Investigation on Double Yielding Behavior Under Tensile Loading in Isotatctic Polypropylene

7/23/2019 Investigation on Double Yielding Behavior Under Tensile Loading in Isotatctic Polypropylene

1/11

Investigation on double yielding behavior under tensile loading

in isotactic polypropylene

Tong Wu 1, Ya Cao 1, Feng Yang , Ming Xiang 1

State Key Laboratory of Polymer Materials Engineering, Polymer Research Institute of Sichuan University, Chengdu 610065, Peoples Republic of China

a r t i c l e i n f o

Article history:

Received 18 November 2013

Accepted 18 March 2014

Available online 4 April 2014

Keywords:

Isotactic polypropylene

Double yielding

Inter-spherulitic deformation

Intra-spherulitic deformation

a b s t r a c t

In this article, a peculiar phenomenon of double yielding was first discovered in isotactic polypropylene

(iPP) under tensile loading. The results of differential scanning calorimetry (DSC), wide-angle X-ray dif-

fraction (WAXD) and polarized light microscopy (PLM) show that all the three samples, which were sub-

jected to different crystallization procedures, only forma-crystals that are composed of radial lamellaeand tangential lamellae. These a-PP samples display different double yield behaviors under tensile load-ing. PP-quenched sample exhibits double yield points when stretched at low cross-head speed (CHS),

while one single yield point appears accompanied with a marked shear band when stretched at high

CHS. However, in the case of PP-annealed, only one yield point appears at low CHS accompanied with

the formation of a large number of crazes in the necked region, meanwhile, a second yield point gradually

develops with increasing CHS. Furthermore, as for PP-isotherm, only one yield point is observed with

homogenous deformation and concomitant whitening along the whole sample at any CHS. Based on

the characterization of crystalline structure changes after yielding, we propose two plastic processes that

contribute cooperatively in the yield process ofa-PP, namely the inter-spherulitic deformation and intra-spherulitic deformation. The inter-spherulitic deformation which is prone to be initiated in the sample of

strongspherulites is predominant in the first yield process, while the intra-spherulitic deformation enters

into action after the appearance of the second yield point in the case of weak spherulites. Moreover, due

to the polydispersity of lamellae thickness, the two deformation processes are co-existed and operatecompetitively in double yielding ofa-PP.

2014 Elsevier Ltd. All rights reserved.

1. Introduction

The plastic deformation of semi-crystalline polymers has been

the subject of numerous investigations for the last forty years.

Many authors have devoted to clarify the mechanism of deforma-

tion of semi-crystalline polymers at small strains that it usually

proceeds via a yield phenomenon which is associated with a mor-

phological change of the material from a spherulitic structure into

a fibrillar one, during this process, the lamellae have been reported

to separate, tilt, untwist, and to undergo inter-lamellar slip [19].

Semi-crystalline polymers were traditionally regarded as materials

which show upon extension only one yield point on the nominal

stressstrain curve[1,35]. However, in recent years, several stud-

ies dealing with polyethylene (PE) have disclosed a singularity in

the shape of the stressstrain curves about yield point [2,1017].

This peculiar feature that consists of a hump in the vicinity of

the upper yield point, raised no comment until Mandelkern [18]

reported well-resolved double yield points for low crystallinity

ethylene copolymers and branched PE under tensile testing. He

ascribed the occurrence of double yield points to the great lamellae

thickness distribution of PE specimens. Furthermore, some authors

also reported double yielding phenomena in polyamide (PA) [19]

and poly(tetramethylene terephthalate) (PTMT) [20].

Young [5] and Argon [6] assumed two main processes of plastic-

ity associated in double yield behavior, namely a fine slip relevant

to a homogeneous shear of the crystal blocks and a coarse slip

involving fragmentation of the crystalline lamellae into blocks.

Seguela and co-workers [7,11,14] reported analogous findings form

tensile testing of PE and related copolymers and showed that the

double yield points are due to the homogeneous shear of the crys-

tal blocks (ductile process) and the slip of crystalline blocks past

each other (brittle process). Furthermore, based on a comprehen-

sive study on the structural changes during the yield process of

PE with different branch contents, Brooks and co-workers [10,13]

http://dx.doi.org/10.1016/j.matdes.2014.03.044

0261-3069/2014 Elsevier Ltd. All rights reserved.

Corresponding author. Tel.: +86 28 85403118; fax: +86 28 85402465.

E-mail addresses: [email protected] (T. Wu), [email protected]

(Y. Cao), [email protected],[email protected] (F. Yang), [email protected]

(M. Xiang).1 Tel.: +86 28 85403118; fax: +86 28 85402465.

Materials and Design 60 (2014) 153163

Contents lists available at ScienceDirect

Materials and Design

j o u r n a l h o m e p a g e : w w w . e l s e v i e r . c o m / l o c a t e / m a t d e s


2/11

constructed a deformation model that can successfully explain the

phenomenology of the double yielding with a clear picture of the

morphological changes involved. They declared that the first yield

point represents the onset of a recoverable reorientation process of

the lamellae within the spherulites in which the orientation can

almost recover upon unloading after 39 days depending on the

density of the sample. The second yield point which related to

the destruction of lamellae byc

axis shear is the onset of necking

and is the beginning of the spherulitic to fibrillar morphological

transformation.

In our study, we discovered the phenomenon of double yielding

in isotactic polypropylene (iPP) under tensile loading, which has

never been reported before. It is well established that iPP is a typ-

ical polymorphic material with several crystal modifications, of

which thea-phase is the most common crystalline form found innormal processing methods [9,2125]. Unlike other semi-crystal-

line polymers whose lamellae always grow radially, the lamellae

ofa-PP can grow in two directions, radially (R-lamellae) and tan-gentially (T-lamellae). The presence of T-lamellae improves the

strength of spherulites by acting like knots and providing anchor

spots when the spherulites deform [9,2130]. This unique cross-

hatched pattern of spherulites might provide peculiar mechanical

properties for a-PP. Lin et al. [26,27] have created microporousmembranes by stretching annealed iPP. They claimed that the

microporous structure was generated by the combination of in-

tra-spherulitic and inter-spherulitic deformations. Nevertheless,

there is no further explanation and explicit relationship between

the crystalline structure and plastic deformation ofa-PP.The uniaxial tensile testing as a function of strain rate was con-

ducted to investigate the double yielding behavior ofa-PP with dif-ferent crystalline structure. The motivation of this study is to

provide deeper understandings of relationship between spherulitic

structure and plastic deformation ofa-PP during yielding.

2. Experimental details

2.1. Materials and sample preparation

A commercially available iPP, model T38F, with a melt flow rate

(MFR) of 2.9 g/10 min (230 C, 2.16 kg),Mw= 3.8 105 g/mol and

Mw/Mn= 4.7, was purchased from Petroleum Chemical Incorpora-

tion (Lanzhou, China). 500 lm thick sheets were produced by pel-lets molding in a pressure of 5 MPa at 200 C. After melting, the

sheet was quickly put into ice water which was approximate 0 C

to obtain the sample PP-quenched. To modify the crystalline

structure, sample designated hereafter as PP-annealed was

heated from the quenched state to the temperature of 140 C in

the oven and was held for 2 h, after that, turned off the oven and

the sample slowly cooled down at about 1 C/min. Furthermore,

after melting at 200

C for 10 min, the sample called PP-isothermwas then placed between another two metal plates at 130 C and

held for 2 h before turning off the heater, which allowed the tem-

perature to gradually drop at 2 C/min.

2.2. Measurements

2.2.1. Tensile testing

Uniaxial tensile experiments were performed in accordance

with ASTM: D882-12 using an MTS Universal tensile testing

machine. Samples were cut into a mold 25 10 0.5 mm3 from

the precursor sheets and then were tested with cross-head speed

(CHS) of 1, 5, 10, 50 and 100 mm/min. All tensile measurements

were carried out at about 25 C.

2.2.2. Differential scanning calorimetry (DSC)

All the calorimetric experiments were carried out using a

Mettler Toledo DSC1 differential scanning calorimeter (DSC) under

nitrogen atmosphere (50 mL/min). Calibration for the temperature

scale was performed using indium as a standard to ensure reliabil-

ity of the data obtained. 5 mg round samples were punched out the

sheets and heated from 25 C to 190 C at a rate of 10 C/min. The

melting temperature (Tm

) of the precursor sheet was determined

from the heating curve. The crystallinity (Xdsc) of the sample was

calculated from enthalpy change values obtained in the heating

curve, and by assuming 209 J/g as the heat of fusion of a 100% crys-

talline sample.

2.2.3. Wide-angle X-ray diffraction (WAXD)

WAXD patterns were recorded with a DX-1000 diffractometer.

The wavelength of Cu Ka was k= 0.154 nm and the spectra wererecorded in the 2h range of 535, a scanning rate of 2/min, and

a scanning step of 0.02. The overall crystallinity, XXRD, was calcu-

lated according to the following equation[31,32]:

XXRD

PAcrystP

AcrystPAamorp

1

whereAcrystandAamorpare the fitted areas of crystal and amorphous

region, respectively.

2.2.4. Polarized light microscopy (PLM)

The samples were cut directly from the molded precursor

sheets and were analyzed using a Leica DMIP polarized light

microscopy, and the morphological photographs of crystallization

were recorded with the aid of a digital camera.

2.2.5. Scanning electron microscopy (SEM)

The SEMexperiments were performed using a Hitachi S3400tED

X SEM instrument to inspect the cryofractured surface of a-PPetched by a mixedacid solution [33]. The samples were gold-coated

and observed under an acceleration voltage of 20 kV.

3. Results and discussion

3.1. Characterization of crystalline structure

The WAXD spectra of the threeiPP samplessubjected to different

crystallization procedures, namely PP-quenched, PP-annealed and

PP-isotherm, are shown inFig. 1a. It can be clearly seen that all

the three samples exhibit four typical diffraction peaks of

a-crystal, namely a1 (110), a2 (040), a3 (130) and a4 (111),(04 1) and (13 1), indicating that onlya-crystals form in these pro-cessing methods [21,22,26,27,32,34,35]. On the other hand, the

melting curves of the three precursor sheets shown inFig. 1b vary

considerably: sample of PP-quenchedhas a lowTm but wide meltingpeak, whereas the PP-isotherm displays a high Tmbut narrow melt-

ing peak. In addition, the morphological characteristics of the three

PP samples listedin Table1 further reveal that the PP-quenched has

the lowest crystallinity, while the crystallinity of PP-annealed and

PP-isotherm, which are obtained from WAXD and DSC testing, in-

creaseabout 10%. It is also worth mentioning that there is no signif-

icant shift in the main melting peak of PP-annealed compared with

PP-quenched, however, a shoulder is observed in both the thermo-

grams of PP-annealed and PP-isotherm. Alamo et al. [36]and Wu

et al. [28] suggested that this low temperature discontinuity is

due to the melting of the T-lamellae. Moreover, the full width of

themelting peak at half maximum (FWHM) of thethreesamplesde-

cline in the sequence of PP-quenched (9.1 C), PP-annealed (6.8 C)

154 T. Wu et al./ Materials and Design 60 (2014) 153163


3/11

and PP-isotherm (3.5 C), indicating that the lamellar thickness dis-tributions decrease gradually.

The PLM photographs (shown inFig. 2) of the three precursor

sheets are performed to investigate the crystalline structures. It is

clear to see fromFig. 2a that the spherulites of quenched sample

are small and dense, for a large number of nuclei form initially

and the crystals grow very slowly in the rapid cooling process. After

annealing at 140 C for 2 h (shown inFig.2b), the spherulitic dimen-

sion increases slightly. Fig. 2c shows that the sample subjected to

Fig. 1. (a) WAXD and (b) DSC spectrum of three iPP samples subjected to different crystallization procedures.

Table 1The morphological characteristics of the three iPP samples subjected to different

crystallization procedures.

Samples Xxrd (%)a Xdsc(%)

b Tm (C)b FWHM(C)b

PP-quenched 50.2 41.1 163.0 9.1

PP-annealed 61.1 50.5 163.1 6.8

PP-isotherm 62.0 50.7 165.0 3.5

a Crystallinity were measured by WAXD.b Crystallinity, melting temperature and FWHMwere measured by DSC.

Fig. 2. PLM photographs of three iPP samples, 400: (a) PP-quenched; (b) PP-annealed; and (c) PP-isotherm.

T. Wu et al. / Materials and Design 60 (2014) 153163 155


4/11

isothermal crystallization has sparse but large spherulites with

clear spherulitic boundaries. Furthermore, according to Norton

and Keller [9], an R-lamellae rich a-spherulite shows a negativebirefringence sign, while ana-spherulite shows positive birefrin-gence sign when the fraction of T-lamellae is abundant. The PLM

photographs in Fig. 2 imply that the sample of PP-quenched mainly

contains R-lamellae rich spherulites, whereas the PP-annealed and

PP-isotherm have plenty of T-lamellae in their spherulites [9,23

25,36]. This finding further corroborates the DSC results.

Above all, it can be concluded that the quenched sample con-

sists of small a-spherulites with predominant R-lamellae, whileannealing improves the number of T-lamellae without increasing

the thickness of the R-lamellae dramatically. On the other hand,

the sample of PP-isotherm contains large spherulites which are full

of T-lamellae.

3.2. Uniaxial tensile testing

The nominal stressstrain curves and photographs taken under

tensile loading to gain a better understanding of the neck formation

of the three iPP samples are shown in Fig. 3. It is very interesting to

note that there are many significant differences among the three

samples in yielding behaviors. For PP-quenched sample (shown in

Fig. 3a), at high CHS (e.g. 50 mm/min), only one typical yield point

is observed and a shear band which is approximately 45 to the

draw direction appears in the necked region. This sharp necking

continues to proceed along the sample upon yielding. However, at

low CHS (e.g. 5 mm/min), the quenched sample exhibits double

yield points. It should be noted that the first yield process in which

crazes are created does not correspond to the formation of marked

neck or any other macroscopic evidence of yield, while sharp neck-

ingis observedjust after thesecondmaximum.It is also worthmen-

tioning that the quenched sample is almost transparent in the

necked region upon yielding. Nevertheless, an oppositetrend occurs

in the sample of PP-annealed (shown in Fig. 3b). At low CHS (e.g.

10 mm/min), only one yield point appears, bus as CHS increases, a

second yield point gradually develops. A particular feature of

Fig. 3b is that the sample completely whitens in the necked region

when the first yield process is largely predominant, implying that

this process leads to microvoid formation [26,27]. Furthermore,

when the second yield process enters into action, a shear band

emerges in the necked region which makes the sample becomes

translucent. On the other hand, in the case of PP-isotherm (shown

in Fig. 3c), only one yield point is observed at any CHS. At first stage

of yielding, whitening and a homogenous deformation proceed

along the whole sample and it breaks down quickly after yielding,

Fig. 3. Thenominal stressstrain curves of the three precursor sheets and corresponding deformation phenomena at different CHS: (a) PP-quenched; (b) PP-annealed; (c) PP-isotherm; and (d) photographs of PP-isotherm stretched to 15% and subsequently relaxed for 15 days.



5/11

indicating a brittle fracture dominates in this process. The peculiar

yield phenomena of iPP suggest that two plastic processes operate

competitively, namely the brittle and ductile natures of the first

yieldingprocess and the second yieldingprocess, respectively. They

are to be related to the crazing and shear banding that contribute

cooperatively in iPP[2,12,37]. This finding is of prime importance

for the understanding of the use properties of iPP, such as puncture

or stress-cracking resistance as well as creating microporous mem-

branes. In addition, photography of PP-isotherm stretched to 15%

with 1 mm/min and relaxedfor 15 days is shownin Fig.3d,it isclear

tosee that onlya smallfraction oforientationrecovers uponunload-

ing, which is different from the studies of Brooks[10,13]who hasshown an almost complete recovery of the homogeneous

deformation within 9 days in the range of first yield point of ethyl-

enehexene random copolymers. It is also noteworthy that the

yielding behaviors of iPP are sensitive to strain rates. In order for a

better understanding of the yielding processes, the nominal

stressstrain curves of PP-quenched and PP-annealed stretched ina wide range of CHS will be discussed in the following paragraph.

Fig. 4. The nominal stressstrain curves of the precursor sheets at different CHSs for: (a) PP-quenched; (b) PP-annealed; and (c) the schematic diagram of characteristic

parameters in yield (the arrow points to the onset of second yield).

Table 2

The characteristic parameters in yield for PP-quenched and PP-annealed at different

CHSs: yield stress, ry; neck width, e1; strain when the second yield happens, e2.

Samples ry (MPa) e1(%) e2 (%)

PP-quenched-1 mm/min 29.2 58.4 31.4



PP-quenched-50 mm/min 31.9 25.5

PP-quenched-100 mm/min 33.7 23.8

PP-annealed-1 mm/min 30.6 16.7



PP-annealed-50 mm/min 35.1 54.7 25.1PP-annealed-100 mm/min 38.0 58.5 27.0

Fig. 5. DSC curves of three precursor sheets and samples stretched to different

ratios with 5 mm/min.



6/11

The nominal stressstrain curvesas a function of strainrates and

corresponding characteristic parameters in yielding are shown in

Fig. 4 and Table 2, respectively. It is clear to see that these iPP sam-

ples show markedly different stressstrain curves in yielding.

Firstly, theyieldstress is in good agreement with conventional yield

point measurements that it ascends with increasing crystallinity

and strain rate [2,4,34,3842]. However, the trends of neck width

(e1) and the onset strain (e2) when second yielding happens differ

considerably. Fig. 4a shows that e2 of the quenched iPP decreases

with increasing CHS and the double yielding phenomenon disap-

pears when the CHS is higher than 10 mm/min. Together with theprevious neck profile in Fig. 3a, it canbe concluded that ductile pro-

cess (i.e. heterogeneous deformation) is preferential for quenched

iPP. On the other hand, e1declines with increasing CHS, which also

corroborates the trend of necking in the deformation regime. Hum-

bert et al.[39] havestudiedthe relationship betweenthe neckwidth

in nominal stressstrain curve and the macroscopic deformation

region of PE. They found that the neck with decreased with increas-

ing CHS, furthermore, the neck width of quenched sample was

wider than that ofannealedone, therefore, they claimed that a sharp

neck should be due to the slip of crystal blocks leading to a highly

heterogeneous deformation, while a diffuse neck is caused by a

homogeneous shearing of the crystal blocks which contains more

tie molecules. However, an opposite occurs in the annealediPP. Sin-

gle yield point accompanied with diffuse necking and whitening is

observed when the CHS is lower than 50 mm/min, indicating that

the brittle process (i.e. homogeneous deformation) is predominant

in the yield process of annealed iPP. However, it is very interesting

to notice that the e1 of PP-annealed ascends monotonously and is

even greater than that of PP-quenched when the CHS exceeds

10 mm/min, which is contradictory to Humberts theory. Thismight

be dueto the cooperative effect of the two yielding processes. In or-

der for a further investigation on this peculiar double yielding phe-

nomenon for the three iPP samples, the changes of crystallographic

and lamellar characteristics during yielding will be discussedin the

next section.

3.3. Characterization of crystalline structures after yielding

3.3.1. DSC characterization

The DSC curves and corresponding characteristic parameters of

three precursor sheets and samples stretched to different ratios

with 5 mm/min are shown inFig. 5andTable 3, respectively. For

Table 3

Thermal analysis of three precursor sheets and samples stretched to different ratios

with 5 mm/min.

Samples Xc(%) Tm (C) FWHM(C)

PP-quenched 41.1 163.0 9.1

PP-quenched-100%-1 m m/min 36.3 161.1 12.5

PP-annealed 50.5 163.1 6.8

PP-annealed-100%-1 mm/min 49.4 162.5 7.4

PP-isotherm 50.7 165.0 3.5PP-isotherm-20%-1 mm/min 50.5 165.0 3.6

Fig. 6. 2D-WAXD patterns of PP-quenched stretched to different ratios with 5 mm/min: (a) PP-quenched-0%, (b) PP-quenched-20% and (c) PP-quenched-100%; and thecorresponding 1-dimensional WAXD spectra: (d) 1D-WAXD; draw direction is horizontal.



7/11

PP-quenched, theXcand Tm decline dramatically while theFWHM

increases from 9.1 C to 12.5 C after stretched to 100%, indicating

that spherulitic deformation with concomitant lamella fragmenta-

tion may take place during the yielding process. However, in the

case of PP-annealed sample which is also stretched to 100%, the XcandTm decline slightly, meanwhile, the FWHM only increases by

0.6 C, implying that no substantial spherulitic to fibrillar morpho-

logical transformation happens in the yield process. A similar trend

is also seen in the case of PP-isotherm sample, which shows almost

the same Xc, Tm and FWHMafter stretched to 20%. Furthermore, alow temperature discontinuity is still observed in the thermogram

of PP-isotherm which suggests that the T-lamellae are still abun-

dant, indicating the existence of intact spherulites after yielding.

3.3.2. X-ray diffraction characterization

WAXD experiments have been performed in order to provide

additional information on the respective contributions of the two

plastic processes in yielding.Fig. 6 displays the 2D-WAXD patterns

and 1D-WAXD spectra of PP-quenched at different draw ratios. The

PP-quenched-0% shows three typical reflections ofa-crystal corre-

sponding to diffraction from the (110), (130) and (040) planes.When the sample stretched to 20%, these reflections appear pre-

Fig. 7. 2D-WAXD patterns of PP-annealed stretched to different ratios with 5 mm/min: (a) PP-annealed-0%, (b) PP-annealed-20% and (c) PP-annealed-100%; sample stretchedwith 100 mm/min: (d) PP-annealed-100%-100 mm/min; and the corresponding 1-dimensional WAXD spectra: (e) 1D-WAXD; draw direction is horizontal.



8/11

dominantly on the equator, indicating partial lamellae orientation

in the second yielding process. It is worth mentioning that the first

yielding process of PP-quenched can hardly be detected by X-ray

diffraction, for the crazing regime is too narrow. Moreover, the

sample shows strong reflections on the equator when stretched

to 100%, which reveals that a marked heterogeneous slip accompa-

nied with lamellae fragmentation is preferential in the yielding

process of quenched iPP[2,10].

In the cast of annealed sample, Fig. 7 shows that when stretched

to 20% with 5 mm/min, which is in the range of first yield point of

PP-annealed, lamellae orientation is hardly observed, indicating a

homogeneous deformation in the first yielding process. Further-

more, when stretched to 100%, the 2D-WAXD pattern in Fig. 7cexhibits slight reflections on the equator, together with the corre-

sponding 1D-WAXD spectra in Fig. 7e, it can be concluded that het-

erogeneous slip accompanied with lamellae fragmentation takes

place just after the occurrence of second yield point. However,

Fig. 7d shows very strong reflections of (110), (130) and (040)

planes on the equator when stretched to 100% with 100 mm/min,

indicating that heterogeneous slip accompanied with dramatic

lamellae fragmentation and orientation is predominant at high

strain rate. This trend is in accord with the photograph of necking

inFig. 3b.

Furthermore, in the cast of PP-isotherm, it is interesting to see

fromFig. 8that the isotherm sample stretched to 20% has almost

the same reflections of (110), (130) and (040) planes as the

precursor sheet, indicating that a homogeneous deformation ispredominant in this yielding process.

3.3.3. SEM analysis of stretched samples

The morphology of the stretched iPP samples was further exam-

inedby SEMin order to investigate the spherulitic structure changes

during yielding. First of all, it should be mentioned that the spheru-

litic structures of PP-quenched and PP-annealed is very difficult to

be distinguished, for their spherulites are too small to be etched

out by the mixed acid solution [33]. However, a significant morpho-

logical difference in the surface is also marked. For the quenched

sample (Fig. 9a), the surface is smooth with a dramatic oriented

fibrillar structure parallel to the draw direction, which further con-

firms that spherulitic destruction accompanied with concomitant

lamellae orientation occurs in yielding process. On the contrary,

as for PP-annealed, it is clear to notice from Fig. 9b that there aremany pores and cracks in its surface and no fibrillar structure is ob-

served when stretched to 100% with 5 mm/min, indicating that a

brittle deformation related to crazing and microvoid formation is

predominant during yielding process. It is consistent with the

results of Lin et al.[26,27]who have produced microporous mem-

branes by stretching annealed PP. Furthermore, the samples of PP-

isotherm with different draw ratio are available for us to clarify

the spherulitic deformationduring the yieldingprocess. As is shown

in Fig. 9c, the inter-spherulitic deformation occurs in the early stage

of yielding (i.e. stretched to 10%). Openings that are almost perpen-

dicular to draw axis are located among the spherulitic boundaries.

As the draw ratio reaches 20% (Fig. 9d), catastrophic inter-spheru-

litic deformation occurs accompanied with plenty of cracks

generated between spherulitic boundaries and sporadic intra-spherulitic deformation is alsoobserved. It is worth mentioning that

Fig. 8. 2D-WAXD patterns of PP-isotherm stretched to different ratios with 5 mm/min: (a) PP-isotherm-0%, (b) PP-isotherm-20%; and the corresponding 1-dimensional

WAXD spectra: (c) 1D-WAXD; draw direction is horizontal.



9/11

few elliptical spherulites are found during the whole yield process,implying that the spherulite of PP-isotherm is too rigid to be

sheared, which is different fromthe phenomenon of double yielding

in PE [2,10,11,1315,17].

3.4. A proposed mechanism of double yielding behavior in a-PP

According to the studies of PE and related copolymers, Brooks et

al. have proposed two main processes of plasticity involved in

double yielding phenomenon, namely a fine slip related to homo-

geneous deformation and a coarse slip which is a heterogeneous

deformation. The first yield point represents the onset of a orienta-

tion process of the lamellae within the spherulites that can almost

recover upon unloading within 9 days, while the second yield point

which is associated with lamellae destruction represents the onsetof necking and the beginning of the spherulitic to fibrillar morpho-

logical transformation [2,10,11,1315,17]. However, the whole set

of experimental data in our study suggests that this classical

deformation model is impractical to explain the double yielding

phenomenon ofa-PP due to its unique spherulitic structures.Unlike other semi-crystalline polymers whose lamellae always

grow radially, the lamellae ofa-PP cangrow in twodirections, radi-ally and tangentially. In the case of R-lamellae or parent lamellae,

crystallites grow along the radius of the spherulite, and the growing

direction coincides with aaxis. T-lamellae or daughter lamellae, on

the other hand, grow epitaxially on the parent lamellae. The pres-

ence of daughter lamellae improves the strength of spherulites by

acting like knots and providing anchor spots when the spherulites

deform. Furthermore, with the aid of this interlocking morphology,thea-PP spherulites are morerigid and capableof maintaining their

spherulitic morphology upon loading [9,23,24,26, 27,29,30, 36,37,43]. Thiscross-hatched pattern of R-lamellae and T-lamellae shown

in Fig. 10a is a unique pattern for a-PP andhas been documented insolution crystallization, bulk crystallization and fiber spinning

[21,22]. The DSC and PLM results of the three iPP samples disclose

that the quenched sample contains small a-spherulites with fewT-lamellae, while the annealing gradually improves the number of

T-lamellae without increasing the thickness of the R-lamellae

dramatically. Furthermore, the sample of PP-isotherm consists of

large spherulites which are full of T-lamellae.

The differences in the spherulitic structures lead to various dou-

ble yielding behaviors in the tensile testing. Combined the photo-

graphs taken under loading and characterization of crystalline

structures after yielding, we propose two plastic processes contrib-

ute cooperatively in yielding ofa-PP, namely the brittle processand the ductile process (shown in Fig. 10b), which are to be related

to the inter-spherulitic deformation and intra-spherulitic deforma-

tion, respectively.

In the case of strong spherulites, i.e. PP-isotherm, PP-annealed

and PP-quenched sample stretched with low CHS, the combined

strength of lamellae and rigid amorphous is sufficiently high to

hold up the lamellar wells and maintain the spherulitic morphol-

ogy ofa-PP. As a result, instead of the fine slip in the first yieldingprocess of PE which involves a homogeneous shear of spherulites,

inter-spherulitic deformation in iPP (i.e. brittle process shown in

Fig. 10b) is predominant in the first yielding process. The openings

at the spherulitic boundaries lead to the formation of microvoid at

the equatorial region perpendicular to the tensile axis, which

results in the appearance of diffuse crazing in the necked region.However, in the case of weak spherulites, i.e. PP-quenched and

Fig. 9. SEM images of four samples after etched: (a) PP-quenched-100%; (b) PP-annealed-100%; (c) PP-isotherm-10%; and (d) PP-isotherm-20%; the arrow represents the

draw direction.



10/11

PP-annealed sample stretched with high CHS, the lamellae are too

weak to sustain the spherulitic morphology, consequently, the

lamellae can easily break down from the lamellar knots by slip,

twinning, or strain-induced martensitic phase transformations.

The lamellae fragments align toward the stretching direction,

resulting in the formation of fibrillar morphology [26,27,29 ,37,

43]. This intra-spherulitic deformation (i.e. ductile process) which

leads to a sharp neck and dramatic shear banding is preferential in

the range of second yielding process ofa-PP and is equivalent tothe coarse slip deformation in the second yield process of PE [2,

10,11,1315,17].

It is also worth mentioning that due to the polydispersity of

lamellae thickness, the two deformation processes are co-existed

and activated selectively in the double yield behavior. Therefore,

the quenched sample which has broad lamellar thickness distribu-

tion only forms a narrow regime of crazing under low CHS, indicat-

ing that the intra-spherulitic deformation is preferential in

PP-quenched. On the other hand, annealing enhances the perfection

of crystals and reduces the lamellar thickness distribution, resulting

in a moderate regime of crazing and slight necked region which

reveals that the effects of inter-spherulitic deformation and intra-

spherulitic deformationare equal. Moreover, the cooperative effects

of the two yielding processes are more pronounced at higher CHS,leading to a gradually broadening neck width. Furthermore, the

sample of PP-isotherm which has narrow lamellar thickness distri-

bution and strong spherulites produces numerous crazes and forms

a homogeneous deformation regime along the whole sample during

yielding, implying that the inter-spherulitic deformationis predom-

inant in this yielding process. However, it is clear to see from Fig. 9c

and d that sporadic intra-spherulitic deformation is still existed in

the yield process of PP-isotherm, leading to only partial recovery

of orientation upon unloading after 15 days (shown inFig. 3d).

In summary, the double yieldingbehavior ofa-PP, whichis asso-ciated withcombination of inter-spherulitic deformation and intra-

spherulitic deformation, can be activated and controlled by the

crystalline structure as well as strain rate. This finding is important

to take full advantage ofa-PP, such as improving its mechanicalproperties and creating microporous membranes.

4. Conclusions

In this article, thedouble yieldingbehavior ofthreea-PP sampleswhich were subjected to different crystallization procedures was

investigated. Theresults of DSC, XRD andPLMof thethreeprecursor

sheets show that quenched samplehas small spherulites with main

R-lamellae, while the annealing improves the number of T-lamellae

withoutincreasing the thicknessof the R-lamellae obviously. On the

other hand, the sample of PP-isotherm has large spherulites whichcontain plenty of T-lamellae.

The nominal stressstrain curves and photographs of three iPP

samples, which are taken in the tensile loading, exhibit various

yielding behaviors. PP-quenched sample discloses double yield

points when stretched at low CHS, while one single yield point

appears accompanied with a marked shear band in the necked

region when stretched at high CHS. An opposite trend is observed

in the case of PP-annealed: only one yield point appears at low CHS

accompanied with the formation of microvoid in the necked

region, whereas a second yield point gradually develops with

increasing CHS. Furthermore, in the case of PP-isotherm, only one

yield point is observed with concomitant whitening and homoge-

nous deformation along the whole sample at any CHS.

Based on the characterization of crystalline structures afteryield through DSC, WAXD and SEM, we propose two plastic pro-

cesses that contribute cooperatively in the yielding process ofa-PP, namely the inter-spherulitic deformation and intra-spherulitic

deformation. The inter-spherulitic deformation which is initiated

in the sample with strong spherulites is predominant in the first

yielding process, while the intra-spherulitic deformation enters

into action in the second yielding process in the case of weak

spherulites. In addition, the two deformation processes are co-ex-

isted in the double yield behavior ofa-PP due to the polydispersityof lamellae thickness.

References

[1]Meinel G, Morosoff N, Pete A. Plastic deformation of polyethylene. I. Change of

morphology during drawing of polyethylene of high density. J Polym Sci PartA-2: Polym Phys 1970;8:172340.

Fig. 10. Schematic ofa-PP: (a) the lamellar morphologies ofa spherulites; and (b) the brittle and ductile processes in yielding.



11/11

[2] Gaucher-Miri V, Seguela R. Tensile yield of polyethylene and related

copolymers: mechanical and structural evidences of two thermally activated

processes. Macromolecules 1997;30:115867.

[3]Matsuo M, Hirota K, Fujita K, Kawai H. Studies on the deformation mechanism

of polyethylene spherulites by the orientation distribution function of

crystallites. Macromolecules 1978;11:10007.

[4]Peterlin A. Molecular mechanism of plastic deformation of polyethylene. J

Polym Sci C: Polym Symp 1967;18:12332.

[5]Bowden PB, Young RJ. Deformation mechanisms in crystalline polymers. J

Mater Sci 1974;9:203451.

[6]Lin L, Argon AS. Structure and plastic deformation of polyethylene. J Mater Sci1994;29:294323.

[7] Darras O, Seguela R. Tensile yield of polyethylene in relation to crystal

thickness. J Polym Sci Part B: Polym Phys 1993;31:75966.

[8] Lee BJ,Argon AS, Parks DM, AhziS, BartczakZ. Simulation of large strainplastic

deformation and texture evolution in high density polyethylene. Polymer

1993;34:355575.

[9] Norton DR, Keller A. The spherulitic and lamellar morphology of

meltcrystallized isotactic polypropylene. Polymer 1985;26:70416.

[10] Brooks NWJ, Unwin AP, Duckett RA, Ward IM. Double yield points in

polyethylene: structural changes under tensile deformation. J Macromol Sci

Part B 1995;34:2954.

[11] Seguela R, Rietsch F. Double yield pointin polyethylene undertensile loading. J

Mater Sci Lett 1990;9:467.

[12] Schrauwen Bernard AG, Janssen Roel PM, Govaert Leon E, Meijer HEH. Intrinsic

deformation behavior of semicrystalline polymers. Macromolecules

2004;37:606978.

[13] Brooks NWJ, Duckett RA, Ward IM. Investigation into double yield points in

polyethylene. Polymer 1992;33:187280.

[14] Seguela R, Darras O. Phenomenological aspects of the double yield of

polyethylene and related copolymers under tensile loading. J Mater Sci

1994;29:534252.

[15] Lucas JC, Failla MD, Smith FL, Wdelkefw L. The double yield in the tensile

deformation of the polyethylenes. Polym Eng Sci 1995:35.

[16] Balsamo V, Muller AJ. The phenomenon of double yielding under tension in

low-density, linear low-density polyethylene and their blends. J Mater Sci Lett

1993;12:14579.

[17] Feijoo JL, Sanchez JJ, Mueller AJ. The phenomenon of double yielding in

oriented high density polyethylene films. J Mater Sci Lett 1997;16:17214.

[18] Popli R, Mandelkern L. Influence of structural and morphological factors on the

mechanical properties of the polyethylenes. J Polym Sci Part B: Polym Phys

1987;25:44183.

[19] Shan GF, Yang W, Xie B, Li Z, Chen J, Yang M. Double yielding behaviors of

polyamide 6 and glass bead filled polyamide 6 composites. Polym Testing

2005;24:70411.

[20] Muramatsu S, Lando JB. Double yield points in poly(tetramethylene

terephthalate) and its copolymers under tensile loading. Polym Eng Sci

1995;35:107785.[21] Bruckner S, Meille SV. Polypropylene. Dordrecht (The Netherlands): Kluwer

Academic Publishers; 1999.

[22] Phillips RA, Wolkowicz MD. Polypropylene handbook. OH (USA): Hanser

Gardner Publications Inc.; 2005.

[23] Varga J. Supermolecular structure of isotactic polypropylene. J Mater Sci

1992;27:255779.

[24] Lotz B, Wittmann JC, Lovinger AJ. Structure and morphology of

poly(propylene): a molecular analysis. Polymer 1996;37:497992.

[25] Poussin L, Bertin YA, Parisot JCB. Influence of thermal treatment on the

structure of an isotactic polypropylene. Polymer 1998;39:42615.

[26] Lin KY, Xanthos M, Sirkar KK. Novel polypropylene microporous membranes

via spherulitic deformation processing perspectives. Polymer

2009;50:467182.

[27] Lin KY, Xanthos M, Sirkar KK. Novel polypropylene-based microporous

membranes via spherulitic deformation. J Membr Sci 2009;330:26778.

[28] Wu CM, Chen M, Karger-Kocsis J. Effect of micromorphological features on the

interfacial strength of iPP/Kevlar fiber microcomposites. Polymer2001;42:199208.

[29] Samios D, Tokumoto S, Denardin ELG. Large plastic deformation of isotactic

poly(propylene) (iPP) evaluated by WAXD techniques. Macromol Symp

2005;229:17987.

[30] Androsch R, Wunderlich B. The link between rigid amorphous fraction and

crystal perfection in cold-crystallized poly(ethylene terephthalate). Polymer

2005;46:1255666.

[31] Huo H, Jiang SC, An LJ. Influence of shear on crystallization behavior of the b

phase in isotactic polypropylene with b-nucleating agent. Macromolecules

2004;37:247883.

[32] Luo F, Geng C, Wang K, Deng H, Chen F, Fu Q, et al. New understanding in

tuning toughness of b-polypropylene: the role of b-nucleated crystalline

morphology. Macromolecules 2009;42:932531.

[33] Olley RH, Bassett DC. An improved permanganic etchant for polyolefines.

Polymer 1982;23:170710.

[34] Mimaroglu A, Yenihayat OF, Celebi A. The influence of thermal history, strain

rate and sample geometry on the deformation behaviour of polymers: use of

the thermovision technique. Mater Des 1995;16:199203.

[35] Xu T, Lei H, Xie CS. The effect of nucleating agent on the crystalline

morphology of polypropylene (PP). Mater Des 2003;24:22730.

[36] Alamo RG, Brown GM, Mandelkern L, Lehtinen A, Paukkeri R. A morphological

study of a highly structurally regular isotactic poly(polypylene) fraction.

Polymer 1999;40:393344.

[37] Phillips A, Zhu PW, Edward G. Simple shear deformation of polypropylene via

the equal channel angular extrusion process. Macromolecules

2006;39:5796803.

[38] Seguela R. On the natural draw ratio of semi-crystalline polymers: review of

the mechanical, physical and molecular aspects. Macromol Mater Eng

2007;292:23544.

[39] Humbert S, Lame O, Vigier G. Polyethylene yielding behaviour: what is behind

the correlation between yield stress and crystallinity? Polymer

2009;50:375561.

[40] Mohammadpour E, Awang M, Kakooei S, Akil H. Modeling the tensile stress

strain response of carbon nanotube/polypropylene nanocomposites using

nonlinear representative volume element. Mater Des 2014;58:3642.

[41] Zebarjad SM, Bagheri R, Lazzeri A, Serajzadeh S. Fracture behaviour of isotactic

polypropylene under static loading condition. Mater Des 2003;24:1059.[42] Zebarjad SM, Bagheri R, Lazzeri A, Serajzadeh S. Dilatational shear bands in

rubber-modified isotactic polypropylene. Mater Des 2004;25:24750.

[43] Coulon G, Castelein G, GSell C. Scanning force microscopic investigation of

plasticity and damage mechanisms in polypropylene spherulites under simple

shear. Polymer 1998;40:95110.


Documents

Investigation on Double Yielding Behavior Under Tensile Loading in Isotatctic Polypropylene