Flow induced crystallization of polymers:Precursors and Nuclei

Flow induced crystallization of polymers: Precursors and Nuclei / by Zhe Ma.Technische Universiteit Eindhoven, 2012.

A catalogue record is available from the Eindhoven University of Technology Library.ISBN: 978-90-386-3299-5

This thesis was prepared with the LATEX2 documentation system.Reproduction: University Press Facilities, Eindhoven, The Netherlands.Cover design: Zhe Ma and Paul Verspaget

This research forms part of the research programme of the Dutch Polymer Institute(DPI), DPI project #714.

Flow induced crystallization of polymers:Precursors and Nuclei

PROEFSCHRIFT

ter verkrijging van de graad van doctor aan deTechnische Universiteit Eindhoven, op gezag van derector magnificus, prof.dr.ir. C.J. van Duijn, voor een

commissie aangewezen door het College voorPromoties in het openbaar te verdedigen

op maandag 3 december 2012 om 16.00 uur

door

Zhe Ma

geboren te Hejian, Hebei, China

Dit proefschrift is goedgekeurd door de promotoren:

prof.dr.ir. G.W.M. Peters

en

prof.dr.ir. H.E.H. Meijer

Contents

Summary ix

1 Introduction 11.1 Background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11.2 Structure-properties relation . . . . . . . . . . . . . . . . . . . . . . . . . 21.3 Processing-structures relation . . . . . . . . . . . . . . . . . . . . . . . . 41.4 Scope of the thesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5

References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

2 Using rheometry to determine nucleation density in a colored systemcontaining a nucleating agent 92.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 102.2 Experimental . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

2.2.1 Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 112.2.2 Rheological characterization . . . . . . . . . . . . . . . . . . . . 112.2.3 X-ray characterization . . . . . . . . . . . . . . . . . . . . . . . 12

2.3 Data analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 122.4 Results and Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . 14

2.4.1 Determination of the nucleation density for NA-iPP . . . . . . . 142.4.2 Reproducibility . . . . . . . . . . . . . . . . . . . . . . . . . . . 162.4.3 Effect of mild flow . . . . . . . . . . . . . . . . . . . . . . . . . 17

2.5 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

3 Pressure quench of flow-induced crystallization precursors 233.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 243.2 Experimental . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

3.2.1 Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 263.2.2 Protocol . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 273.2.3 X-ray characterization . . . . . . . . . . . . . . . . . . . . . . . 29

3.3 Results and Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . 303.3.1 Reference experiment: no pressure quench after flow . . . . . . . 303.3.2 Pressure quench after flow . . . . . . . . . . . . . . . . . . . . . 31

v

vi Contents

3.3.3 Pressure quench after annealing . . . . . . . . . . . . . . . . . . 343.3.4 Inverse quench by depressurization . . . . . . . . . . . . . . . . 38

3.4 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39

Appendices. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42

4 Short-term flow induced crystallization in isotactic polypropylene:how short is short? 454.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 464.2 Experimental . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48

4.2.1 Material . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 484.2.2 Methods . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 484.2.3 Optical microscopy . . . . . . . . . . . . . . . . . . . . . . . . . 50

4.3 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 504.3.1 Rheological evolution during flow . . . . . . . . . . . . . . . . . 504.3.2 Structural evolution . . . . . . . . . . . . . . . . . . . . . . . . 52

4.4 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 564.4.1 SAXS delay time . . . . . . . . . . . . . . . . . . . . . . . . . . 574.4.2 Implications of the viscosity rise . . . . . . . . . . . . . . . . . . 594.4.3 Conditions of the viscosity rise . . . . . . . . . . . . . . . . . . . 60

4.5 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62

5 The influence of flow induced precursors and nuclei on crystallizationof isotactic polypropylene 655.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 665.2 Experimental . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 68

5.2.1 Material . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 685.2.2 Flow device . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 685.2.3 X-ray characterization . . . . . . . . . . . . . . . . . . . . . . . 685.2.4 Birefringence . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72

5.3 Results and Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . 735.3.1 Flow-induced distinguishable nuclei/precursors . . . . . . . . . . 735.3.2 Isothermal crystallization . . . . . . . . . . . . . . . . . . . . . . 75

5.4 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 87References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 87

6 High-stress shear induced crystallization in isotactic polypropyleneand propylene/ethylene random copolymers 896.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 906.2 Experimental . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 91

6.2.1 Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 916.2.2 Flow device . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 926.2.3 X-ray characterization . . . . . . . . . . . . . . . . . . . . . . . 926.2.4 Optical microscopy . . . . . . . . . . . . . . . . . . . . . . . . . 93

Contents vii

6.3 Depth sectioning method . . . . . . . . . . . . . . . . . . . . . . . . . . . 936.4 Results and Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . 96

6.4.1 iPP homopolymer . . . . . . . . . . . . . . . . . . . . . . . . . . 966.4.2 Propylene/ethylene random copolymers . . . . . . . . . . . . . . 996.4.3 Quantification of nuclei . . . . . . . . . . . . . . . . . . . . . . . 102

6.5 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 106References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 106

7 Conclusions and recommendations 1097.1 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1097.2 Recommendations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111

References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 113

Samenvatting 115

Acknowledgements 119

Curriculum vitae 121

List of publications 123

Summary

Semi-crystalline polymers cover over two thirds of commercial applications ofpolymeric materials. With most industrial processing technologies, semi-crystallinepolymers are brought in the molten state and then deformed for different reasons:i.e. to create specific properties like a high modulus with fiber spinning or to shapegeometrically complex parts as with injection molding. The deformation not onlyaccelerates the crystallization kinetics but also can cause crystalline structures to changefrom isotropic spherulites to the highly oriented cylindrical structures (e.g. shish-kebab) which determine the ultimate properties. Therefore, understanding the interplaybetween flow fields and the resulting crystalline structures is of great importance. Itenables the design of processing procedures and tailoring final product properties. Theobjective of this thesis is to reveal how a flow field affects the crystal structures bystudying the early stages of the crystallization process which are dominated by thecreation of shear-induced precursors and nuclei.

Polymer crystallization involves two steps: nucleation and growth. Nucleationprovides quantitative (more or less) and qualitative (isotropic or oriented) nuclei andthese are the templates for further crystal growth that, ultimately, will fill the full space.The nucleation step is very sensitive and is often controlled by additives (nucleatingagents) or/and imposing a flow field. The formation and features of nuclei are, therefore,the key factors that determine the crystalline structures. Probing and quantifying nucleiis the main work of this thesis.

According to the resulting morphology, nuclei can be divided into two groups: point-like nuclei and fibrillar nuclei. The former give rise to spherulites, the latter mainlyinduce oriented structures like the well-known shish-kebab; i.e. fibrils with transverselamellae. Oriented nuclei can be further classified into row nuclei and shish nuclei bywhether they can be directly observed with X-ray characterizations. In the first part ofthis thesis, the creation of point-like nuclei is studied for an isotactic polypropylene witha nucleating agent, with and without applied shear deformation. For such a system,nucleation is dramatically enhanced by both the nucleating agent (U-Phthalocyanine)and the flow. The nucleation density is that large and, therefore the crystallite sizeso small that optical microscope is not suitable to count nuclei numbers. Therefore, asuspension-based rheological model is used to quantify the nuclei density.

ix

x Summary

In the second part, precursors of row nuclei in a bimodal Polyethylene (PE) arestudied. Since row nuclei cannot be observed directly by X-ray characterizations, thecrystallization step is triggered by increasing the pressure which raises the equilibriummelting temperature and thus the under-cooling: a so-called pressure quench. It isshown that shear-induced precursors can be generated at a temperature close to thenominal melting temperature and the unstable ones can relax.

Next, the formation of shish nuclei during flow is studied by using fast (30 frame/s),time resolved X-ray scattering in combination with rheological measurements. It isfound that a critical shear rate exists for the formation of shish nuclei within very shortflow times (0.25 s). Shish precursors are formed during flow and these were found todevelop into shish afterwards, during or after flow depending on the flow strength.

Besides the external effects, the nuclei formation is also influenced by the molecularstructure. Therefore, in the last part, the effect of the molecular architecture onshear-induced crystallization is studied for isotactic polypropylene (iPP) and twopropylene/ethylene random copolymers with varying ethylene monomer contents. Thesethree grades had very similar rheological behavior; there is only a small but importantdifference in the longest relaxation time. Flow enhanced nucleation density was foundto be the lowest for the iPP homopolymer which is due to the shorter longest relaxationtime. However the reduced growth rate due to the added ethylene monomer leadsto slower overall crystallization kinetics of the random copolymers compared to thehomopolymer at similar thermal conditions.

Chapter one

Introduction

1.1 Background

Human history is closely accompanied by the progress in material development.

Advances in material are crucial for affecting peoples way of life because it can

significantly change the way of working and improve output efficiency. The consequence

of the material evolution is particularly obvious at the primary stages of civilization.

For example, the innovation of iron farming tools such as plows and sickles, made the

farming processes much faster and easier and, as a consequence, people produced enough

food to survive and also gained extra time to exploit other activities. Considering such

great importance, anthropologists even define historical epochs by the materials used

at that moment, such as the Stone, Bronze and Iron ages. To meet the continuously

increasing need of new functions, synthetic materials appeared and have been widely

applied. Among current synthetic creations, plastic is quite prominent in our society

due to its wide range of applicability. Just looking around, plastic products can be

easily found at any time and any place. They may be a coffee stirrer, mobile shell,

computer screen, bumper, greenhouse, construction and building parts, and so on.

Actually, the development of truly synthetic plastics started only around one hundred

years ago. The first completely synthetic one is poly(phenol-co-formaldehyde), a densely

cross-linked thermoset polymeric material, well known as Bakelite, commercialized in

1909 and applied for electrical appliances and phonograph records [1]. After its birth,

plastic has developed dramatically during the past century. On one hand, more and

more new plastics were synthesized and manufactured. Some major examples are:

poly(vinyl chloride) (PVC) in 1920s, polyamide (Nylon) 66 in 1930s, linear high-density

polyethylene (HDPE) and isotactic polypropylene (iPP) in 1950s, poly(p-phenylene

terephthalamide) (known as Kevlar) in 1960s. [1, 2]

1

2 Chapter 1

1950 1960 1970 1980 1990 2000 2010

0

50

100

150

200

250

300

2010: 265

2009: 250

2002: 204

1989: 99

1976: 47

Mtonne

year

1950: 1.5

Figure 1.1: World plastics production between 1950 and 2010. Data are obtained fromreference [3]. Only continuous growth is shown and temporary fluctuation is notincluded.

On the other hand, the total amount of the plastic being used also increased

tremendously, as shown in Figure 1.1. The world plastic production has risen from

1.5 million tonnes in 1950 to 265 million tonnes in 2010, [3] equal to an averaged annual

growth rate of 9% during sixty years.

Polymers can be amorphous or semi-crystalline, depending on the degree of ordering

of the molecular arrangement. Due to crystallization, semi-crystalline polymers can

still be used as a solid material above the glass transition temperature. For example,

although the glass transition temperature of HDPE is -80C, [1] its high crystallinity

enables HDPE to be used as a solid-material at normal environmental temperatures.

Therefore, semi-crystalline polymers, especially polyolefins such as HDPE and iPP, have

become the most widely used polymeric materials, see Figure 1.2.

1.2 Structure-properties relation

Ultimate properties, e.g. mechanical and optical, of end-use plastic products strongly

depend on the structure of the material [4, 5]. The manufacturing history includes two

major processes: a) synthesis, starting with small molecules and b) processing, where

the final product is given its shape. Therefore, the general issue structure includes

two classes of defining features: the intrinsic chemical structures of the macromolecules

and the physical arrangement structure of these macromolecules.

The intrinsic chemical structure includes chemical composition (various monomers,

homo-polymer/co-polymer, etc.), chain structure (linear/branched/star/crossed, ran-

dom/block, etc.) and configuration, molecular weight, and so on. These molecular

Chapter 1 3

Figure 1.2: European plastics demand by polymer type in 2010. [3]

properties are controlled and dominated by the polymerization process.

For amorphous polymers, considerable knowledge has been developed to understand

the relationship between the basic molecular structure and mechanical behavior. Semi-

crystalline polymers contain both amorphous and crystalline phases where an identical

polymer chain often is involved in these different phases. The crystal phase has the

hierarchic structures of conformation, crystal lattice, lamellar crystal, spherulitical or

cylindrical assembling form of lamellae, etc. These crystal micro-structures cover a

broad range of length scales from sub-nanometer up to microns. Polymer crystallization

is determined by the molecular structures and, often in a very strong way, on the

processing conditions. Therefore, when discussing structure-properties relation of semi-

crystalline polymers, specific features of the crystal microstructure like crystallinity,

lamellar thickness and orientation, should be taken into account. A typical and

illustrative example [6] is shown in Figure 1.3, where an injection molded plate of HDPE

is shown. From the optical micrographs it is observed that different structures appear

in the thickness direction. These skin and core layers are due to the varying cooling

and flow conditions as experienced by the material at different positions. Moreover,

the thickness of these layers strongly depends on the position along flow path. Finally,

for the same polymer material of HDPE, these distinct structures lead to completely

different mechanical properties between specimens cut from the same plate but at

different locations and in different directions. The failure mode varies from brittle

to tough, the latter showing necking in one case and homogeneous deformation in the

other.

Next to molecular properties, full understanding on how crystal structures are formed

during processing and how these subsequently affect the properties is a prerequisite of

tailoring the ultimate properties of plastic products. This thesis focuses on the second

step in the above knowledge chain, i.e. how polymer crystallizes under processing-

relevant conditions.

4 Chapter 1

70x70x1mm

injectionofpolymer

A B

C

A

B

C

Figure 1.3: Variation in microstructure over the thickness in an injection molded HDPEplate. Right column shows the resulting mechanical properties of threespecimens cut from the same plate but different positions. [6]

1.3 Processing-structures relation

Semi-crystalline polymers are often processed from the molten states and subjected

to flow to transport the material and shape the product. Flow gradients affect

crystallization by altering the amount, size and orientation of nuclei. In this way,

processing-relevant flows are able to accelerate the crystallization kinetics by several

orders of magnitude and induce anisotropic structures. According to the resulting

morphology, flow-induced nuclei can roughly be divided into two groups: point-like

nuclei and oriented fibrillar nuclei. During quiescent crystallization, only point-like

nuclei are formed. The average crystal structure is isotropic and may grow into large

spherulites because of the relatively low nuclei density, see Figure 1.4a. Depending on

the competition between flow strength and molecular mobility, there are mild and strong

flows. Mild flow forms extra point-like nuclei, i.e. increase nuclei density which results

in more but smaller spherulites, see Figure 1.4b. In contrast, strong flow completely

changes the isotropic morphology into oriented crystals. As an example highly oriented

shish-kebab structures are shown in Figure 1.4c. [7]

If the flow should effectively alter the crystallization behavior, the nuclei that are

generated should be stable. However, in many cases flow-induced ordered structures

are so-called precursors, instead of stable nuclei. These flow-induced precursors have

Chapter 1 5

(a) (b) (c)

Figure 1.4: Typical morphologies of: a) less and large spherulites; b) more but smallspherulites; c) oriented shish-kebab [7].

a degree of ordering (position or/and orientation; intra- or/and inter- molecule) and

favor further development to nuclei for sufficient under-cooling, or may relax back to

amorphous melt in case of relatively high temperatures. When the unstable precursors

are inclined to disappear, the slow kinetics of precursor relaxation may cause significant

memory effect on subsequent crystallization, since cooling stabilizes and reactivates the

residual precursors. [8]

By stretching polymer chains and aligning segments parallel to each other, flow

promotes nucleation in polymer. On the other hand, foreign substances, known as

nucleating agent (NA), are added to the polymer melt and the additional surfaces

lower the surface energy change in an efficient way and decrease the critical size for

nucleation [9]. Moreover, specific nucleating agents can control the phase modification of

growing crystals. For example, -NA is utilized to induce formation of iPP hexagonal -

crystals [1012]. Understanding how the amount and orientation of nuclei are controlled

by flow or/and NA is a key step in revealing processing-structure relations.

1.4 Scope of the thesis

The aim of this thesis is to gain understanding on how polymer crystallization under

processing-relevant conditions (flow and NA) takes place by exploring the creation and

evolution of precursors/nuclei. During processing, a continuous flow imposed during

solidification, may influence both nucleation and crystal growth. To separate the

influence of flow on these two aspects from rheological changes due to crystallization,

Janeschitz-Kriegl and co-workers introduced the short-term shearing method [13].

Flow duration is so short that structure formation and a related change in the viscosity

can be minimized. The flow-induced precursors/nuclei can be revealed by direct

observations, using state of the art experimental techniques, or by examining subsequent

crystallization. In this thesis, we use both approaches to study (probe or/and quantify)

precursors/nuclei induced by various flow strength. Moreover, the validity of short-

6 Chapter 1

term flow is verified for strong shear conditions. The polymers investigated are grades

of a high density polyethylene (HDPE) and an isotactic polypropylene (iPP). They are

not only the most used polymers but also the simplest macromolecules and provide

suitable model systems for the academic research.

First of all, generation of extra point-like nuclei is studied in Chapter 2. Addition of

nucleating agents (U-Phthalocyanine) and imposition of mild flow are both utilized to

enhance formation of point-like nuclei in an iPP system. A suspension-based rheological

model [14,15] is applied to quantify nucleation density. In this way, the nucleation effects

of this nucleating agent and mild flows applied are assessed.

Oriented nuclei are investigated in Chapter 3-6. Oriented nuclei can be classified

as row nuclei and shish, according to whether such nuclei can be detected by X-ray

characterization.

Chapter 3 focuses on the precursors of the X-ray invisible row nuclei. Shear-induced

precursors cannot be observed directly by X-ray. Therefore, after increasing the under-

cooling, subsequent crystallization kinetics and orientation evolution is examined to

reflect the underlying precursors. A pressure quench method is presented, which

provides an efficient way of achieving sufficient under-cooling in a fast way for triggering

crystallization and effectively lightens up precursors.

In the case of very strong shear, X-ray observable shish appear during flow. In

Chapter 4, the combination of a slit flow with time resolved X-ray measurements

provides the possibility to study in-situ structure formation of iPP during flow. In this

way, shish formations during and after flow can be distinguished. Also the evolution

of the average rheological behavior is tracked. The appearance of shish structure and

the rise of the apparent viscosity identify the validity of the conditions for short-term

flow, i.e. how short is short depends on shear strength.

For the weakest flow conditions studied in Chapter 4, X-ray and rheology fail to detect

shear-induced structures during flow. Therefore, in Chapter 5, birefringence is employed

to probe the potential precursors in these flows. Subsequent crystallization features

such as kinetics, orientation and -crystals are tracked in time, in order to establish a

correlation between the precursors/nuclei and their influences on crystallization.

The effect of flow strength in relation with molecular structure (defects present in the

molecular chain and high molecular tail) on crystallization is explored in Chapter 6, by

comparing the amount of oriented nuclei in an iPP and two ethylene/propylene random

copolymers with various ethylene content. A pressure-driven slit flow device [16, 17]

is employed to impose strong flows with a wall stress of up to 0.11 MPa. With this

novel device, the depth sectioning method [18] can be applied to extract the polymer

crystallization in certain layers and correlate that with the specific stress. In this way,

the effect of molecular structures on flow induced nucleation can be revealed.

Chapter 1 7

References

[1] L.H. Sperling. Introduction ot Physical Polymer Science. Wiley, fourth edition edition, 2006.

[2] D. I. Bower. An Introduction to Polymer Physics. Cambridge University Press, 2002.

[3] Plastics Europe. Plastics - the facts 2011. an analysis of european plastics production, demand

and recovery for 2010. Technical report, 2011. Http://www.plasticseurope.org.

[4] B. A. G. Schrauwen, R. P. M. Janssen, L. E. Govaert, and H. E. H. Meijer. Macromolecules

37(16):60696078, 2004.

[5] B. A. G. Schrauwen, L. C. A. Von Breemen, A. B. Spoelstra, L. E. Govaert, G. W. M. Peters,

and H. E. H. Meijer. Macromolecules 37(23):86188633, 2004.

[6] B. A. G. Schrauwen. Deformation and Failure of Semicrystalline Polymer Systems: Influence of

Micro and Molecular Structure. Ph.D. thesis, Eindhoven University of Technology, 2003.

[7] B. S. Hsiao, L. Yang, R. H. Somani, C. Avila-Orta, and L. Zhu. Physical Review Letters

94(11):117802, 2005.

[8] D. Cavallo, F. Azzurri, L. Balzano, S. S. Funari, and G. C. Alfonso. Macromolecules 43(22):9394

9400, 2010.

[9] B. Wunderlich. Macromolecular Physics, Vol 2: Crystal Nucleation, Gorwth, Annealing.

Academic Press, New York, 1976.

[10] J. Garbarczyk and D. Paukszta. Polymer 22(4):562564, 1981.

[11] W. Stocker, M. Schumacher, S. Graff, A. Thierry, J. C. Wittmann, and B. Lotz. Macromolecules

31(3):807814, 1998.

[12] F. Luo, C. Geng, K. Wang, H. Deng, F. Chen, Q. Fu, and B. Na. Macromolecules 42(23):9325

9331, 2009.

[13] S Liedauer, G Eder, H Janeschitz-Kriegl, P Jerschow, W Geymayer, and E. Ingolic. International

Polymer Processing 8:236244, 1993.

[14] R. J. A. Steenbakkers and G. W. M. Peters. Rheologica Acta 47(5-6):643665, 2008.

[15] J. W. Housmans, R. J. A. Steenbakkers, P. C. Roozemond, G. W. M. Peters, and H. E. H. Meijer.

Macromolecules 42(15):57285740, 2009.

[16] G. Kumaraswamy, R. K. Verma, and J. A. Kornfield. Review of Scientific Instruments 70(4):2097

2104, 1999.

[17] G. Kumaraswamy, A. M. Issaian, and J. A. Kornfield. Macromolecules 32(22):75377547, 1999.

[18] L. Fernandez-Ballester, D. W. Thurman, and J. A. Kornfield. Journal of Rheology 53(5):1229

1254, 2009.

Chapter two

Using rheometry to determine

nucleation density in a colored

system containing a nucleating

agent

Abstract

A new suspension-based rheological method was applied to experimentally study the

crystallization of a nucleating agent (NA) filled isotactic polypropylene. This method

allows for determination of point nucleation densities where other methods fail. For

example, optical microscopy can fail because nucleation densities become too high to

be counted (materials with effective NA) or crystallites are not easily visible (colored

materials), while differential scanning calorimetry does not allow the effect of flow to be

studied. Both quiescent and mild-shear induced crystallization were investigated. The

results show that the addition of a nucleating agent increases the nucleation density by

six decades for quiescent crystallization. The effect of shear on crystallization in the

presence of a nucleating agent was assessed and it is demonstrated that, at least for

this system, the effect of applied shear is much smaller than the effect of the nucleating

agent.

This chapter is based on : Zhe Ma, Rudi J.A. Steenbakkers, Julien Giboz, Gerrit W.M. Peters.Rheologica Acta 50:909915, 2011

9

10 Chapter 2

2.1 Introduction

Semi-crystalline polymers cover over two thirds of the products used in our daily life.

Adding a nucleating agent to a semi-crystalline polymer is a common way to control

crystallization and tailor mechanical and optical properties [14]. The interface between

the nucleating agent and the polymer melt has a high surface energy, which reduces

the energy barrier associated with the formation of a nucleus. Introducing a large

amount of nucleating agent particles therefore greatly increases the nucleation density

and consequently accelerates crystallization [5]. In general, a higher nucleation density

leads to more desirable properties. Several methods are available to determine the

nucleation density in polymer crystallization. These include (a) microscopy to directly

count the number of nuclei in the melt, (b) differential scanning calorimetry (DSC) [6],

or (c) dilatometry [7]. The latter two indirectly measure the crystallization evolution in

order to calculate the nucleation density using a kinetics equation. However, for colored

nucleating agent systems, the nucleation density can become too large to be counted

and also can change the optical properties dramatically [4], making quantification of

the nucleation density difficult with optical microscopy (Figure 2.1).

Figure 2.1: Morphology during quiescent crystallization at T = 151 C.

Flow is another crucial factor affecting polymer crystallization, especially the

nucleation step. The number of nuclei can increase dramatically [8, 9] and for a strong

enough flow the formation of shish-kebab structures will occur [10]. The combined

effect of nucleating agent and flow has hardly been studied. In this work we will restrict

ourselves to point-like nucleation which will occur for moderate flow only. How much

a flow will change the kinetics of nucleating agent crystallization depends on the type

of nucleating agent [11]. Knowledge of the individual effect of nucleating agent and

shear on nucleation density is a prerequisite for understanding the combined nucleation

mechanism and to provide input for models to predict resulting structures.

Chapter 2 11

To achieve this goal, conventional DSC and classical dilatometry are not suited since

they do not allow one to impose a flow on the material. A new type of dilatometer,

the Pirouette [12], developed in our group [13, 14], does allow for applying shear to

the sample and thus for using the measured time dependent specific volume for our

purpose; determining the nucleation density. However, in this chapter we will use a

rheometer that first serves as flow device (with the possibility to vary the shear rate

and the shear time) and, subsequently, as a mechanical spectrometer that measures

the changing complex modulus due to the progressing crystallization process. Using an

analytical relation between space filling and the modulus, it is possible to determine

the nucleation density. This solves both problems: dealing with large numbers of nuclei

and poor visibility of crystalline structures. The aim of this study is to explain this

approach and to demonstrate its effectiveness for colored nucleating agent systems, for

both quiescent and flow-induced crystallization.

2.2 Experimental

2.2.1 Materials

The polymer used in this study is an isotactic polypropylene (iPP; HD601CF,

Borealis, previously known as HD120MO). It has a weight average molecular weight

Mw = 365 kg/mol and a polydispersity Mw/Mn = 5.4 [9]. The nominal melting

temperature is 163 C. The polymer was compounded with an organic nucleating

agent, U-Phthalocyanine of molecular weight 310 kg/mol [15] at a concentration of

0.2wt%. This nucleating agent was also used by Lee Wo and Tanner [16] who found

non-spherical crystallites that we did not observe, see section 2.4. From the neat iPP

and the artificially nucleated material (NA-iPP), 1.1-mm-thick plates were injection

molded and from those plates circular disks were cut with a diameter of 8mm. The

NA-iPP samples were blue due to the coloring effect of the nucleating agent.

2.2.2 Rheological characterization

For the rheological measurements, a Rheometrics ARES rheometer with a plate-plate

geometry was used. Samples were first heated to 230 C and kept on that temperature

for 10 min to erase the thermal and mechanical history. Next, the melt was cooled to

the desired temperature at a rate of 15 C/min and kept at this temperature for the

isothermal crystallization. During cooling, gap adjustment was performed continuously

to compensate for the thermal shrinkage of tools and sample. Two minutes of delay

time was used to equilibrate the sample temperature before starting the dynamic

measurements. Temperatures for crystallization were chosen between 133 and 140 C

12 Chapter 2

for the neat iPP and between 143 and 151 C for the NA-iPP. Experiments with shear

were carried out at temperatures of 148 and 151 C. Steady shear at a rate of = 60 s1

was imposed for shear times ts = 2, 4 and 6 s. Small-amplitude oscillatory shear

measurements were employed to track the time evolution of the storage modulus (G)

and loss modulus (G) using an angular frequency of 5 rad/s and a strain of 0.5%. All

experiments were performed in a N2 atmosphere to prevent the material from degrading.

2.2.3 X-ray characterization

Polypropylene can crystallize in different phases. To check if both the neat iPP and

the NA-iPP had crystallized in the -phase, the samples were analyzed afterwards using

X-ray scattering. Wide angle X-ray diffraction (WAXD) measurements were carried

out at the Dutch-Belgian (DUBBLE) beamline BM26B of the European Synchrotron

Radiation Facility in Grenoble, France. [17] A Photonica CCD detector with 20041335pixels of 44m 44m was placed at 178mm. The wavelength used was 1.033 A andthe exposure time was 10 s. WAXD data were integrated with the software FIT2D.

2.3 Data analysis

A linear viscoelastic version of the three dimensional generalized self-consistent

method of Christensen et al. [18] is used to couple the amount of space filling, caused

by point-like nucleation and spherulitic growth, to the measured dynamic (or complex)

modulus. See also the work of Christensen et al. [19,20]. This model has been validated

experimentally in our previous work [21]. If spherulites are formed, the relative dynamic

modulus f G = G/G0 can be obtained from

Af G2 +Bf G + C

= 0 , (2.1)

where G and G0 are the complex dynamic modulus of the suspension and the

amorphous phase, respectively. The complex coefficients A, B, and C depend on

space filling , the ratio of the complex moduli of the amorphous (G0) and crystalline

phase (G1), and the Poisson ratios of both phases, 0 and 1. Notice that all moduli

are frequency and temperature dependent. Expressions for the coefficients are given in

Appendix A of Steenbakkers et al. [21]. In this case, space filling is the unknown and

is obtained by sloving Eq. (2.1) using the measured f G.

Next the space filling has to be related to the nucleation density. For a fixed number

density of nuclei N(T ), the Kolmogorov Avrami Evans equation [2225] describes the

Chapter 2 13

progress of space filling in time during isothermal crystallization at temperature T ,

(t, T ) = 1 exp(4pi

3N(T )G3(T )t3

)(2.2)

from which the nucleation density can be obtained:

N(T ) = 3 ln (1 )4piG3(T )t3

(2.3)

with growth rate G(T ) depending on the temperature only. The main requirement here

is that the growth rate is independent of the nucleation density, even when this density

is influenced by a nucleating agent or by flow. Flow experiments are so-called short

term shear experiments; the flow can generate extra nuclei but the growth takes place

after the flow has stopped. The growth rate for iPP is well known [26]. Note that there

are different empirical and theoretical relations describing the temperature dependence

of the growth rate, see Eq. (2.4) [8] and Eq. (2.5) [27]:

G(T ) = Gref exp[cG(T TG,ref)2] (2.4)

G(T ) = G0 exp

[ U

R (T Tg + T)]

exp[GT

2m (Tm + T )

2T 2 (Tm T )]

(2.5)

One could think of using other approaches than Eq. (2.1) by applying more simple

models [2832] or empirically based scaling laws between the space filling and the storage

modulus [3335]. However, it was demonstrated by Steenbakkers and Peters [21] that

such approaches do perform less than the suspension model as used here, and therefore,

we will not apply them to our results. Summarizing, by measuring the complex modulus

G(t, T, , ts) for different temperatures, shear rates , and shear times ts, we can

determine the space filling (t, T ) using Eq. (2.1) and the nucleation density N(T, , ts)

using Eq. (2.3), provided that only spherulitic growth from predetermined point-like

nuclei occurs.

14 Chapter 2

2.4 Results and Discussion

2.4.1 Determination of the nucleation density for NA-iPP

First of all, Figure 2.1 shows that the structures formed during crystallization with

nucleating agent do not show any unusual structure as was found by Lee Wo et al. [16] for

a similar system (a commercial iPP with different concentrations of U-Phthalocyanine),

so the above method can be applied to convert the time evolution of the rheological

properties into kinetics of space filling. Figure 2.2 shows the time evolution of the

storage modulus G of the nucleating agent system during crystallization at different

temperatures. It should be noticed that for T = 145 C the initial value of G is not

the same as those for the other temperatures. The method is based on the assumption

of isothermal crystallization. If the degree of undercooling is large with respect to

the cooling rate, crystallization already sets in during cooling. The problem is made

even worse due to the 2-min delay time we used in order to equilibrate the sample

temperature. This is the case for T = 145 C; crystallization has already started before

we begin to track the rheological evolution (indicated by an increased initial value of

G). Faster cooling rates are required when studying higher levels of undercooling.

Nevertheless, we still want to show this result to demonstrate the limitations of the

method.

10 100 1000 10000

10

4

10

5

151

o

C

148

o

C

147

o

C

145

o

C

10

6

10

7

10

8

G' (Pa)

time (s)

Figure 2.2: Time evolution of the storage modulus for NA-iPP samples crystallized atdifferent temperatures.

The related space filling, determined by Eq. (2.1), is shown in Figure 2.3. For the high

experimental temperature range used here, the growth rate is better captured by Eq.

(2.4) [8, 26], so this expression was applied. Figure 2.4 shows the diffraction patterns

of NA-iPP samples after quiescent crystallization and demonstrates that nucleating

Chapter 2 15

10 100 1000 10000

0.0

0.2

0.4

0.6

0.8

1.0

151

o

C

148

o

C 147

o

C 145

o

C

space filling

time (s)

Figure 2.3: Time evolution of space filling derived from the dynamic modulus for NA-iPPsamples crystallized at different temperatures.

9 10 11 12 13 14 15

145

o

C

147

o

C

148

o

C

(-131)

(111)

(130)

(040)

Intensity (a.u.)

2 (

o

)

(110)

151

o

C

Figure 2.4: One-dimensional WAXD curves of NA-iPP samples crystallized at differenttemperatures.

agent crystallization results in the same -modification as crystallization of the neat

iPP. Consequently, the same growth kinetics, Eq. (2.4), applies. The number density

of nuclei determined by Eq. (2.3) is plotted versus space filling in Figure 2.5. It is

nearly constant between 0.1 and 0.9, except for the experiment at T = 145 C,where crystallization has already set in before the dynamic measurements have started.

Much higher values are found in the early stage, but N(T ) becomes nearly constant

when 0.4 for all temperatures. The influence of the unknown initial space filling isrelatively smaller for higher degrees of space filling, i.e., in later stages of the process.

We have taken the average value for a space filling between = 0.5 and = 0.9 as a

reasonable approximation. Nucleation densities from all experiments are plotted as a

function of the experimental temperature in Figure 2.6. It can be concluded that this

type of nucleating agent is very effective as it increases the nucleation density by up to

16 Chapter 2

0.0 0.2 0.4 0.6 0.8 1.0

151

o

C

148

o

C

147

o

C

145

o

C

10

14

10

16

10

18

10

20

N (m

-3

)

space filling

10

22

Figure 2.5: Nucleation density versus space filling for NA-iPP samples crystallized atdifferent temperatures.

130 135 140 145 150 155

10

13

10

15

10

17

10

11

NA+iPP

iPP

N (m

-3

)

temperature (

o

C)

10

19

Figure 2.6: Nucleation density of iPP and NA-iPP versus temperature.

six decades. Notice that the temperature dependency of the nucleation density is very

similar for the neat iPP and the NA-iPP, and that the results for T = 145 C, although

less reliable for < 0.4, are well in line with the results for higher temperatures.

Results for relatively low temperatures should be treated with some caution.

2.4.2 Reproducibility

First we compare the results of Housmans et al. [9], obtained with the same method

and for the same iPP, to the results presented here. For a temperature T = 138 C,

they found, for quiescent conditions, a nucleation density N = 8 1011m3, while weobtain the (interpolated) value N = 6 1012m3. We ascribe this difference to the

Chapter 2 17

sample preparation procedure. Our neat iPP and nucleated samples were processed by

injection molding, which means that the sample preparation step already induced extra

nuclei due to the applied (uncontrolled) flow. The samples used by Housmans et al. [9]

were prepared by means of compression molding.

130 135 140 145 150 155

temperature (

o

C)

1

st

Series NA+iPP

1

st

Series iPP

2

nd

Series NA+iPP

2

nd

Series iPP

10

11

10

13

10

15

10

17

10

19

N (m

-3

)

Figure 2.7: Nucleation density of iPP and NA-iPP versus temperature for two series ofexperiments.

We repeated our experiments to check for reproducibility. The time lapse between the

two series of experiments was about 8 weeks. The results of these repeated experiments

(2nd series) are shown in Figure 2.7, together with the previous results (1st series). We

want to stress the importance of a good temperature control. An error of 1 C typically

gives a factor two difference in the number of nuclei. Notice that in the second series

of experiments we managed to get good results for temperatures as low as T = 143 C

while in the first series we already encountered problems at T = 145 C.

2.4.3 Effect of mild flow

Figure 2.8 shows the evolution of G during crystallization at 148 C under quiescent

conditions and after different flows (fixed shear rate and variable shear time). For a

shear rate of 60 s1, a shear time of 2 s shows clearly accelerated kinetics. Further

increase of the shear time to 4 s does not change the kinetics much. The acceleration

of crystallization seems to become (nearly) independent of the shear time beyond 4 s,

indicating that the shear-enhanced point-like nucleation saturates. This effect was also

observed by Housmans et al. [9] for this and two other iPPs. The time evolution of

the storage modulus of the sheared nucleating agent system has the same shape as in

the quiescent nucleating agent system, meaning that the growth mechanism is the same

and only spherulites are formed. This implies that our method can still be applied

18 Chapter 2

to determine space filling for crystallization after shear. Moreover, the results can be

compared to those of Housmans et al. [9].

10 100 1000 10000

0

2

4

time (s)

10

4

10

5

10

6

10

7

10

8

G' (Pa)

shear time (s)

6

Figure 2.8: Time evolution of the storage modulus for NA-iPP under quiescent conditionsand after short term shear ( = 60 s1, ts = 2, 4, 6 s) at T = 148

C.

For flow-induced crystallization experiments, an adequate measure of the flow

strength is the Weissenberg number, based on the stretch relaxation time of the high

molecular weight (HMW) tail of the molecular weight distribution:

Wis(T, ) = HMWs (T ) (2.6)

Recent modeling work [36] suggests that the creation rate of point-like nucleation

precursors depends not only on the average stretch of the HMW molecules, but also

explicitly on the temperature: the prefactor of the creation rate was found to be

proportional to the time-temperature shift factor aT . Therefore we use aTWis as a

criterion to compare experiments, which scales as a2T , since HMWs in Eq. 2.6 scales

with aT as well. From the linear viscoelastic data and the stretch relaxation time HMWs

reported by Housmans et al. [9], we find Wis(148C) = 13, Wis(151

C) = 12, and

aT (151C)Wis(151

C)

aT (138 C)Wis(138 C)= 1.1 , (2.7)

where Wis(138C) = 8 for the strongest flow applied by Housmans et al. [9] to the same

neat iPP as used here. Hence, these two data sets are reasonably comparable in terms

of flow conditions.

The calculated nucleation densities are plotted in Figure 2.9. The saturated

nucleation density at 148 C after shear is around 40 1016m3, six times higher than7 1016m3 for quiescent crystallization, i.e., adding typically 1017m3 nuclei. This

Chapter 2 19

0 2 4 6

10

17

10

16

10

15

10

18

151

o

C

148

o

C

N (m

-3

)

shear time (s)

Figure 2.9: Nucleation density of NA-iPP under quiescent conditions and after shear atT = 148 C and T = 151 C ( = 60 s1).

130 135 140 145 150 155

temperature (

o

C)

1

st

Series NA+iPP

1

st

Series iPP

1

st

Series shear

2

nd

Series NA+iPP

2

nd

Series iPP

10

11

10

13

10

15

10

17

10

19

N (m

-3

)

Figure 2.10: Nucleation densities of iPP and NA-iPP under quiescent and shear conditions.

elevated number is much more than what was added by imposing shear to the neat iPP

in the experiments of Housmans et al. [9]; mild flow raised the point-like nucleation

density of the neat iPP by one to two decades before the shear effect saturated, i.e.

adding typically 1013m3 nuclei. This indicates that, in the presence of nucleatingagent, shear-induced point nucleation can be much more effective. This should be

directly related to the nucleating agent, i.e., the local amplification of flow effects due

to the presence of the nucleation particles, see for example the work of Hwang et al. [37],

and not to a change in the rheological properties, which we expect to be very small due

to the, in this respect, still very low space filling of nuclei.

20 Chapter 2

2.5 Conclusions

A new method to determine nucleation densities, based on rheometry, was used

on quiescent and sheared samples of a neat and an artificially nucleated isotactic

polypropylene. All calculated nucleation densities are summarized in Figure 2.10. It

is quantitatively shown that U-Phthalocyanine is very effective for nucleating isotactic

polypropylene. Moreover, it was found that the effect of shear is enhanced by the

presence of the nucleating agent. This rheological method is easy to apply since it

requires a standard rheometer, available in most academic and industrial labs.

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Chapter 2 21

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8398, 2006.

Chapter three

Pressure quench of flow-induced

crystallization precursors

Abstract

We developed a novel protocol to study the mutual influence of shear flow and

pressure on crystallization of polymers. Here, we have applied this protocol, named

pressure quench, to polyethylene with a bimodal molecular weight distribution. With

pressure quench, the undercooling, required to initiate crystallization of flow-induced

precursors generated at high temperature, is obtained by increasing pressure, i.e.,

leaving the specimen isothermal. We find that pressure enhances the effect of shear. In

particular, results show that the pressure quench effectively lightens up shear-induced

precursors which otherwise are not observable, even with high-resolution synchrotron

X-ray scattering. A pressure quench in combination with SAXS and WAXD gives

insight into the early stages of crystallization. In this chapter we focus on the use of

WAXD since it provides all the information required to demonstrate our main issues.

We conclude that precursors with different stability can be formed during shear and,

with annealing, the least stable ones relax back to the melt. Finally, it is demonstrated

that when pressure is released after crystallization, an inverse quench takes place and

crystalline structures partially melt, similar to an increase of the temperature.

This chapter is based on : Zhe Ma, Luigi Balzano, Gerrit W.M. Peters. Macromolecules 45:42164224, 2012

23

24 Chapter 3

3.1 Introduction

Semicrystalline polymers represent the largest group of commercial polymeric

materials. Their properties strongly depend on the crystalline structures that form

during processing. [14] Apart from molecular features, the structures formed are

controlled by processing variables such as flow, pressure, and temperature. These

variables are often investigated individually, but there is also a remarkable mutual

influence that is still unexplored. Laying ground for unraveling the relation between

pressure and flow is the topic of this chapter.

It is established that shear flow is able to generate precursors of crystallization in

molten polymers. [58] Flow-induced precursors are metastable domains where stretched

molecules assume a packing order in between the amorphous and the crystalline states.

They can be considered as the cradle of flow-induced nuclei as, in suitable conditions,

e.g. at low temperature, they nucleate and grow into crystalline structures. In this way,

crystallization kinetics is accelerated, and the final morphology of the material is altered.

Therefore, monitoring formation and evolution of crystallization precursors is crucial to

rationalize the effect of flow on the final morphology of polymeric manufactures.

Precursors are often invisible to scattering techniques such as SAXS and WAXD

because their structure is neither crystalline nor densely packed or their concentration

is too low to produce significant effects. In these cases, they are studied indirectly by

interpreting features (such as kinetics and morphology) of the subsequent crystallization

process. [915] Interestingly, precursors can be generated and survive for long times

also when temperature is around and sometimes above the melting point. [9, 1114]

Especially in these cases, cooling to a lower temperature is a convenient way to trigger

crystallization. For instance, Hsiao et al. [9] found that, for polyethylene (PE), after

flow ( = 20 s1, ts = 12 s) at 134C, no structure could be observed with WAXD.

Nevertheless, after bringing the system at 129 C, oriented crystals started to grow.

Practically, the cooling step should be fast and the final temperature not too low in order

to keep the nucleation and growth separated on the time resolution of the experiment.

This is commonly achieved by cooling with a liquid or gas medium. Small lab-scale

devices, such as the commercially available Linkam shear cell, can reach cooling rates

up to 30 C/min. Whereas, flow cells designed for strong flows and high pressures are

typically restricted to lower values as they are constructed out of relatively large masses

of metal. For example, the average cooling rate for the original design of the Multi-Pass

Rheometer (MPR) is 1 C/min. [16]

When performing flow around the melting point followed by cooling to a lower

temperature, the lifetime of precursors becomes an important parameter as for long

cooling times (compared with the lifetime); a fraction of the precursors dissolves and

thus mitigates the apparent effect of flow on crystallization. Balzano et al. [17] suggested

Chapter 3 25

that, shortly after flow unstable structures dissolve on the time scale of the reptation of

the longest molecules of the melt. Therefore, experimental protocols are required that

allow for studying the details (number, size, morphology, and dynamics) of precursors

in the very early stages while having control over their relaxation behavior. The goal

is to separate the nucleation from the growth step. In this light, cooling directly to

room temperature is often not a valid option since the fast overgrowth of crystalline

structures would obscure nucleation. On the other hand, performing experiments at

high temperature, close to the nominal melting temperature, seems a better solution

as growth is very slow in these conditions. The drawback is that structures created at

high temperatures are very sporadic and easily fall beyond the detection limits even of

high-resolution methods such as synchrotron X-ray scattering.

An alternative way to perform controlled quenching is to utilize a pressure quench. In

other words, apply pressure and, in this way, shift the melting temperature (according

to the ClausiusClapeyron relation) and thus effectively increase the undercooling

without actually changing the sample temperature. This methodology can be directly

implemented in certain slit-flow devices such as the one developed in Eindhoven [18]

(mounted on the MPR) and offers some clear advantages:

The cooling step can be instantaneous.

Temperature gradients and undershoots can be avoided (especially important for

thick samples).

The cooling step is easily reversed by depressurizing so complex undercooling

histories can be applied simply by varying pressure.

Pressurization has been used as an alternative way to shift the phase boundary

in studies on phase separation. [19, 20] In particular, it was found that during phase

separation of polymer blends the general features of nucleation are independent of

whether the undercooling is obtained by decreasing temperature or by increasing

pressure. [20] Therefore, pressure quench provides an effective way to obtain

undercooling. Moreover, it also reflects the practically important, mutual influence

between flow and pressure in polymer processing techniques.

In this chapter, we demonstrate the use of pressure quench to investigate the early

stages of nucleation of flow-induced precursors. We focus on a bimodal blend of PE

containing 3 wt% high molecular weight tail to simplify the rheological classification [21]

of the flow conditions and to enhance the flow-induced formation of crystallization

precursors. [2228] Only a relatively mild increase of pressure is used in this work,

since a few degrees of undercooling is already effective to accelerate the flow-induced

crystallization growth. Moreover, a too high pressure (on the order of kbar) may

induce the hexagonal phase [29, 30] of PE, which is beyond the scope of this work.

26 Chapter 3

Theoretical aspects on the origin of flow-induced precursors are not discussed here.

Details can be found in a previous paper of our group. [31] For example, the issue

if flow creates precursors directly from the melt or if precursors (so called dormant

precursors [32, 33]) are always present and flow only changes their size to increase the

number that can be activated by a quench, was dealt with in detail.

3.2 Experimental

3.2.1 Materials

A bimodal polyethylene blend containing 3wt% of high molecular weight tail was

used in this work. The ultrahigh molecular weight polyethylene (UHMWPE) has a

weight-averaged molecular weight Mw = 1480 kg/mol and polydispersity Mw/Mn = 2.

[34] The linear low molecular weight polyethylene (LMWPE) matrix, supplied by Basell

Polyolefine GmbH (Frankfurt, Germany), has a Mw = 45 kg/mol and polydispersity

Mw/Mn = 3. The critical overlap concentration of high molecular weight molecules can

be calculated with [35, 36]

c =3Mw

4piRg23/2NA(3.1)

where Rg2 is mean-square radius of gyration of chain related to the molecular weight byRg21/2 = 0.46Mw1/2, [37] is the density andNA is Avogadros number. The estimatedcritical concentration of UHMWPE is around 0.35wt%, i.e., much smaller than the

3wt% in our bimodal blend, meaning that a significant number of entanglements exist

between the UHMWPE molecules.

The bimodal system was prepared by solution blending to achieve mixing at a

molecular scale. The UHMWPE was first dissolved in a xylene solution at 130 C, and

subsequently LMWPE was added to dissolve, where the concentration of total PEs is

2.5% with an antioxidant (IRGANOX1010) added at a concentration of 2000 ppm. This

solution was stirred for 1 h under a nitrogen atmosphere. Next, the hot xylene solution

was poured into a large excess of stirred cold methanol. The precipitated gel was filtered

and washed with methanol several times and then dried in vacuum at 80 C for 2 days.

After further addition of 2000 ppm of antioxidant (IRGANOX1010), in order to avoid

degradation during sample preparation, the bimodal blend was compression molded at

160 C into strips to fit the test cell.

Chapter 3 27

3.2.2 Protocol

The slit flow device developed in Eindhoven [18,38] is an evolution of the Multi-Pass

Rheometer of Eland Engineering Co Ltd. (UK) [39] (see Figure 3.1). The flow cell is

specifically designed for online scattering and therefore equipped with diamond windows

(opening angle of 45). The specimen (L = 200mm, W = 6mm, and H = 1.5mm)

is confined between two servo hydraulically driven rectangular pistons. An important

advantage of this slit-flow device is the possibility to impose and release pressure. The

maximum pressure is 800 bar and the maximum temperature 250 C. The procedures

for flow, pressurization, and depressurization are illustrated in Figure 3.1.

(a) (b) (c)

Figure 3.1: Schematic of the flow device and flow, pressurization, and depressurizationprocedures. Moving directions of the pistons are indicated by arrows. Heatingrods and the heating/cooling channels are not shown.

Moving the two pistons in the same direction introduces a shear flow to the sample

(Figure 3.1a), while moving the two pistons toward or away from each other will

pressurize (Figure 3.1b) or depressurize (Figure 3.1c) the sample, respectively, without

causing any flow at the observation point (center of the slit). The pressure difference

during flow is recorded by means of two pressure transducers.

The experimental protocol used in this chapter is shown in Figure 3.2. The polymer

in the test cell is first heated up and annealed at 190 C for 10 min to erase the memory

of previous thermal and mechanical histories and then cooled to 134 C at which shear,

pressurization, and depressurization are performed. To avoid temperature fluctuations,

the cell is stabilized by means of an oil bath, while the top and bottom barrels are

always kept at high temperature (190 C) in order to have the pressure transducers

functioning properly. Next, flow is imposed with a piston speed of 15mm/s for 0.8 s.

28 Chapter 3

time

pre

ssure

shear

tem

pera

ture

ta

Figure 3.2: Experimental protocol. The annealing time ta is between cessation of flow andpressurization.

The shear rate and Weissenberg number (Wi) distributions over the slit thickness were

estimated considering the rheology of the material and are shown in Figure 3.3 (see also

the Appendix). The shear rate varies nonlinearly from 0 at the slit center to a maximum

of 67 s1 at the wall. The Weissenberg number, Wi = Rouse, is used to estimate the

molecular stretch induced by flow. When Wi > 1, molecules become stretched, and

this is a requirement for the creation of precursors. [40] The longest Rouse times of

UHMWPE and LMWPE at 134 C are 4.9102 and 4.5105 s (see section 3.3.3 forthe calculation), respectively, so the Weissenberg number for molecular stretch at the

wall is 3.3 for HMW tail but nearly zero for LMW matrix.

In order to obtain good filling of the slit and rule out wall slippage, a reference

pressure of 50 bar was kept on the specimen. The average pressure during flow,

PAverage = Plower +12Psensor, where Plower and Psensor are the lower pressure (around

15 bar) and the pressure difference between the two pressure sensors (see Figure 3.4),

is around 70 bar totally, so the influence of flow on the pressure level is negligible.

A pressure of 300 bar was chosen as the elevated pressure level to increase the melt

undercooling. Figure 3.2 also shows that pressurization can be applied right after flow

(ta = 0) or after annealing (ta = 22 min) with the same flow. Both protocols were

employed in this work. The former (ta = 0) is used to see if any precursors were formed

during flow, and the latter (ta = 22 min) is to show how precursors develop with time.

Chapter 3 29

centerwall wall

0.15 0.30 0.45 0.60 0.75

0

1

2

3

4

LMWPE matrix

distance from center (mm)

UHMWPE

We

issen

be

rg n

um

be

r

0.75 0.60 0.45 0.30 0.15 0.00

0

20

40

60

80

distance from center (mm)

she

ar

rate

(1/s

)

Figure 3.3: Distributions of shear rate and Weissenberg numbers for stretching HMW tailand LMW matrix over the slit thickness direction.

3.2.3 X-ray characterization

X-ray measurements were carried out at the Dutch-Belgian (DUBBLE) beamline

BM26B of the European Synchrotron Radiation Facility (ESRF) in Grenoble, France,

with a wavelength of 0.95 A. [41] A Frelon detector with a resolution of 2048 2048pixels of 48.8m 48.8m, and placed at a distance of 0.195m, was used for wide-angle X-ray diffraction (WAXD). The measuring time of each WAXD frame, including

exposure and readout of data, was 8.5 s.

WAXD images were processed with the software package FIT2D to obtain intensity

versus scattering angle (2) profiles. Crystallinity was calculated after deconvolution

of the total intensity scattered by the crystalline (Acrystal) and amorphous (Aamorphous)

domains:

X =Acrystal

Acrystal + Aamorphous 100% (3.2)

Flow-induced crystallization depends on the strength of the local flow field. In the

slit geometry, the shear rate changes from a minimum in the center to a maximum at

the wall, as shown in Figure 3.3, so shear-induced crystalline layers will form mostly

next to the two walls. X-ray going through the whole sample thickness will scatter from

the two shear layers, and therefore, values calculated in this way do not represent the

real crystallinity in certain specific layer but rather an apparent crystallinity averaged

over the optical path of the X-ray beam through whole sample.

30 Chapter 3

3.3 Results and Discussion

3.3.1 Reference experiment: no pressure quench after flow

The formation and development of flow-induced precursors are investigated with a

shear pulse ( = 67 s1, ts = 0.8 s) under 50 bar at 134C. First of all, the structural

development during flow was examined through the rheological response of the polymer

melt. Figure 3.4 shows the evolution of pressure difference during flow, which first

increases rapidly, subsequently shows a small overshoot, and reaches a steady-state

plateau. This behavior can be related to the usual nonlinear rheological response of

the sheared molten polymer. Scelsi et al. [42] reported a buildup of pressure difference

(HDPE, 130 C, Wi for stretch is estimated around 60) during flow in a channel with

a contraction in the middle section and associated it to continuous crystal formation

concentrated in the region of slit exit. Such a pressure difference buildup during flow

was not observed in our experiments (134 C, Wi = 3.3). Therefore, our results imply

that crystals do not form during flow or that the viscosity does not change even if

precursors form (as discussed later).

0.0 0.2 0.4 0.6 0.8

0

20

40

60

80

100

120

140

160

pressure difference (bar)

time (s)

experiment 1

experiment 2

Figure 3.4: Pressure difference during flow. Two experiments are shown to demonstrate thereproducibility of the flow history.

Based on the nano- and mesoscale structures of crystalline planes and density

differences, both SAXS and WAXD were employed to examine, at different length scales,

structural aspects of flow-induced precursors. However, in-situ synchrotron SAXS and

WAXD data acquired after flow did not show any observable signal, which confirms

that no detectable crystallization takes place neither during flow nor within the next 20

min when the sample is kept isothermal (data not shown since no detectable signal was

found). More in detail, despite the Wi number for stretching the HMW tail is larger

Chapter 3 31

than unity in the region close to the wall (see Figure 3.3); neither crystalline nor densely

packed scatterers are observed. This can be a consequence of a low concentration of

precursors, i.e., below the experimental detection limit. Therefore, crystallization has to

be triggered to reflect the features of precursors. For PE, WAXD is sufficient to fulfill

this purpose of characterizing crystallization features (i.e., orientation, crystallinity,

twisted lamellae) which will be shown in the following results. Therefore, we use only

WAXD [43] to track PE crystallization and show that a pressure quench is a suitable

method to visualize precursors even at these very low concentrations.

3.3.2 Pressure quench after flow

In order to investigate the formation of flow-induced precursors, after the application

of the same flow as in section 3.3.1, a pressure quench, with the pressure history shown

in Figure 3.5, is performed. The pressure is raised linearly from the reference pressure

of 50 to 300 bar without overshoot. The synchronous top and bottom pressures indicate

that the pressure field is homogeneously distributed over the specimen and that there

is no flow at the observation point (center of the slit). With pressure, undercooling is

generated and immediately triggers crystallization, illustrated by the WAXD images in

Figure 3.6.

0 10 20 30

0

100

200

300

pressure - top

pressure - bottom

pressure (bar)

time (s)

Figure 3.5: Pressure profile during a pressure quench.

As soon as the maximum pressure is achieved, arched (110) and (200) reflections

appear in the equator direction (Figure 3.6a), indicating the formation of the

orthorhombic unit cell with c-axis aligned in the flow direction. According to Keller

et al. [6], the diffraction pattern of Figure 3.6a can be produced both by the extended

chain crystals of shish and by the folded chain crystals of untwisted lamellae.

32 Chapter 3

(a) tc = 0s (b) tc = 8.5s (c) tc = 17s (d) tc = 102s

Figure 3.6: 2DWAXD patterns under a pressure quench after flow. Flow is along the verticaldirection. The reference time is the moment when 300 bar is reached.

For the purpose of this work, it is not essential to make such a distinction. The most

important observation is the presence of highly oriented crystals. In fact, as the sample

is pressurized in absence of extra flow, these oriented crystals originate from deformed

molecules involved in precursors induced by the flow during the shear step.

As crystallization proceeds, the (110) reflection broadens in the azimuthal direction

(Figure 3.6b) and, at tc = 17 s, splits into off-axis diffraction together with the clear

appearance of off-axis (200) diffractions (Figure 3.6c). These off-axis (110) diffractions

are the result of the lateral growth of twisted lamellae. During this process, the

lamellar propagation direction holds perpendicular to the flow direction, but crystal

units rotate along the b-axis. [6] Twisting leads to randomization in the orientation

of c-axes. Keller et al. [6, 44] proposed two extremes of lamellar orientation for flow-

induced PE crystallization: Keller/Machin I and Keller/Machin II. The former

(KM-I) corresponds to the fully twisted lamellae producing off-axis (110) and meridional

(200) diffractions. The latter refers to the flat, nontwisted lamellae (all c-axes parallel

with flow direction) producing equatorial (110) and (200) diffractions. In the transition

from one to another an intermediate state is observed showing the off-axis (110) and

(200) patterns. [6, 45] Therefore, the azimuthal features of the (200) diffraction reflect

the lateral growth of lamellae. [6, 9, 45]

In a later stage, the diffraction of isotropic structures is also observed, e.g. isotropic

(110), (200) diffractions at 102 s (see Figure 3.6d). These two crystallization processes,

fast oriented crystallization and slow isotropic crystallization, characterized by different

degrees of orientation and time scales can be rationalized by considering that the

inhomogeneous flow field in the slit gives rise to inhomogeneous molecular stretch as

shown by Figure 3.3. At the wall, where the stress reaches a maximum, molecules

experience the largest stretch and many flow-induced precursors are generated. These

are the domains where highly oriented crystals are formed upon pressurization. Toward

the centerline, molecules experience little or no stretch and the probability of forming

oriented precursors vanishes. In this part of the sample, isotropic crystallization can

be triggered purely by pressure. The WAXD images (Figure 3.6a-d) are indicative of

Chapter 3 33

structures averaged over the optical path of the X-ray beam (thickness of the samples),

and therefore, they contain information on both these two processes.

Next, the effect of precursors on crystallization kinetics is illustrated in Figure 3.7a

where the apparent crystallinity is plotted as a function of time for shear-induced

crystallization under a pressure quench and two reference experiments (shear-induced

crystallization without pressure quench and quiescent crystallization with pressure

quench). It is evident that when shear is applied, crystallization starts as soon as

the 300 bar pressure is reached. On the other hand, without shear, crystallization

starts only after 25 s from reaching 300 bar. Eventually, without applying pressure,

crystallization of the sheared polymer does not occur in the experimental time window

of 20 min. The results show that there is a remarkable interplay between flow and

pressure on crystallization kinetics and morphology.

0 20 40 60 80 100

0

5

10

15

flow + Pressure Quench

quiescent + Pressure Quench

flow

apparent crystallinity (%)

time (s)

(a) (b)

Figure 3.7: (a) Crystallinity evolution for shear-induced crystallization with pressure quench(), quiescent crystallization with pressure quench (), and shear-inducedcrystallization without pressure quench (). (b) 2D WAXD pattern of quiescentcrystallization under pressure quench at tc = 34 s.

From the above results, we conclude that oriented precursors are generated by the

flow, but prior to pressurization, they are invisible to X-rays (see section 3.3.1). It is

clear that a pressure quench effectively provides the additional undercooling that triggers

crystallization by raising the equilibrium melting temperature. The equilibrium melting

temperature of PE at 1 bar is T 0m = 414.6 K. [46] Its increase with pressure, described

by the Clausius-Clapeyron relation, can be approximated by T (p) = T 0m+(dTpm/dp)p,

where dT pm/dp = 35.2 K/kbar. [46] In this way, the equilibrium melting temperatures

corresponding to 50 and 300 bar are calculated as 416.4 and 425.2 K, respectively.

With pressurization from 50 to 300 bar, the undercooling T pexp = T pm Texp increases

34 Chapter 3

from T 50bar134 9K to T 300 bar134 18K (experimental temperature 407.15 K). Thus, apressure quench from 50 to 300 bar increases the undercooling by 9 K within 25 s. Such a

pressure quench, around 22 K/min, is comparable to a temperature quench in small lab-

scale device. On the other hand, it is much more efficient than lowering the temperature

of such large metal device with a cooling medium, especially since the cooling rate

decreases when approaching the target temperature, and thus it usually takes much

longer time (minutes) to stabilize at the desired undercooling (see the Appendix).

The azimuthally homogeneous (110) diffraction in Figure 3.7b shows that pressurizing

does not cause flow at the observation point which might have influenced crystallization.

Concluding, oriented crystals (Figure 3.6) and faster kinetics (Figure 3.7a) show

that precursors can form in a sheared melt, where only the HMW chain is stretched.

This demonstrates that the HMW tail determine the formation precursors, consistent

with the findings of Mykhaylyk et al. [15] A pressure quench triggers crystallization

and effectively lightens up these flow-induced precursors which could not be detected

otherwise for these conditions (134 C, 50 bar).

3.3.3 Pressure quench after annealing

The effect of partial dissolution of flow-induced precursors can be semi-qualitatively

estimated by introducing an annealing step (during which pressure is kept at 50 bar)

between the flow pulse and the pressure quench. During the annealing step, whose length

is arbitrarily fixed at 22 min, no structural changes were observed by SAXS/WAXD

(data not shown). Figure 3.8 shows the 2D WAXD patterns after applying a pressure

quench to the sheared and annealed specimen. The equatorial (110) diffraction in Figure

3.8a indicates the survival of orientation in the precursors with annealing.

(a) tc = 0s (b) tc = 8.5s (c) tc = 34s (d) tc = 93.5s

Figure 3.8: 2D WAXD patterns of structures after flow, 22 min annealing under 50 bar anda final pressure quench. Flow is along the vertical direction.

The difference with crystallization without annealing (section 3.3.2) is the meridional

(200) diffractions (Figure 3.8c) that arises from a higher degree of randomization in the

Chapter 3 35

c-axes of the unit cells. [6] This arises from the fact that, in the annealed sample, the

lamellae can grow laterally over a longer distance; i.e., their nucleation points are further

apart than in the nonannealed sample and also means that the annealed sample has a

lower nucleation density.

This observation suggests that the lifetime of some of the flow-induced precursors is

longer than 22 min (at 134 C and 50 bar), whereas for others it is not.

The question is whether the long lifetime of precursor relates to the rheological time

scales of the material. Considering the molecular weight between entanglements Me =

828 g/mol, [47] the numbers of entanglements per chain, Z, for HMW tail and LMW

matrix are 1787 and 54, respectively. Because of the significant degree of overlap of long

molecules and of the low StruglinskiGraessley number Gr = ZHMW/Z3LMW = 0.01,dynamic tube dilution can be neglected. [47] In other words, the relaxation time of

HMW molecules is not reduced by the LMW matrix. According to the tube model,

the Rouse time Rouse and reptation time D, responsible for stretch and orientation

relaxations, can be calculated from

Rouse = eZ2 (3.3)

D = 3eZ3

(1 1.51

Z

)2(3.4)

where e is the entanglement equilibration time, around 7109 s for PE at 190 C. [47]Without considering molecular weight distribution and the effect of 50 bar pressure,

the estimated reptation times of the HMW tail and LMW matrix at 134 C (Ea =

21.8 kJ/mol) are 243 and 0.005 s, respectively, and the Rouse times are 4.9 102and 4.5 105 s, respectively. The striking feature is that the longest relaxation time,DHMW 4 min (predicted with Mw), is much shorter than the annealing time, 22min.

Concerning precursor relaxation, previous experimental studies [17, 38, 48] showed

that the relaxation of the most unstable shear-induced bundles, observable with

SAXS, follows the reptation dynamics of the longest chains whereas stable precursors

survive on time scales much longer than the rheological ones.

Systematical studies on relaxing shear-induced nucleation precursors that are

invisible to SAXS were done by Alfonso and co-workers. [14] They found that, in iPP,

at temperatures as high as 190 C (above nominal melting temperature but not beyond

equilibrium melting temperature), nucleation precursors can survive much longer than

the longest rheological relaxation times, and that the characteristic survival time can

be increased by a stronger or longer flow. The change in relaxation times implies that

other effects are determining the dissolution of precursors. This idea is supported by the

different relaxation behaviors of the flow-induced helices in iPP as observed by An et

36 Chapter 3

al. [49] They suggested that interactions between flow-induced helices of iPP dominate

the dissolution rates and the helices with interactions relax slower than those without.

Considering the long lifetime of X-ray unobservable oriented precursors, in our

results, it is reasonable to infer that interactions between PE chain segments

(comparable to very local crystallization events) contribute to the long lifetime beyond

the rheological times. According to the classical nucleation theory, a precursor below the

critical size tends to relax; the interaction decreases the relaxation kinetics. Once the

total interaction (contributed by the increased volume as well) is sufficient, the volume

free energy overcomes the surface energy, and the precursors can develop to nuclei that

are stable. Long-term stable precursors were also observed by Mykhaylyk et al. [15] for

hydrogenated polybutadiene; the precursors survived within the experimental window

which was as long as 10 h. Interestingly, in our experiments, when these segmental

interactions are initiated during flow, they do not cooperate significantly to change the

viscosity of the whole melt as observed from the pressure difference in the presence of

precursors (see Figure 3.4).

0

10

20

30

40

50

60

70

80

90

10

20

30

40

A

z

i

m

u

t

h

a

l

a

n

g

l

e

(

o

)

I

n

t

e

n

s

i

t

y

(

a

.

u

.

)

time

0s

110s

51s

(a)

0

10

20

30

40

50

60

70

80

90

10

20

30

40

A

z

i

m

u

t

h

a

l

a

n

g

l

e

(

o

)

I

n

t

e

n

s

i

t

y

(

a

.

u

.

)

time

0s

85s

(b)

Figure 3.9: Change in the azimuthal distribution of (200) diffraction during pressure quenchcrystallizations for (