adhesive energy, etc., quantitatively on the scale of nanometer. The so-called nanomechanical mapping technology has been applied on the study of polymers, polymer blends, and polymer composites.24−28 It provides a new insight into probing fine details of polymer surfaces and interfaces, which was recently reviewed.29 In this paper, we report the submicrometer-scale mechanical study on the uniaxial drawing of PLA. PLA is supposed to be used at room temperature for commodity applications, accordingly, it is uniaxially tensile deformed at room temperature. Both the submicrometer morphology and on-site microscopic mechanical properties are acquired quantitatively at different drawing ratios. The results are discussed to elucidate the deformation process and fracture mechanisms. Received: September 18, 2012 Revised: October 24, 2012 Published: November 2, 2012 8770

dx.doi.org/10.1021/ma301961h | Macromolecules 2012, 45, 8770−8779

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Article

EXPERIMENTAL SECTION

Material and Preparation. The PLA used for study is purchased from NatureWorks (USA), with the brand name of 4032D. The PLA contains about 2 mol % of D-isomer, it has weight-average molecular weight of 207 kDa and polydispersity of 1.73. PLA samples were first hot pressed into sheets with thickness of 1 mm at 200 °C by a thermo compressor, which was then followed by quenching at 20 °C to form amorphous sheets. The sheets were transparent and isotropic, no relaxation or internal stress could be detected. Macroscopic Mechanical Characterization. Uniaxial tensile deformation was carried out using a Shimazu AG-Xplus tester machine with a load cell of 10 kN. PLA samples were cut into dumbbell shape and tested at a constant crosshead speed of 1 mm/min. The samples were tensile deformed at different ratios to obtain specimens at different deformation stages. Five specimens were tested for each experimental condition. DSC Measurements. DSC was performed on TA Q200 apparatus. The temperature and heat flow were calibrated using high-purity indium samples according to standard procedures. Neat PLA and the craze part of deformed PLA were cut into small pieces and measured. The samples of about 5 mg were put into Tzero aluminum pans and then scanned at a heating rate of 10 °C/min under nitrogen gas flow. The crystallinities were determined by subtracting the area under the cold crystallization peak ΔHc from that of the melting endotherm ΔHm and dividing by the heat of fusion of 100% PLA crystalline ΔH0m, shown in the following equation

χC =

Figure 1. Force−distance curve for adhesive contact.

δ = (z − zc) − (Δ − Δ0);

(2)

The JKR contact is described as the follows.

ΔHm − ΔHc ΔHm0

(Δ0 < 0)

F=

K 3 a − 3wπR − R

δ=

a2 2F + 3R 3aK

6wπRF + (3wπR )2

(3)

(4)

where F, a, R and w are the force, the contact radius, the tip radius and the adhesive energy, respectively. Note that F is equal to kΔ. K is socalled reduced modulus and has the following relationship with Young’s modulus, E

(1)

The heat of fusion used for 100% crystalline PLA is 93 J/g.30 Nanomechanical Mapping. The nanomechanical mapping requires to be characterized on a flat surface. Thus, neat and deformed PLA were cut at −80 °C using a Leica EM UC6 Ultramicrotome. The sections of deformed PLA face on the machine direction were cut to avoid possible relaxation. The cutting direction has an angle of 45° to the tensile direction to distinguish possible effect of knife marks on the results. The measurements were operated in force volume (FV) mode, in which force−distance curves were collected over randomly selected surface areas at a resolution of 128 × 128 pixels. All the measurements were performed on Bruker MutiMode AFM with NanoScope V controller under ambient conditions. The characterization was made immediately after microtome to prevent possible relaxation. Each force−distance curve measurement were carried out at a constant ramp rate of 4.48 Hz using an E scanner and silicon cantilevers with nominal spring constant, k, of 3 N/m. The actual spring constant was measured by thermal tune method, and it changed with different tips. Since the radius of the cantilever tips is important for the result and might change during the FV test, it was measured before and after every FV test by testing the tip-check standard sample. SPIP software was used to calculate the actual tip radius. The trigger threshold of the cantilever deflection was set to 3.0 nm and thus the trigger force was about 9 nN. Same position of PLA specimen was tested twice and no reduce in Yong’s modulus can be found. Therefore, plastically deforming the sample surface can be avoided in the 9 nN trigger force. Force−distance curves were obtained and then analyzed using a procedure developed by our group.31,32 The analysis procedure is based on the Johnson−Kendall−Robert (JKR) contact mechanics.33 Figure 1 shows the schematic force−distance curve for hard (dashed) and soft (solid) materials. First, an abrupt change in cantilever deflection is observed from the point (zc, 0) to the contact point (zc, Δ0), which is called jump-in phenomena. After that, the force increases, crossing the horizontal axis (z0, 0), where the apparent force exerted on the cantilever becomes zero due to the balance between the elastic repulsive force caused by sample deformation and the adhesive attractive force. During the unloading process, a much larger adhesive force is observed beyond the original contact point. Finally the maximum adhesive force is realized and succeeded by a sudden decrease in the contact radius and jump-out of the cantilever. The sample deformation, δ is calculated by the formula

E=

3(1 − ν 2) K 4

(5)

where ν is Poisson’s ratio and that of PLA is 0.36.34 These equations are unable to be converted to an explicit function of δ and F. Accordingly, curve-fitting analysis is not easily applicable. The “twopoint method”, which was proposed by Sun et al.,20 is applied as an alternative solution. The JKR curve is drawn so as to cross the two typical points. One is the point where the attractive force and the repulsive force become equivalent (δ, F) = (δ0, 0), and another is the point where the adhesive force becomes a maximum (δ, F) = (δ1, F1). Using these two points, JKR equations are converted to the following algebraic formula, where K is explicitly represented,

K=

1.27F1 R(δ0 − δ1)3

(6)

Figure 2 shows the data processing of the nanomechanical mapping. Deflection−displacement curves were collected from FV test. The 128 × 128 curves were then converted to force−deformation curves by eq 2, also shown in Figure 2. No sharp jump-in can be found in approaching curve, which might be attributed to possible capillary effect. The moduli were calculated by applying JKR two-point method as described. Although the calculation can be affected by many effects, which has been recently discussed by Tsukruk et al.,29 reliable and quantitative submicrometer-scale results can be acquired if all the experiments and calculation are carefully executed. As shown in the force−deformation curve in Figure 2, the theoretical curve calculated by JKR theory and the experimental data are in good agreement. The “profile broadening” effect due to the tip−sample convolution should be considered in the evaluation of fibril width. Equation 7 is used to correct the true width t ≅ 4 Rr

(7)

where t is the apparent width of the fibrils, r and R are true radii of fibrils and nominal radii of tips. 8771


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Figure 2. Deflection−displacement and force−deformation curves of soft (left) and hard (right) part of PLA. The red curve is the approaching curve and the blue curve is the withdrawing curve. The red line in the force−deformation curve is calculated directly by JKR theory.

Figure 3. Stress−strain curve of PLA.

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RESULTS AND DISCUSSION

fractured PLA. Neat PLA samples are transparent, which makes it easier to observe the structure change during tensile deformation. The delocalization of polymer chains is induced by weakening the structure perpendicular to the loading direction, giving rise to small crazes as shown in the PLA samples under deformation. The crazes tend to appear in the adjacent position and leave large spaces without any crazes, which can be ascribed to the dynamical heterogeneities of amorphous polymers.35 With further deformation, more and larger crazes appear all through the samples, as observed in the post yielding samples. Fracture happens when the crazes can no

Typical tensile stress−strain curve of PLA is shown in Figure 3. PLA is brittle at room temperature and breaks at a small elongation. The macroscopic Young’s modulus is calculated as the slope of the initial linear part and it was 2.8 ± 0.4 GPa. Deformed PLA samples at an elongation of 0.3 mm, 0.7 mm and the fractured samples, marked in Figure 3, were selected for further characterization. Figure 4 shows the photos and microscope graphs of neat PLA, deformed PLA before yielding, after yielding, and 8772


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Figure 4. Photos and microscopic images of neat, before yielding, after yielding and fractured PLA. The orientation of the crazes is perpendicular to the loading direction.

Figure 5. DSC curves of of neat PLA, PLA before yielding, PLA after yielding, and fractured PLA.

longer resist the load. There is no distinct difference between post yielding and fractured samples from the microscopic graphs. DSC measurements were carried out to investigate the crystallinity change due to deformation, as shown in Figure 5. The glass transition temperature Tg, the cold crystallization temperature Tc, the melting point Tm, and the integral of the melting endotherm, ΔHm, together with the crystallinity χc, are collected and summarized in Table 1. The result shows that the crystal behavior might be effected by the deformation process. Neat and before-yielding PLA samples had similar Tc, whereas the obvious decrease of Tc could be observed for postyielding and fractured PLA. The cold crystallization was also enhanced by the deformation, reflected by the increased area of the cold crystallization peak. This phenomena indicates that the chains near crazes are tensed and stress is remained in the polymer matrix after deformation. The remained stress benefits the

Table 1. Summary of the DSC results of PLA neat before yielding post yielding fracture

Tg (°C)

Tc (°C)

Tm (°C)

ΔHm (J/g)

χc (%)

64.8 64.6 64.8 63.9

113.9 114.8 111.7 110.2

166.5 166.6 167.8 167.6

4.48 5.66 7.59 9.23

4.8 6.1 8.2 9.9

rearrangements of polymer chains and results in the enhancements of the cold crystallization.10 It is also possible that local orientation might occurr in the craze region. In order to induce molecular mobility, this local order has to be destroyed, which is reflected in the growing area of exothermic cold crystallization peak in DSC curves.36 Moreover, although PLA was deformed below its Tg, the crystallinity increased slightly after deformation. 8773


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Figure 6. Topography mapping, modulus mapping and modulus histogram of neat PLA.

Figure 7. Topography mapping and modulus mapping of initial craze of PLA. The red box indicates the enlarged area shown in Figure 8. Tensile direction is along y-axis.

The nanomechanical mapping was employed to investigate the microscopic morphology and local mechanical property. Neat specimens were first characterized, as shown in Figure 6. The microscopic Young’s modulus of neat PLA specimen had single peak distribution and the average value of 3.5 ± 0.9 GPa, which was slightly larger than the macroscopic value. Similar result has been also found in Poly(ε-caprolactone) (PCL).27 The interpretation of this phenomenon is that impurities and defects in the polymer matrix might lower the macroscopic Young’s modulus, while these problems could be, to a great degree, suppressed in the microscopic characterization.

When PLA was deformed below its glass-transition temperature, catastrophic localization in terms of craze formation perpendicular to the loading direction occurred as observed in microscope photos. The development of crazes were studied by carefully collecting the nanomechanical mapping results of crazes in the samples before yielding, post yielding and subjected to fracture. Figure 7 shows the data collected at the initial part of a craze. Similar to flexible polymers, Fibrillar structure appeared in the origin of cracks. The fibrillar structures in PLA only occurred in the middle part of stress concentration region. The average modulus of the fibrillate part 8774


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Figure 8. Topography mapping and modulus mapping of the center part of PLA craze.Tensile direction is along y-axis.

Figure 9. Topography mapping and modulus mapping of craze of PLA. Blue and red boxes indicate the enlarged area of center part and edge part respectively shown in Figure 11. Tensile direction is along y-axis.

was 3.0 ± 1.4 GPa, which was slightly lower than neat PLA. Whereas in the case of flexible polymers, PCL for the comparison, the fibril formation occurs all over the area through the polymer matrix and have an average value much higher than neat specimens.27 Figure 8 shows the nanomechanical mapping on the enlarged area at the fibrillar part of PLA crazes. The diameter of PLA fibrils was in the range of 30−60 nm. Moreover, the fibrillar structure in PLA were not strictly along the tensile direction. It might be induced by shear force during deformation. According to Riggleman et al, molecular rearrangements are localized and occur through clusters of molecules moving in union near the stress concentration region during deformation of a polymer glass.37 The mobility is enhanced dramatically in these regions and the local polymer matrix then enters a flow region. The onset of flow leads to the movements of polymer chains and consequently the formation of crazes. With further deformation, crazes developed via the initiation of cavities and finally the coalescence of voids. It is needed to note that the effect of height difference and tip size on the

results must be considered when investigating the modulus of the crazes. The cantilever might be incapable of detecting the bottom part of craze if it is too deep. To avoid this problem, a large number of crazes were characterized and we managed to acquire the reliable submicrometer-scale mechanical properties of craze tips, as shown in Figure 9. The nanomechanical mapping results show that the modulus in the edge of crazes kept on increasing to 6.0 ± 2.5 GPa. Associated with the DSC results, the high modulus could be ascribed to the stress concentration and possible orientation. While in the center part the matrix remains fibrillar structures with a reduced average modulus of 3.0 ± 1.9 GPa, the possible explanation of this phenomenon is that chain scission/slippage must occur during crazing to permit the generation of the fibrillar structure, and modulus is lowered as a consequence. Furthermore, the applied local stress may lead to an increase in segmental mobility,38,39 giving rise to the decreases in viscosity and moduli. The standard derivation of the whole part was increased accordingly, and the whole area in Figure 9 had an average modulus of 4.5 ± 2.6 GPa. It is consistent with finding of Rottler et al.40 that in 8775


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introduced by craze. Deformed PLA shows highly heterogeneous nature of force distribution. With the initial and growth of crazes, the modulus histograme is much widened in both increase and decrease direction. It could be ascribed to the compromise of two folds: one is the stress-induced chain scission and the increase of chain mobility that could reduce the modulus, and the other is the local strain-induced orientation and crystallization that could increase the modulus. The center part and the edge part of crazes were studied in smaller scale. Figure 11 shows the morphology and mechanical properties of these two regions. The typical fibrillar structure orientated to the tensile direction appeared in the center part, which was primarily the consequence of the onset of flow induced by the increased mobility. Modulus mapping image shows that the fibrils are consist of short fibrils, which has been reported by Miller et al.41 Diameter of the fibrils are in the range of 20−40 nm, which is smaller than the fibrils of initial craze. The modulus was reduced due to tearing-induced chain scission and slippage. Furthermore, the center part is the craze fringe, in which direction the craze propagates. The microfibrils are gradually drawn up to their final draw ratio as the craze grows by generation of new fibrils near the craze tip. The fibrils

the craze, the tension is mostely carried by the covalent backbone bonds, and the exponential distribution of tension was found in the craze zone. The overlapped modulus histogrames of neat PLA, initial craze and mature craze is shown in Figure 10 to describe the local modulus change

Figure 10. Moduli histograms of neat PLA, initial craze, and mature craze.

Figure 11. Topography mapping and modulus mapping of the center part (left) and edge (right) of PLA craze. Tensile direction is along the y-axis. 8776


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Figure 12. Topography mapping, modulus mapping, and modulus histogram of before yielding (left) and post yielding (right) PLA.

from neat specimens. The average modulus is 3.3 ± 1.0 GPa and 3.1 ± 1.0 GPa for before yielding and post yielding PLA, respectively, and the modulus histograms are quite similar to that of neat PLA. The PLA matrix is intrinsically heterogeneous material and responds differently to external stress. Even in the fractured samples there are large area of polymer matrix maintaining undeformed. It might be argued that the stress in this part could be relaxed after the load is released. However, since PLA was tensile deformed and characterized at room temperature, which was much lower than its Tg, the relaxation process was supposed to be pretty slow. The modulus values also might be affected by the microtoming process, while according to the nanomechanical mapping of these regions,

breakdown as the consquence of chain scission and slippage in the craze zone42 results in the enlargement of crazes and finally the fracture occurs. On the other hand, the craze edge along tensile directions behave differently. In this part the polymer matrix showed highly orientated structure and had the highest modulus of 4.9 ± 1.5 GPa. Besides, it is possible that this region contributes to the increase in the crystallinity revealed by DSC. Polymer chains staying taut in this region not only increase the modulus but also benefit the chain arrangements and thus the crystallinity is increased.10 The regions away from crazes at different elongations were also investigated by nanomechanical mapping, as shown in Figure 12. The modulus mapping shows no distinct difference 8777


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Table 2. Summary of the Average Modulus in Different Deformation Ratios and Regions of PLA initial craze modulus (GPa)

craze

neat macro

neat micro

before yielding

post yielding

edge

center

edge

center

2.8 ± 0.4

3.5 ± 0.9

3.3 ± 1.0

3.1 ± 1.0

4.0 ± 1.4

3.0 ± 1.4

6.0 ± 2.5

3.0 ± 1.9

Figure 13. Illustration of the deformation of PLA at room temperature.

dispersion in PLA matrix. The middle part in the crazes has the lowest local Young’s modulus of 3.0 ± 1.9 GPa, and that is quite near the macroscopic Young’s modulus which is 2.8 ± 0.4 GPa. It is suggested that the macroscopic Young’s modulus depends on the weakest part of the matrix and the strength and resilience of PLA might be improved if the polymer matrix responds uniformly to the stress.

knife marks had no evident effect on the modulus mapping results. Therefore, these effects can be ignored and the calculated modulus is reliable. The microscopic modulus in different samples and regions were collected and shown in Table 2. These results show the PLA under deformation can be divided into two regions. One is the region near the craze part and the other is where far away from the craze region, illustrated in Figure 13. The craze region undergoes extremely high stress whereas there are regions where no stress is imposed. The brittle nature of PLA can be explained by this phenomenon. The strong heterogeneity is shown at the early deformation stage. The stress tends to be concentrated in random regions, where the local molecular mobility increases dramatically. However, in adjacent regions the molecular mobility remains unchanged. With further deformation, the catastrophic localization takes place in terms of crazes perpendicular to the loading direction, where the stress is concentrated. Afterward, the stress-induced chain scission and slippage give rise to reduced Yong’s modulus and the crazes develope. The fracture happens when the crazes can no longer sustain the increased stress. Meanwhile, the polymer matrix in the regions away from crazes bears no stress and shows similar Young’s modulus as neat specimens. In brief, polymer chains in some regions undergo extremely high stress whereas no stress is imposed in the other regions during the deformation process. Accordingly, the brittle nature of PLA at room temperature can be ascribed to the inhomogeneous stress

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CONCLUSION

PLA was uniaxially tensile deformed at room temperature. The macroscopic Youngs modulus calculate by the slope of initial stress−stain curve was 2.8 ± 0.4 GPa. DSC and AFM was used to characterize PLA at different elongations and thus the deformation mechanism was phenomenologically studied. DSC results suggested that the stress-induced orientation and crystallization might occur during the tensile deformation. The cold crystallization temperature was lowered and the crystallinity was increased from 4.8% for neat PLA to 9.9% for fractured PLA. Moreover, since crazes were formed at places where the stress concentration points located, it is possible that the difference in crystallinity are likely associated with the crazes. Neat PLA and before-yielding PLA had few crazes and share similar Tc and χc. While the decreases in Tc and increases in χc could be observed in postyielding and fractured PLA, in which plenty of crazes appeared. Nanomechanical mapping enabled us quantitatively analyze the submicrometer-scale topography together with the local 8778


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(14) Singamaneni, S.; LeMieux, M. C.; Lang, H. P.; Gerber, C.; Lam, Y.; Zauscher, S.; Datskos, P. G.; Lavrik, N. V.; Jiang, H.; Naik, R. R.; Bunning, T. J.; Tsukruk, V. V. Adv. Mater. 2008, 20, 653. (15) Markutsya, S.; Jiang, C.; Pikus, Y.; Tsukruk, V. V. Adv. Funct. Mater. 2005, 15, 771. (16) Sahin, O.; Erina, N. Nanotechnology 2008, 19, 445717. (17) Sahin, O.; Magonov, S.; Su, C.; Quate, C. F.; Solgaard, O. Nat. Nanotechnol. 2007, 2, 507. (18) Burnham, N. A.; Colton, R. J. J. Vac. Sci. Technol. 1989, 7 (4), 2906. (19) Domke, J.; Radmacher, M. Langmuir 1998, 14, 3320. (20) Sun, Y.; Walker, G. C. Langmuir 2004, 20, 5837. (21) Akhremitchev, B. B.; Walker, G. C. Langmuir 1999, 15, 5630. (22) Magonov, S. N.; Elings, V.; Whangbo, M. H. Surf. Sci. 1997, 375, L385. (23) Tsukruk, V. V.; Sidorenko, A.; Gorbunov, V. V.; Chizhik, S. A. Langmuir 2001, 17, 6715. (24) Magonov, S. N.; Sheiko, S. S.; Deblieck, R. A. C.; Moller, M. Macromolecules 1993, 26 (6), 1380. (25) Shulha, H.; Kovalev, A.; Myshkin, N.; Tsukruk, V. V. Eur. Polym. J. 2004, 40, 949. (26) Wang, D.; Fujinami, S.; Liu, H.; Nakajima, K.; Nishi, T. Macromolecules 2010, 43 (13), 552. (27) Liu, H.; Fujinami, S.; Wang, D.; Nakajima, K.; Nishi, T. Macromolecules 2011, 44, 1779. (28) Sun, Y. J.; Walker, G. C. Langmuir. 2005, 21 (19), 8694. (29) McConneya, M. E.; Singamanenia, S.; Tsukruk, V. V. Polym. Rev. 2010, 50 (3), 235. (30) Fischer, E. W.; Sterzel, H. J.; Wegner, G.; Kolloid., Z. Z. Polymer 1973, 251, 980. (31) Nakajima, K. In Polymer Physics from Suspensions to Nanocomposites and Beyond; John Wiley & Sons: Hoboken: NJ, 2010; Chapter 3, p 129. (32) Nakajima, K. In Current Topics in Elastomers Research; Bhowmick, A. K., Ed.; CRC Press: Boca Raton, 2008, 579. (33) Johnson, K. L.; Kendall, K.; Roberts, A. D. Proc. R. Soc. London, Ser. A 1971, 324, 301. (34) Lee, J. T.; Kim, M. W.; Song, Y. S.; Kang, T. J.; Youn, J. R. Fiber. Polym. 2010, 11, 60. (35) Park, J. W.; Ediger, M. D.; Green, M. M. J. Am. Chem. Soc. 2001, 123, 49. (36) Meijer, H. E. H.; Govaert, L. E. Prog. Polym. Sci. 2005, 30, 915. (37) Riggleman, R. A.; Lee, H.; Ediger, M. D.; de Pablo, J. J. Soft. Matter. 2010, 6, 287. (38) Loo, L. S.; Cohen, R. E.; Gleason, K. K. Science 2000, 288 (5463), 116. (39) Capaldi, F. M.; Boyce, M. C.; Rutledge, G. C. Phys. Rev. Lett. 2002, 89 (17), 175505. (40) Rottler, J.; Robbins, M. O. Phys. Rev. E. 2003, 68, 011507. (41) Miller, P.; Buckley, D. J.; Kramer, E. J. J. Mater, Sci. 1991, 26, 4445. (42) Hui, C. Y.; Ruina, A.; Creton., C.; Kramer, E. J. Macromolecules 1992, 25, 3948.

mechanical properties. The PLA matrix showed the strong heterogeneity evolution during deformation. The stress prefers to be concentrated on random parts, while the polymer matrix is too rigit to prevent the stress concentration. Accordingly, the polymer matrix behaves differently under deformation; it can be classified into two regions. One is the region far away from crazes, where the local Young’s modulus remains unchanged regardless of the deformation process. The average Young’s modulus in this region were 3.5 ± 0.9 GPa, 3.3 ± 1.0 GPa, and 3.1 ± 1.0 GPa for neat, before-yielding, and postyielding PLA, respectively. The other region is the craze region where the stress concentration points located and the local molecular mobility is increased. The nanomechanical mapping on the initial craze showed the increased Young’s modulus in the edge part while the modulus was slightly lowered in the center part. With further deformation, the chain orientation and the stressinduced crystallization led to the increased Young’s modulus as high as 6.0 ± 2.5 GPa in the edge. Whereas affected by the simultaneous occurrence of the chain scission/slippage and the stress-induced orientation, the center part showed a reduced modulus of 3.0 ± 1.9, which is quite close to the macroscopic value. The microscopic deformation mechanism of PLA is thus studied phenomenologically. The results show that the brittle nature of PLA at room temperature is due to polymer’s intrinsically heterogeneous nature and the lack of chain entanglements. The stress tends to be concentrated in some parts, craze develops after the local plastic flow and finally fracture occurs at low strain. While the other parts remain unchanged during the whole deformation process. We anticipate that the submicrometer-scale nanomechanical observation provided here will lead to the better understanding on the fracture mechanism of amorphous polymer under deformation and thus help to improve the mechanical properties of PLA.

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AUTHOR INFORMATION

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by Word Premier International Research Center Initiative (WPI), MEXT, Japan.

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REFERENCES

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Quantitative Nanomechanical Investigation on Deformation of Poly

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