@
........... ..
MASSACHUSETTS INSTITUTE
OF TECHNOLOGY
SEP 2 0 2000

by
Rami Lokas
B. S. Mechanical Engineering, Georgia Institute of Technology, 1998
Submitted to the Department of Mechanical Engineering on May 8, 2000, in partial fulfillment of the requirements for the degree of
Master of Science in Mechanical Engineering
Most descriptions of polymers start at room temperature and end at the melting point. Cryogenic testing is rare for even the most common polymers. Considering the increased use of polymers at low temperatures (eg: thermal and electrical insulations, support elements for cryogenic devices, low-loss materials for high-frequency equipments) this seems to be a great lack. This thesis seeks to provide data on the behavior of several polymers under low temperature testing. The polymers tested are high density polyethylene, polyvinyl chloride, polypropylene, and polyetherimide.
The mechanical tests they underwent were compression tests and compact tension tests. The temperature range was from -150'C to room temperature. The yield and stiffness values as a function of temperature are presented. They were all found to be increasing with decreasing temperature.
Several parameters determine the fracture behavior of ploymers : temperature, time, plasticity, chain orientation, and adiabatic heating. The main topic of these investigations is the temperature dependence. Time and loading rate where kept constant at 2 mm/min. in all the tests. The polymers were as recieved in an almost isotropic form. The fracture toughness of the polymers increased with decreasing temperature except for polyetherimide. However, HDPE and PVC had a precipitous drop at tests below the glass transition.
Fractography was used to study and understand the process of crack propagation.
In most cases there were well defined regions of ductile and brittle crack propagation with clear transitions. These regions correlated well with the load displacement plots of the fracture tests. Explanations were ventured to explain the morphological features of the fracture surfaces such as yielding due to adiabatic heating and crack bifurcation due to crack velocities and stress distributions.
Thesis Supervisor: Ali S. Argon
Title: Professor
2

I want to acknowledge my mother, Dr. Olivia Shenouda, and my father, Mr. Farouk
Lokas, to whom I am eternally indebted. I cannot forget my love, Christine Chen, who continues to cancel my debts and my brother, Karim Lokas with whom I am currently developing my debt. I am greatly privileged to have my name written on the same page as Prof. Ali S. Argon. I do not think I am deserving of such an honor.
I have not only learned much about materials behavior from him but I have also learned much about life and civil human interaction.
I want to thank all the professors from whom I have learned : Prof. M. Boyce,
Prof. D. Parks and Prof. L. Anand. I am also thankful for the best office mates anyone could ask for : Kevin Bass, Hang Qi, Heather Dunn, Franco Capaldi, Matts
Danielsson, Rebecca Brown, Ethan Parsons, Jin Yi, Steve Xia, Jeremey Gregory, Greg
Nielson, Yu Qiao, Cheng Su, Harish Rajaram, Jinchul Hong, Brian Gearing, Prakash,
Tom Arsenlis, Jennifer Shin, and the friendliest Una Callinan. They made my stay fun and they so willingly offered their help whenever I needed it. I couldn't have done this thesis without them (especially Qi Hang, Greg Nielson and Una Callinan).
I also want to thank my friends who encouraged me along the way and offered me unquantifiable support: Brad Geving, Kurt Romondt, Ese Adebayo, Scott Davis,
William Ochan, the Wilson family, Tom Lin, the Boyle family, Brian Johnson with a specially heart fealt thank you to Darbon Go a friend who is more than a brother.
Finally I want to thank my Lord and Savior Jesus Christ who made this wonderful creation for us to study and admire and wonder about. Without whom nothing that is, is.
3

1 Introduction 10
2 A Review of Polymers and Fracture Mechanics 13
2.1 Overview of Crystalline and Amorphous Thermoplastics . . . . . . .
13
2.2 HDPE . . . .... ... . .
... . .. . . . .. . . . . . . . .. ..
... 14
2.3 PVC . . . . .. .. .... .... . . . . . .. . . .. . . . . . . . . . . .
15
2.4 U ltem . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15
2.5 Polypropylene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15
2.6 Literature Review of Polymer Fracture Testing . . . . . . . . . . . . . 16
2.7 Review of Fracture Mechanics . . . . . . . . . . . . . . . . . . . . . . 18
3 Experimental Techniques 22
3.1 The Test and Equipment . . . . . . . . . . . . . . . . . . . . . . . . . 22
3.2 ASTM Standards . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22
3.3 Evaluation of Temperature Dependent Material Properties . . . . . . 23
3.4 Stress Intensity Factor Calculations . . . . . . . . . . . . . . . . . . . 25
4 Temperature Dependence of Fracture Toughness
4.1 HDPE .. .
........ ... . .. . .. .
... . .. . . . .. . .. .
30
30
4.1.1 Crack Tip Opening Displacement . . . . . . . . . . . . . . . . 33
4.1.2 Crack Velocity . . . . . . . . . . . . . . . . . . . . . . . . . . 34
4.2
4.1.3 Modified CT Specimens . . . . . . . . . . . . . . . . . . . . . 36
PVC ... ..... ........ ... . . . . . . . . . . . . .. . . . . . . .
37
4

4.3 Ultem . . . .. ... ... .... .. ... . . . . . . . . . ... . . . . . .
38
4.4 Polypropylene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39
5 Fracture Surface Topographies 77
5.1 Specimen Preparation . . . . . . . . . . . . . . . . . . . . . . . . . .
77
5.2 HDPE . . . ... ... ..... .. . .. . . . .. . . . . .. . .. . . . .
77
5.2.1 Fracture Surface Topography . . . . . . . . . . . . . . . . . .
77
5.2.2 Adiabatic Heating . . . . . . . . . . . . . . . . . . . . . . . . . 79
5.2.3 Brittle Fracture Surface . . . . . . . . . . . . . . . . . . . . .
81
5.2.4 Causes of Bifurcation . . . . . . . . . . . . . . . . . . . . . . . 81
5.3 PVC . . . . ..... .... ... . . . . . . . . . . . . . . . . . .. . . .
86
5.3.1 Transition Crack Length . . . . . . . . . . . . . . . . . . . . . 87
5.4 U ltem . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 87
5.5 Polypropylene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 88
6 Conclusions and Recomendations 108
6.1 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 108
6.2 Suggestions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 109
5

3-1 Comparing two methods of displacement measurements in a -50
0
C
HDPE Compact Tension Test (COD gage vs. crosshead displacement) 26
3-2 The dimensions of the compact tension specimens . . . . . . . . . . . 27
3-3 The compression test setup inside the temperature chamber (the screws were not tightened so that the PVC placed inside can be seen so as to understand the placement of the specimens) . . . . . . . . . . . . . . 28
3-4 Modifying the data from the compression setup by factoring-in the stiffness of the setup (the above plot is the unmodified data and the lower plot is the modified data) . . . . . . . . . . . . . . . . . . . . . 29
4-1 Compression tests on HDPE at various temperatures . . . . . . . . . 43
4-2 Temperature dependence of yield strength in HDPE . . . . . . . . . . 44
4-3 Temperature dependence of elastic modulus in HDPE . . . . . . . . . 45
4-4 Unmodified plots of fracture toughness tests on HDPE . . . . . . . . 46
4-5 Modified plots of fracture toughness tests on HDPE . . . . . . . . . . 47
4-6 The bifurcated cracks in HDPE . . . . . . . . . . . . . . . . . . . . . 48
4-7 Temperature dependence of the critical stress intensity factor in HDPE 49
4-8 CTOD as a function of temperature in HDPE . . . . . . . . . . . . . 50
4-9 Plastic zone size as a function of temperature in HDPE . . . . . . . . 51
4-10 Unmodified plots of fracture toughness tests for 1" thick HDPE (with side grooves) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52
4-11 Modified plots of fracture toughness tests for 1" thick HDPE (with side grooves) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53
6

4-12 Temperature dependence of the critical stress intensity factor in 1" thick HDPE (with side grooves) . . . . . . . . . . . . . . . . . . . . . 54
4-13 Compression tests on PVC at various temperatures . . . . . . . . . . 55
4-14 Temperature dependence of yield strength in PVC . . . . . . . . . . . 56
4-15 Temperature dependence of elastic modulus in PVC . . . . . . . . . . 57
4-16 Unmodified plots of fracture toughness tests on PVC . . . . . . . . . 58
4-17 Modified plots of fracture toughness tests on PVC . . . . . . . . . . . 59
4-18 Temperature dependence of the critical stress intensity factor in PVC 60
4-19 Compression tests on ULTEM at various temperatures . . . . . . . . 61
4-20 Temperature dependence of yield strength in ULTEM . . . . . . . . . 62
4-21 Temperature dependence of elastic modulus in ULTEM . . . . . . . . 63
4-22 Unmodified plots of fracture foughness tests on ULTEM . . . . . . .
64
4-23 Modified plots of fracture toughness tests on ULTEM . . . . . . . . . 65
4-24 Temperature dependence of the critical stress intensity factor in ULTEM 66
4-25 CTOD as a function of temperature in Ultem . . . . . . . . . . . . . 67
4-26 Plastic zone size as a function of temperature in Ultem . . . . . . . . 68
4-27 Compression tests on polypropylene at various temperatures . . . . . 69
4-28 Temperature dependence of yield strength in Polypropylene . . . . . . 70
4-29 Temperature dependence of elastic modulus in Polypropylene . . . . . 71
4-30 Unmodified plots of fracture toughness tests on Polypropylene . . . .
72
4-31 Modified plots of fracture toughness tests on Polypropylene . . . . . . 73
4-32 Temperature dependence of the critical stress intensity factor in Polypropylene .......... .................................... 74
4-33 Plastic zone size as a function of temperature in Polypropylene . . .
.
75
4-34 CTOD as a function of temperature in Polypropylene . . . . . . . . . 76
5-1 The cavitated region and the brittle crack origin in HDPE . . . . . . 89
5-2 Closeup of the cavitations and the tapped crack front in HDPE . . .
90
5-3 Brittle crack origin and hackle marks in HDPE at lower temperatures 91
7

5-4 A schematic demonstration of (a) kinked and (b) forked crack geome-
5-5 tries and teh associated nomenclature . . . . . . . . . . . . . . . . . .
Shear yielding in HDPE . . . . . . . . . . . . . . . .
92
93
5-6 The intermediate region during the cracks transition from ductile to brittle fracture in HDPE . . . . . . . . . . . . . . . .
5-7 A closeup of the intermediate region in HDPE . .
. .
5-8 Typical brittle fracture surface . . . . . . . . . . . . .
5-9 PVC fracture surface photograph . . . . . . . . . . .
5-10 Stable to unstable crack growth in PVC . . . . . . .
. . . . .
94
. . . . .
95
. . . . .
96
. . . . .
97
. . . . .
98
5-11 A craze in PVC . . . . . . . . . . . . . . . . . . . . .
5-12 The initial crack transition in PVC . . . . . . . . . .
5-13 A closeup of the initial crack transition in PVC .
. .
. . . . .
99
. . . . .
100
. . . . .
101
5-14 The mountainuos region that defines the final stages o f the high speed brittle crack . . . . . . . . . . . . . . . . . . . . . . . 102
5-15 Brittle crack origin in ULTEM . . . . . . . . . . . . .
5-16 Crazing in ULTEM . . . . . . . . . . .
103
104
5-17
5-18
5-19
Transition from hackle marks to wave pattern in HDPE . . . . . . . .
Multiple crazing at tapped crack front in PP . . . . . . . . . . . . . .
Microvoiding in PP . . . . . . . . . . . . . . . . . . . . . . . . . . . .
105
106
107 ft
8

2.1 Material Properties of the Polymers Supplied by their Manufacturers 21
4.1 HDPE's Measured Values . . . . . . . . . . . . . . . . . . . . . . . . 41
4.2 PVC's Measured Values . . . . . . . . . . . . . . . . . . . . . . . . . 41
4.3 Ultem's Measured Values . . . . . . . . . . . . . . . . . . . . . . . . . 42
4.4 Ploypropylene's Measured Values . . . . . . . . . . . . . . . . . . . . 42
9

The mechanical behavior of metals have been studied more extensively than polymers.
On the whole, their behaviour is well understood and modeled. Polymers, on the other hand, have not had the benefit of similar intensive study and thus their behavior and modelling is still lacking consideraly in comparison. This becomes a serious issue when it is considered that they are becoming increasingly more widely used and are replacing their metal counterparts in many structural engineering and biological applications, and the like.
Considering the microstructural differences between metals and polymers, it is not difficult to see that models developed for metals do not necessarily extend to polymers without modifications. As an example, the theories based on the micromechanisms of rate dependence of flow stress and toughenening, namely dislocations, in metals do not explain the behavior in polymers since polymers do not have similar crystalline structures as metals have.
Most of the modeling of polymers currently is based on phenomenological information rather than physical. This is clearly a disadvantage since each model would be very limited in its application. In 1963 Geil wrote:
"Although rapid advances have been and still are being made in our knowledge of the morphology of crystalline polymers, the gaps in our knowledge are at present ... both wide and numerous. In addition, little ... is known
10

of the morphology of noncrystalline polymers."
Although it has been nearly 40 years, much remains to be understood on polymer morphology and behavior. Therefore, polymers need to be studied more throroughly for them to be better understood.
The following studies have been carried out on the subject of polymer brittleness.
Fracture behavior is an essential field of study for structural applications of materials since they have a tendency to fail in a brittle manner. Fracture toughness is studied in this thesis to find its temperature dependence at below room temperature conditions.
This information is not readily available in the literature. When it is available the studies do not provide the desired details nor offer sufficient explanations. Therefore this study was undertaken to gather more needed information to supplement the currently available literature. Some explanantions of the observed behavior have been attempted.
This study was performed on four polymers
" High Density Polyethylene (HDPE)
* Polyvinyl Chloride (PVC)
" ULTEM (a ployetheremide)
" Polypropylene (PP)
Their known properties and morphologies are briefly explained in Chapter 2. Included for each material are mechanical behavior data and properties obtained from manufacturers and literature. This is followed by a review of cryogenic polymer testing literature and fracture mechnaics principles.
The brittleness and temperature dependence of the critical stress intensity factor for fracture in these polymers was studied by performing compression and fracture toughness tests at temperatures ranging from room temperature down to -1450C.
These tests and the data reductions are explained in Chapter 3. This includes the
ASTM standards and the test equipment.
11

In Chapter 4 the results of the tests are presented. The results are then interpreted and used to calculate other fracture related values such as crack tip opening displacement, plastic zone size, and crack velocity. These values are used to aid the interpretation of the results and the understanding of the behavior of the polymer.
Through this process an explanation of the specific behavior patterns of the polymers is formulated.
Scanning Electron micrographs were obtained of the fracture surfaces. They are presented with explanations in Chapter 5. The features of the surfaces are pointed out and addressed. Also their correlations with the fracture toughness information is delineated to better understand the fracture process they underwent. In the study of the morphologies, issues of crack propagation, ductile to brittle transitions, and adiabatic heating are discussed. Considering local stress intensities and crack tip stress distributions an understanding of crack bifurcation in HDPE is developed.
In Chapter 6 conclusions and suggestions for further work are formulated.
12

Three polymers used in this study were crystalline thermoplastics while Ultem was a glassy polymer. The crystalline properties of certain polymers have been studied in the past but the structure and behavior of semicrystalline polymers is still a topic of much investigation. Though many of the unit cells of crystalline phases are well established the molecular arrangements are not. Crystallizable polymers can be melt-crystalized or grown by precipitation from dilute solutions. Many are known to be spherulitic with amorphous layers arranged between the lamellae rays of the spherulite. The solution grown polymers tend to be lamellar with the chains folded in a regular manner and the layers separated by amorphous regions.
Crystalline polymers are found in a variety of morphological forms ranging from single crystals to semicrystalline polymers in which the non-crystalline component can be either rubbery or glassy. They can also be oriented. Unlike amorphous polymers the fracture behavior of crystalline polymers is greatly affected by structure and morphology and thus their fabrication is an important factor. Moreover, the applica-
13

tion of fracture mechanics is not always straightforward even for isotropic polymers because they are not always brittle and their deformation can be non-linear with large scale yielding in the vicinity of the crack. This however is not a concern at low temperatures because brittleness sets in.
Above the glass transition temperature (Tg) thermoplastic polymers are able to deform extensively in tension or compression. Below Tg the amorphous region is much less compliant and though the material could exhibit a higher yield stress it does not easily flow plasticaly. The variety of mechanical properties that have not always been available makes them worthy of study.
Glassy thermoplastics have been extensively studied for their fracture behavior.
The reasons for this are many. Thermoplastics are relatively simple from a structural viewpoint, compared with either thermosets which have a complex three-dimensional molecular structure or semicrystalline polymers with a great variety of morphological forms. They are also very experimentally well behaved in the sense that the behavior is very consistent because unlike crystalline polymers their material properties are not strongly dependent on their fabrication. Also, since their bulk deformation is approximately linear elastic and the plastic zone is limited to the crack tip, they are well suited for Linear Elastic Fracture Mechanics (LEFM).
A brief listing of some of the properties provided by the manufacturers of the polymers used in this study is given in Table 2.1.
HDPE (High Density PolyEthylene) falls under the category of olefin polymers. It is also known as linear polyethylene. It is composed of a row of carbon atoms each of which are surrounded by two hydrogen atoms. Its unit cell is orthorhombic. Typically
HDPE is very highly crystalline. This causes it to have a relatively high melting point, moderate stiffness, tensile strength, and hardness. Polyethylene is used in injection molding of housewares (eg: bottles) and toys. It is also used in piping and fabrics.
The HDPE used in this study was manufactured by Spartech Plastics (821 Clark
14

Street, Conneaut, OH 44030-1213, 1-800-325-5176)
PVC (PolyVinyl Chloride) falls under the vinyl polymers which used to be the largest group of thermoplastics before the olefins surpassed them. They are only slightly crystalline and the crystals are not large and not well defined. Thus, they are primarily amorphous but tend to retain a discrete particle structure. PVC is generally unstable.
These two facts lead to the complication that PVC is rarely in pure form as other polymers.
The PVC used in this study was manufactured by Geon Company (One Geon
Center, Avon Lake, Ohio 44012, 1-800-438-4366). The sheets were grey in appearance due to the manufacturer injecting coloring during the compounding process. Geon claims this has no affect on its properties and behavior.
Polytherimide (PEI) is an amorphous engineering thermoplastic characterized by high heat resistance, high strength and modulus, excellent electrical properties that remain stable over a wide range of temperatures and frequencies, and excellent processibility.
It derives its trademark name, Ultem, from the General Electric Co. which is the resin producer.
The Ultem used in this study was manufactured by AL Hyde (1 Main st., Grenloch,
NJ 08032; (856) 227-0500).
Polypropylene, like HDPE, is linear and has high degrees of crystallinity which imparts to it high stiffness and tensile strength. It is used primarily in filaments (eg: seat covers) and fibers. It is also used for articles such as luggage made by injection
15

molding.
Spartech Plastics, the supplier of HDPE was also the source of polypropylene used in this study.
Some of the relevant properties of the four polymers used in this study are listed in Table 2.1
Fracture toughness as used in fracture mechanics is a material property in that it is independent of test method and geometry. However it can be temperature dependent and because of polymer viscoelasticity it can be time-dependent too. The characterizing parameter when considering fracture energy is Jc. In the case of linear visco-elasticity, as is the case with polymers, it is more convenient to express this parameter as G, or its equivalent Kc. The 'I' refers to mode I fracture which is the only fracture mode considered in this thesis. Linear Elastic Fracture Mechanics can be utilized for polymers at cryogenic temperatures because yielding is minimal causing brittle fracture.
Glass transitions are important to the study of polymers because they mark significant changes in the behavior. Several of the studies Kinloch and Young referred to in their book Fracture Behavior of Polymers [19] have found some influence of secondary glass transitions on Gmc. However, their dependence on loading rates and material fabrication makes them difficult to study for universal application since each material batch will have a unique glass transition that is affected by the rate of deformation and testing environment. This requires referring behavior to well designed standard reference states. The fracture behavior of polymers above room temperature, where many primary glass transitions occur in industrially interesting polymers, has been studied extensiely. Cryogenic fracture behavior of polymers has not been well investigated and this includes secondary glass transitions. Below is a helpful summary of the literature avialable on cryogenic fracture testing of polymers.
Marshall et. al. [24] showed that in PMMA, KI, increased with decreasing tem-
16 i

perature. Atkins et. al. [11] showed that GI, had a similar temperature dependence.
They also investigated the dependence of these parameters on crack velocity. They found that a critical crack tip opening criterion can be applied over the entire range of tests because it remained constant at ~ 1.6pm [24]. Prior to these studies Marshall et.
al. [25] investigated an added complication to cryogenic testing. They found that the testing medium affects the polymer behavior. The PMMA exhibited greater craze growth rates in active environments. The active environment used was methanol.
Parish and Brown [27] demonstrated that this was the case for liquid nitrogen and liquid helium environments too. They showed that the results and glass transitions presented by Johnston and Beardmore [13], and Beardmore [12] (who did not always specify the environment) on PS, PC, and PMMA were environmentally affected.
Mai and Atkins [22] tested the dependence of Kc on low temperatures and crack velocity in PS while Parvin and Williams [28] did similar studies on PC. Fraser and
Ward [18] showed that a constant critical crack growth criterion could also be applied to PVC as was applied to PMMA by Morgan and Ward [26] a year earlier.
Sims [30], using a trouser leg tear test, studied the dependence of GI, on temperature in PP for temperatures between -60'C and room temperature. The relationship between GI, and temperature was found to be positive. However, this relationship applied to tests below the glass transition. He also found that at temperatures approaching the glass transition (- - 15'C) there was a rapid increase in fracture energy but he did not show the temperature dependence beyond this point. Mai and Williams
[23] showed that the plane strain KIc dependence on temperature in PP and Nylon 6 is minimal but that the plane stress Kc increased with decreasing temperature. Mai and Williams also suggested a constant critical crack-opening displacement criterion in the region of secondary transitions.
Fernando and Williams [17] demonstrated that K, for single edge notched PP specimens in bending and tension remained relatively constant from -150'C to -75"C and then rose with increasing temperature. They also showed that modified PP had a decreasing K, with decreasing temperature until brittleness at -100 0 C was reached, below which K, remains constant. Williams, [1] also mentions that PP has
17

an essentially constant K, (4.7 MPav/ ii) from -100 C to 201C.
Williams [21 also showed that single edge notched PVC specimens in bending and tension had a very slowly rising negative dependence of K, on temperature down to
-1200C below which a pronounced increase is observed. He explained that this could be due to the condensation of liquid nitrogen. In her thesis, Lee [21] pointed out that unnotched tensile PVC specimens underwent a transition from discontinuous yielding to a brittle failure mode at approximately -60C.
Chan and Williams [161 tested single edge notched PE in bending and found it behaved in the same fashion as the PVC with a similar transition temperature.
However, the PE they used was not a homopolymer.
Finally, Saatkamp and Hartwig
used compact tension specimens with chevron shaped crack fronts to study crack propagation in HDPE specimens at -196 0
C. They found a strong increase in the fracture energy associated with uncontrolled crack propagation which they explained by local crack tip heating effects that raised the temperature to above the glass transition causing higher crack resistance from plasticity.
There are many factors to be considered in the study of polymer fracture behavior.
As has been explained, the glass transition is very specific to the material batch, test rate and environment. KIc has similar dependences and is influenced by the glass transition. The dependence of KIc on temperature is a topic of controversy.
Some of the explanations given above for abrupt changes in this dependence were environmental effects, adiabatic heeating effects, and glass transitions.
From the theory of fracture mechanics, a quantity called the stress intensity factor,
K, can be defined that characterizes the severity of the crack situation as affected by crack size, stress, and geometry. In defining K, the material is assumed to behave in a linear-elastic manner, according to Hooke's Law, so that the approach used is called
Linear Elastic Fracture Mechanics.
18

A given material can resist a crack without brittle fracture occuring as long as this K is below a critical value K, which is a property of the material called the fracture toughness. Values of K, vary widely for different materials and are affected
by temperature and loading rate, and secondarily by the thickness of the member and the environment.
Considering a crack in the center of a wide plate:
K = O-ovi (2.1) where
G-o = far field stress a = half the crack length
This equation only applies if a is much smaller than half the width of the plate.
Therefore, the critical far-field stress is:
-coG = c (2.2)
Hence, longer cracks have severe effects on the materials strength.
For Linear Elastic Fracture Mechanics to apply the plastic zone of the crack tip must be small compared to the crack length and other geometric dimensions.
In general terms, K characterizes the magnitude (intensity) of the stresses in the vicinity of an ideally sharp crack tip in a linear-elastic isotropic material. This region is labeled the K-field. The variation of stresses around the crack tip in the K-field are given in Sect. 5.2.4. Their general form is:
=
K
1 o- 2- = fij (6) where
-ij = the various stress components
K
1
= K in mode I (tensilemode) r = radial distance from crack tip fij (0) = angle dependent function of the K field
19
(2.3) i

Fracture toughness testing of polymers in ambient conditions cannot usually be considered LEFM. Under cryogenic conditions, however, LEFM usually applies to fracture testing of polymers.
20

Table 2.1: Material Properties of the Polymers Supplied by their Manufacturers
HDPE PVC ULTEM PP
Melt Index (g/10 min)
Density (g/cm
3
)
0.31 -
0.949 1.42
-
1.27
0.06
0.9
Tensile Strength (MPa) 26.06 50.3 104.8 34.5
Tough to Brittle Transition Temperature
Heat Deflection Temp. (A 66 psi)
< -76oC
720C 82
2150C -28.8-OOC
0
C 2100C 98.9-104.4"C
21

Compact tension (CT) specimens were used to test the fracture toughness of the various polymers. To find the temperature dependence of the fracture toughness, the
CT specimens were tested at temperatures varrying from -10'C down to -145
0
C.
The loading pin displacement rate was 2 mm/min. on a screw driven Instron machine
(Model 5582). The chamber used to maintain temperatures was also supplied by
Instron (Model 3009). Liquid nitrogen was the cooling medium in the chamber.
An available COD gage was initially used to measure the crack opening displacement but was soon abandoned because it had neither the required extension range nor was it usable at the low temperatures. The machine crosshead displacement was used instead of the COD measurements. This proved to be sufficient for the purpose
(see Fig. 3-1).
The polymers were ordered as 10 mm thick sheets from which the CT specimens were machined. The dimensions had to comply with ASTM standards (ASTM Standard no. E399 and D5045 for polymers) as shown in Fig. 3-2 with B as thickness and W as width. The crack length, a, was about 2.7 cm (Table 3.1) such that a/W~ 0.53.
22

A size criterion that must be satisfied to achieve plane strain conditions is:
B, a, (W a) > 2.5(KQ/o-) 2 (3.1) where KQ is the trial Kmc value and a-, is the yield stress for the specific temperature and loading rate. The yield stresses for the different polymers at different temperatures are given in Sect. 3.3.
Once the above criteria (Eq. 3.1) are satisifed plane strain and limited plasticity in the ligaments are ensured and Kc, for the testing conditions, is established. The energy release rate GI, can then be obtained from :
=
(1 v
E
2 )K2 (3.2) but for polymers, E must be obtained at the same time and temperature conditions as the fracture test because of viscoelastic effects. Many uncertainties are introduced
by this procedure and it is considered preferable to determine GI, directly from the energy derived from integration of the load versus displacement curve up to the same load point used for Krc as stated in Sect. 6.2 for future work.
Compression tests on 7 mm tall cylinders with 9.6 mm diameters were carried out at a displacement rate of 2 mm/min. (same rate as the fracture toughness tests) to establish the yield strength for each polymer at the chosen temperatures. This was needed for evaluating the fracture toughness at the various temperatures. An apparatus (Fig. 3-3) was designed to achieve compression from an Instron used in the tension mode. The fixtures resembled a pair of links. As the displacement between the Instron's crossheads increased the polymer, sandwiched between the fixtures, was compressed.
23

The machine and setup stiffness (Fig. 3-4) were determined and used to correct the data collected. True stress and true strain were calculated from load-displacement data using the sequence of equations below: r = initial radius (4.8 mm)
10 = initial length (7 mm)
1 = current length
A
0
= initial cross sectional area
A = current cross sectional area
Al= length increment
F = load
Eeng
= engineering strain
= true strain a
= true stress
6 eng =
E= ln(1 + Eeng)
Ao = 7rr 2
A
=
Aoe'
F
- = -(3.7)
A
The 0.2% offset yield strength was also extracted from these plots by translating the elastic portion of the slope 0.002 strain units to the right and taking the value of stress at the point of intersection betwen it and the curve.
One issue that arose during testing was lubrication. To achieve this 0.05 mm thick
Teflon sheets with WD-40 lubricant were used at room temperature. This form of lubrication fails at low temperatures. Most other lubricating materials and agents also fail at such low temperatures. Therefore, the cryogenic compression tests were conducted without a lubricant. The test apparatus surfaces that were in contact with
(3.3)
(3.4)
(3.5)
(3.6)
24

the specimens were polished to a 0.3pm mirror finish in order to improve slipping between the specimens and the fixtures.
3.4 Stress Intensity Factor Calculations
The factor KQ, the conditional or trial Kc, must first be established in order to calculate Kjc. To determine KQ, FQ must be evaluated from the load displacement plot. where FQ is the greater of the two:
1. The load at the intersection of the load-displacement curve with a line that has a slope that is 0.05 smaller than that of the elastic loading line.
2. Maximum load withstood by the specimen.
Once FQ is obtained, KQ is calculated using :
FQ
KQ = B
2
(3.8) where (0.2 < x < 0.8) f (2 + x)(0.886 + 4.64x 13.32x
2 + 14.72x
3 5.6x
4 )
(1 X)
(39) where:
FQ = load as determined above
B = specimen thickness (10 mm)
W = specimen width from center of pinholes to edge (2") a = crack length measured as an average of 3 measurements on dif f erent crack f ronts x =a/W
It is then determined whether KQ is consistent with the size and yield strength of the specimen according to Eq. 3.1. If all the critera are satisfied then KQ represents
Kc-
25

1.4
1.2
1
0.81
0
0
0.61
0.41
0.2
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
/
0
0 0.5
-
1 displacement (mm)
1.5
I
1-
COD gage
Instron displacement
2
I
Figure 3-1: Comparing two methods of displacement measurements in a -50
0 C HDPE
Compact Tension Test (COD gage vs. crosshead displacement)
2.5
26

W=5.14 cm a=2.7 cm
0.275W=1.4 cm
1.2W=6.1 cm
2.79 cni
O
4.19 cm
0.25W=1.27 cm
B=1.27 cm
Figure 3-2: The dimensions of the compact tension specimens
27 i

Figure 3-3: The compression test setup inside the temperature chamber (the screws were not tightened so that the PVC placed inside can be seen so as to understand the placement of the specimens)
28

400
C- 300
unmodified data using the compression settup
- materials real behavior
_ _ _ _ _ _ _ _ _-__ _
U)
,200
U)
100
-
0
C 0.1
0.2
0.3 0.4
true strain (mm/mm)
0.5
0.6
-
0.7
400
--- modified data using the compression settup
- materials real behavior
D- 300
CO)
CD,
,200
/
/
100
O'
0 0.1
0.2
0.3 0.4
true strain (mm/mm)
0.5
0.6
Figure 3-4: Modifying the data from the ness of the setup (the above plot is the modified data) compression setup by factoring-in the stiffunmodified data and the lower plot is the
-
0.7
29

The polymers were tested in compression to first establish their temperature dependent material properties. The temperature dependence of fracture toughness for these polymers was determined from compact tension tests. Fracture toughness calculations
(Eq.'s 3.8 and 3.9) depend on specimen geometry and loading behavior but in order to ensure plane strain and small scale yielding the size criteria (Eq. 3.1) must be satisfied. Flow stress is needed to verify the Kc test's validity using the size criteria.
The compression tests supplied the flow stress parameter and its variance with temperature. Thus, the temperature dependence of the polymer's fracture toughness is determined. This is presented in the following sections with discussions.
The compression specimens were tested to establish the material constants of specific batches and their temperature dependence. Figure 4-1 shows the stress strain plots of the compression tests at the various temperatures. The properties extracted from these are listed in Table 4.1 alongside other values. As can be clearly seen from the plots, the HDPE undergoes a brittleness transition between -70
0
C and -110'C. The glass transition temperature of the amorphous component of the material, as stated
30

by the manufacturer (Table 2.1), is below -76
0
C. In most of the literature it ranges from -73
0
C to -118
0
C for different grades of PE having different molecular weights.
This transition is evident in the material's inability to withstand much deformation as can be seen in Fig. 4-1. At -110'C and -1451C the specimens fracture just beyond the yield point at about 0.16 units of true strain. The slight dip in the stress strain curve of the HDPE at -70
0
C is due to a circumferential crack that developed but did not become unstable. This may be the result of environmental cracking. Parrish and Brown [14] have shown that PE is affected by liquid nitrogen at increasingly low temperatures. They observed that PE in liquid nitrogen fractured at lower stresses than in a helium environment or vacuum. The N
2 appeared to induce cracks in the samples. The reduced surface free-energy due to the absorption of the N
2 was given as the cause of the brittle behavior. The compression tests presented in Fig. 4-1 could be exhibiting environementally induced brittleness but it cannot be verified unless compared to tests in other mediums.
The temperature dependence of the yield point exhibits a very linear trend in Fig.
4-2 as does stiffness in Fig. 4-3. The yield point increases about 4.3 times from -10
0
C down to -145'C. Elastic modulus increases about 1.4 times. Brittleness is associated with increased stiffness and strength and thus the lower temperatures should yield lower Kic's, however, this is not the case.
The load-displacement plots of the compact tension test results for the various temperatures are presented in Fig. 4-4. A peripheral issue that should be addressed before proceeding is the stiffnees of the system. The fracture toughness test fixtures lacked adequate stiffness due to many connectors (screw connectors and pinholes).
This is apparent in the initial non-linearity of the unmodified load-displacement plots in Fig. 4-4. They were adjusted by making displacement corrections for system compliance and loading-pin penetration into the samples. The corrections were made
by discarding the initial non-linear portion of the curves, extending the linear portion down, and transposing to zero displacement. The adjusted plots are shown in Fig.
4-5.
The behavior is generally linear with very precipitous drops at the point of frac-
31 i

ture. The specimens failed by unstable brittle crack propagation. In the lower temperatures the crack bifurcated and developed into two unstable cracks. These two unstable cracks caused the specimen to separate into three parts as shown in Fig. 4-6.
The angle between the two new crack fronts was roughly 570 t 50 and was relatively unaffected by the temperature. This behavior has not been reported on by most researchers of HDPE who performed fracture tests in cryogenic conditions. However, several of these researchers resorted to the use of compact tension specimens with chevron shaped crack fronts to avoid the cracks bifurcation. At the higher temperatures cracks did not branch. For example at -10'C there were no signs of bifurcation and the material fractured along the crack plane. The load-displacement curve shows some previous local yielding prior to fracture and the fracture surface had a distinguishable milky craze-like region in the initial stages of the crack. The test at -30
0
C failed in a similar manner to that at -10'C but had evidence of bifurcation. However, neither of the two branched cracks became unstable. The corresponding curve also shows some non-linearity prior to fracture.
The fracture toughness calculated from Eq. 3.8 exhibited a linear rise with decreasing temperature (Fig. 4-7) from -10'C down to -70
0
C. This rise has been documented by other researchers and it coincides with the crack's tendency to bifurcate. It rises from about 4.5 MPavjmi, at -10"C, to about 9 MPaj , at -70
0
C.
However, below -70
0
C, this linearly rising trend is curtailed. The fracture toughness ceases to be temperature dependent and settles down to about 7 MPa/-T. This abrupt reduction in the critical stress intensity occurs below the low temperature brittleness transition. The temperature dependence of Kc has been a matter of controversy. Hartwig 161 mentioned that some authors such as Parrish and Brown
[271 expected that Kc is influenced by secondary glass transitions. Hartwig stated:
"Most properties of amorphous polymers are influenced by several weak glass transitions well below the primary glass transition temperatures.
They arise from unfreezing of the mobilities of specific molecular groups
.... Glass transitions hav been found down to 30K. However, for most polymers they do not occur below 120K."[7]
32

Chan and Williams [15] showed that Kc rose with decreasing temperature. They aslo observed a similar glass transition in the same regime. They found that the glass transition temperature was absent from the PE copolymer.
The rise in Kc is unusual because lower temperatures generally cause materials to be increasingly brittle and become defect sensitive. Considering the Crack Tip Opening
Displacement (CTOD, 6) some insight can be given into this negative temperature dependence. Thus,
S 8 a n[sec( )]
7rE 2a-, which for small stresses can be simplified to give:
(4.1)
6 =
__(
UOE
v 2 ) for plane strain (4.2) where
6 Crack tip opening displacement o= yield strength
E elastic modulus o = far field stress a = crack length
K
1
= stress intensity factor
Inserting the values for the range of temperatures from -10
0
C to -70
0
C the
CTOD remains constant at about 400 pm. Figure 4-8 shows this constancy is maintained for the range of temperatures above the glass transition in the amorphous component. The constancy of the CTOD has been verified in several polymers over a range of crack velocities and temperatures and has been termed critical crack opening displacement or a crack tip opening criterion. Marshal et. al. [24] found a 1.6 pm critical CTOD for PMMA over a range of temperatures while Parvin and Williams
33

[28J showed that the critical 6 for PC is approximately 37 pm for a lower range of temperatures. Though the calculation of 6 is questionable because it depends on three temperature dependent parameters, Morgan and Ward [26] measured this constancy in PMMA. Fraser and Ward [181 repeated that work for PC with the same conclusion. Given the constancy of the critical CTOD, the plastic resistance must be rising at a high rate (Fig. 4-2 and Table 4.1) to compensate for the rise in the square of the stress intensity factor (overlooking E because it does not rise much as seen in Fig. 4-3 and Table 4.1). The increased yield strength is giving the brittle material the ability to withstand defects through elastic means while the local crack tip region is undergoing plasticity at higher flow stresses. Below the glass transition this no longer continues because the material remains elastic and a fracture condition is reached before local yielding. This also applies to the rising trends in PVC and polypropylene as is discussed below. Ultem however is simply brittle, with no evidence of plasticity and thus it follows the downward fracture toughness trend that is expected of brittle materials. The temperature dependence of the fracture toughness of Ultem is understood by the occurrence of crazing that accompanies the fracture as explained in Sect. 5.4.
The tendency to bifurcation of cracks in HDPE may be related to the crack velocity.
However, a more plausible explanation is given in Sect. 5.2.4. Crack velocities are an important factor that have been studied in many like investigations. Crack velocities affect Kmc and glass transitions. Another important outcome of the crack velocity is heating which is discussed in Sect. 5.2.2 in relation to fracture surface features.
Although crack velocity is primarily governed by loading rate, temperature and environmental influences are considerable. Since the loading rate was maintained at
2mm/min for all the tests the variation in crack velocity is a result of the temperature and environmental effects. Vincent and Gotham [32] were among the first to show that GI, rises with crack velocity. Marshall et. al. demonstrated a similar relationship between Kc and i for various temperatures.
34

The crack velocity, 6, can be given by: t t (4.3) where t is the time elapsed and rp can be either the plastic zone (as expressed in
Sect. 5.2) or the Dugdale zone for materials that craze. Either zone is described by the same equation:
(4.4)
0 with the value of the constant c depending on which zone size is considered. Using c = y from Eq. 5.1 in Sect. 5.2, the plastic zone size is seen to be diminishing rapidly with decreasing temperature as illustrated in Fig. 4-9. However, this does not mean that the crack velocity is diminishing also since t ,the time elapsed, is diminishing at a greater rate. The plastic zone size at -50
0
C was more than 85% of the plastic zone size at -70
0
C while the elapsed time for fracture at -50
0 C was less than 75% that of -70
0
C. This can be seen more clearly if Eq.s 4.2, 4.3, and 4.4 are combined with
Eq. 5.1 to give:
K c (60)1/2
8o n Eo?" (4.5) where E0 is from a power-law dependence, o = EoEot-, and n, related to molecular relaxation, is a constant in the order of .1. The CTOD has been shown in Sect.
4.1.1 to remain relatively costant. Yield strain is effectively constant. Therefore all the values in the parenthesis are essentially of no bearing to the relationship between
K, and e. Since Kc practically doubles while E grows less than 50%, & must be increasing.
With increasing crack velocities the maximum stress shifts to either side of the plane of the crack causing a tendency towards forking. However, this usually happens in the final stages of crack propagation, preceded by a significant amount of stable crack growth, and where the velocities are approaching sonic levels. On the other hand, the cryogenic conditions under which these cracks occur cause extremely high
35

crack velocities at the early stages of crack propagation. Some cryogenic fracture testing on polymers measured crack speeds up to 1/3 the speed of sound [8] [3].
Other investigations
have shown there is a transitional crack speed that defines a ductile to brittle transition and that this transitional crack speed decreases with decreasing temperature. The higher crack velocities at the lower temperatures could cause this bifurcating behavior however this is not a sufficient explanation since this does not occur in the other polymers that experienced similar high crack velocities.
A more cogent explanation is given in Sect. 5.2.4. Though these stress distributions existed in the other polymers their fracture surfaces showed evidence of crazing and initial plasticity. When HDPE showed signs of cavitation the specimens did not fracture but in the lower temperatures, with no signs of cavitation, the specimens bifurcated. Therefore the energy absorbing mechanisms of deformation caused the cracks to remain in plane while their absence allowed the brittle cracks to deviate.
The measured Kc values were clearly affected by crack forking. In order to confirm the validity of the fracture toughness values, more tests were conducted with thicker specimens. The 1 inch thick specimens (as opposed to the prior 0.5 inch specimens) were machined with a groove along the two sides of the specimens in order to confine the travel path of the crack so as to not allow it to bifurcate. The thickness across the grooved area, which is the reduced specimen thickness, was 0.5 inch. The unmodified and modified plots are presented in Fig. 4-10 and 4-11. The temperature dependence of the critical stress intensity factor in these specimens is commensurate to the thinner specimens thus validating them. The bifurcation does not have any significant effect on the measure of the fracture toughness. Figure 4-12 not only demonstrates the similarity in Kic's, it also demonstrates that there is an abrupt decrease in fracture toughness below -70
0
C re-confirming the britteleness transition in the samples without side grooves.
36

The material properties of PVC are presented in Table 4.2. The manufacturer reports that the glass transition temperature is about 70 0
C while the majority of the literature usually reports it to be about about 80
0
C. The tests however show that there is another brittleness transition at lower temperatures. HDPE exhibited such a transition when the specimens in the compression tests began to fracture at small plastic strains as explained above in Sect. 4.1. PVC showed the transition by a visible and quantifiable increase in the modulus of elasticity. This is apparent in Fig. 4-13 when comparing the linear slopes of the -110"C and -145
0
C tests to the remaining tests.
Fig. 4-15 emphasizes this sudden change in stiffness. The elastic modulus remains relatively constant at 1.1 GPA from -10
0
C down to -70
0
C with a sudden rise to
1.5 GPa at -110 0 C below which it remains constant once again. The yield strength, however, has a linearly rising trend with decreasing temperature (Fig. 4-14). Below
-700C the yield point seems to rise parabolically, though not enough data has been collected to confirm this behavior. The yield point increased roughly 1.85 times from
77 MPa at -10
0
C to 219 MPa at -145
0
C.
This unforeseen low temperature transition also arose in the fracture toughness tests (Fig. 4-18). There was a linearly rising relationship between Kc and temperature down to -70
0
C below which there is an abrupt drop in value from about 7
MPaV m to about 4 MPaV 7i. The unmodified curves in Fig. 4-16 and the modified curves in Fig. 4-17 (for an explanation of the modifications see Sect. 4.1) show evidence of plasticity in the tests. The -10
0
C tests proved to be invalid Kc tests due to excessive plastic deformation. This is obvious from the curves which also show stable crack growth. All the tests leading down to -70
0
C had some degree of plasticity which is confirmed by investigating more closely the linearity of the curves. This, however, was not the case for -110'C and -145
0 C, again emphasizing a transitional behavior.
The cracks often wandered away from the expected crack plane. They did not always result in a clear fracture of the specimens in two pieces at the higher temper-
37

atures since the fractures did not become completely unstable. When the specimens were then finally torn apart at very high rates in order to examine the fracture surfaces, the cracks bifurcated near the end of the specimens for high crack velocity reasons elaborated upon in Sect. 4.1.2.
4.3 Ultem
Unlike HDPE and PVC, Ultem showed no transitional behavior. The compression tests presented in Fig. 4-19 show slowly and uniformly rising yield strengths and stiffness. This is better seen in Fig. 4-20 and Fig. 4-21. The yield point rises with decreasing temperature in a linear manner from 120 MPa at -10
0
C to 185 MPa at
-145'C. The elastic modulus rose at a slower rate, from 1.2 GPa to about 1.4 GPa.
The glass transition temperature of Ultem is around 200 0 C as stated by both in the literature and by the manufacturer. Table 4.3 contains the experimentally verified material properties of Ultem.
Fig. 4-19 shows the stress strain plots at the different temperatures. The 201C test showed an anomolous dip at about 0.175. This dip is due to poorly centering the specimen in the compression chamber. This resulted in an offcentered radial plastic expansion and contact with the wall of the compression chamber. The constraint caused a lateral force shifting the specimen over and an associated load drop.
Ultem behaved in a very brittle manner. The curves in Fig. 4-23 are linear with precipitous drops illustrating the instability of the cracks that signify brittle behavior.
Ultem had a relatively temperature independent Kc with no transitions throughout the range of temperatures from -145
0
C to -10
0
C, given in Fig. 4-24. The Kc values decreased slightly from 6.5 MPaVx/ at --10C to 5.5 MPaIm at -145'C. The result of the -110
0
C test is attributed to experimental error.
The brittleness of Ultem can aslo be seen by its smaller plastic zone size and
CTOD shown in Fig.'s 4-26 and 4-25. The CTOD is essentially constant throughout the range of temperatures.
The fracture surface of the Ultem CT specimens exhibited a wave-like pattern of
38

lines. The waves bowed away from the crack front and had smaller separations in the initial stages of crack growth. As the crack propagated the spacing between the lines diminished. These waves have the appearance of hackle and rib marks. Generally, hackle and rib marks appear when cracks depart from their plane of fracture and are usually relatively steep ridges but Ultem had a fairly flat fracture surface and the ridges were not steep. These features are presented in Sect. 5.4.
4.4 Polypropylene
The temperature dependent material properties of Polypropylene are listed in Table
4.4. Fig. 4-27 show the stress strain curves of the compression tests at different temperatures. Much like Ultem, there is no apparent change in behavior over the temperature range of -10'C to -145'C. The literature and the manufacturer place the T of the amorphous component of polypropylene between 0
0
C and -25'C. This is in agreement with the curves in Fig. 4-27. As explained in Sect.'s 4.1 and 4.2 the brittleness transition correlates with a sudden rise in the stiffness and reduced strain to fracture. Below -10
0
C the compression tests show a gradually increasing stiffness with decreasing temperature. The highest Kc test temperature was -30 0
C and thus no brittleness transition could be seen by the fracture toughness tests.
The 0.2% offset yield as a function of temperature is plotted in Fig. 4-28 as is the elastic modulus in Fig. 4-29. The temperature dependence of the yield strength follows the pattern found in PVC (Fig. 4-14). It is initially linear and then becomes slightly parabolic. However, the rise is greater than that of PVC. The yield strength at -145
0
C (120 MPa) is 4 times larger than the yield strength at -10'C (30 MPa).
The temperature dependence of the elastic modulus increases at a decreasing rate and is an almost linear relationship rising from 0.7 GPa to 1.2 GPa for the same temperature range.
As with PVC, polypropylene showed significant plasticity effects at the higher temperatures in fracture toughness tests as seen in Fig. 4-30 and Fig. 4-31. The temperature dependence of KIc (Fig. 4-32) also showed a very slight upward trend
39

as HDPE and PVC, but without a drop at low temperatures. This trend, as with the stiffness trend, showed minimal temperature dependence. Fig. 4-33 shows the calculated plastic zone size as a function of temperature. It follows the trend exhibited by the other polymers. However, the CTOD shows to be slight temperature dependence
(Fig. 4-34).
The fracture surfaces of the tested CT specimens were identical for all the test temperatures. The cracks did not deviate from the expected crack plane path and the fracture surfaces were smooth. At the higher temperatures, the crack did not become unstable, yet, the surface maintained the same appearance as the fracture surface of the specimens that underwent unstable crack growth at the lower temperatures. The only observable difference between the specimens was a shear lip along the side of the specimens in the higher temperature experiments which is a result of loss of the triaxility of stresses at the surface.
40

(OC)
-10
-10
-30
-30
-50
-50
-70
-70
-70
-70
-110
-145
Temperature Crack
Length
Table 4.1: HD
Yield
PE's Measured Values
Elastic Critical Stress
Strength Modulus Intensity Factor
(a, cm)
2.7457
2.772
2.754
2.6387
2.69
2.661
2.63
2.63
2.752
2.685
2.725
2.705
(UO, MPa)
47.36
47.36
72.47
72.47
88.32
88.32
118.4
118.4
118.4
118.4
164.71
203.86
(E, GPa) (K4c MPav6 )
0.94 4.63
0.94
1.1
1.1
4.73
5.93
1.3
1.3
1.48
9.96
6.73
7.62
11.02
1.48
1.48
1.48
1.41
1.56
8.81
8.25
9.59
6.89
7.13
(OC)
-10
-10
-30
-30
-50
-50
-70
-70
-70
-70
-110
-145
Temperature Crack
Length
(a, cm)
2.735
2.6483
2.7103
2.6967
2.625
2.74
2.595
2.5925
2.7917
2.748
2.68
2.762
Table 4.2: PVC's Measured Values
Yield Elastic Critical Stress
Strength Modulus Intensity Factor
(0-0,
MPa) (E, GPa) (K,,, MPav/'m)
76.89
76.89
1.08
1.08
6.46
5.64
90.34
90.34
1.11
1.11
3.21
3.58
105.77
105.77
113.95
113.95
113.95
113.95
151.64
218.96
1.16
1.16
1
1
1
1
1.48
1.48
4.93
5.05
8.02
8
5.28
5.16
4.1
3.83
41

(OC)
-10
-10
-30
-30
-50
-50
-70
-70
-70
-70
-110
-145
Temperature
(a, cm)
2.624
2.6343
2.7167
2.74
2.6
2.6
2.54
2.54
2.756
2.756
2.64
2.681
Crack
Length
Table 4.3: Ultem's Measured Values
Yield Elastic Critical Stress
Strength Modulus Intensity Factor
(E, GPa) (KC, MPa
/m-) (o, MPa)
122.78
122.78
1.18
1.18
6.5
6.56
124.68
124.68
137.41
137.41
158.36
158.36
158.36
158.36
171.28
186.5
1.22
1.22
1.28
1.28
1.2
1.2
1.2
1.2
1.44
1.31
5.76
5.69
6.46
5.42
6.04
6.35
5.92
5.46
4.27
5.64
(OC)
-10
-10
-30
-30
-50
-50
-70
-70
-70
-70
-110
-145
Temperature
Table 4.4: Ploypropylene's Measured Values
Crack Yield Elastic Critical Stress
Length
(a, cm)
2.7303
2.7287
3.008
2.743
2.717
2.762
2.64
2.64
2.862
2.777
Strength
(o, MPa)
30.83
30.83
37.21
37.21
45.04
45.04
52.54
52.54
52.54
52.54
Modulus
(E, GPa)
0.73
0.73
0.89
0.89
0.91
0.91
1.01
1.01
1.01
1.01
Intensity Factor
(KIC, MPav/-j)
1.97
2
3.45
2.04
3.93
3.58
3.76
4.27
3.59
3.61
2.78
2.757
84.16
119.98
1.09
1.17
3.97
4.04
42

250
200-
'-i 150 -
1V cn
100-
50+-4
I I o 20
0
C x -10 OC
+ -300C
-704C
-1100C
-1450C
H
0 0.1 0.2 0.3 true strain (mm/mm)
0.4
Figure 4-1: Compression tests on HDPE at various temperatures
0.5 0.6
43

250
I I I I
200-
'2 150 a..
C',
CD,
100 -
0
0
50- 0
0 1
-160 -140 -120
1
-100
1
-80 -60 temperature (OC)
-40 -20
Figure 4-2: Temperature dependence of yield strength in HDPE
0 20
44

2
1.8 --
1.6-
1.4-
I
0
I I
01.2-
0
-n
E
-
CO 0.8-
0.6-
0.4-
0.8-
0
0.2-
0
-160 -140 -120 -100 -80 -60 temperature ( 0 C)
-40 -20
Figure 4-3: Temperature dependence of elastic modulus in HDPE
0 20
45

2500
2000-
1500 -700
-
-
-100C
300C
-30'C
-
_
5000
~~-50'C
-70 OC
700C
-700
-
-70'C
-11 00C
4500
1000-
500-
-5000
0
0 0.5 1 displacement (m)
1.5
Figure 4-4: Unmodified plots of fracture toughness tests on HDPE
2 2.5
x 10-3
46

2500
2000
15001
0
10001
500
/
/
/1
I
I>
-
S-
100C
-300C
-300C
-500C
-
-7000
-700C
700C
~ -70'C
-1100C
-1400
0
0 0.5
1 1.5
displacement (m)
2
Figure 4-5: Modified plots of fracture toughness tests on HDPE
2.5
x 10-3
3
47

Figure 4-6: The bifurcated cracks in HDPE
48

15
S
E
(U
10-
U)
C0
0
U,
0)
0
CD
0 glass transition
0
0
-160 -140 -120 -100 -80 -60 temperature (4C)
-40 -20
Figure 4-7: Temperature dependence of the critical stress intensity factor in HDPE
0 20
49

1
E
0.6o 0.5-
0
0.4-
0.9-
0.8-
0.7-
0.3-
0.2-
0.1
1 I 1 1 1 1
-160 -140 -120 -100 -80 temperature (0C)
-60 -40
Figure 4-8: CTOD as a function of temperature in HDPE
1
-20 0
50

1 .5
E a)
N
1-
N
I I
0
0
0
-160 -140 -120 -100 -80 temperature (0C)
-60 -40
Figure 4-9: Plastic zone size as a function of temperature in HDPE
-20 0
51

2500 --
2000-
1500-
0
1000-
500 -
0
/
0.5
-
-100C
..........400C
-400C
700C
-- -700C
1
- 1100C displacement (m)
1.5 ---1400C
-- 400
Figure 4-10: Unmodified plots of fracture toughness tests for 1" thick HDPE (with side grooves)
X
2.5
10-3
52

2500
I I I I
2000-
1500-
1000-
500-
0
0 0.2
//-40
0.4 0.6
~ - 1
-400C
0
-700C
-700C
~~-11 00C
--- 110 C
0.8
~-1 400C
1 -- 140*C 11.4 displacement (m)
1.6
Figure 4-11: Modified plots of fracture toughness tests for I" thick HDPE (with side grooves)
1.8
X 10-
2
53

15
E
101-
CO
(D>
C
U)
0
5 x x x x x x x
0 -
-150 -100 -50 temperature (0C)
Figure 4-12: Temperature dependence
HDPE (with side grooves) of the critical stress intensity factor in 1" thick
0
54

300
250-
200-
U)150 --
Uw
100-
50
I
C7
0 200C
X
10 c x--100C
-500C
-70'C
-110C
-1 450c
0 0.1 0.2 0.3 true strain (mm/mm)
0.4
Figure 4-13: Compression tests on PVC at various temperatures
0.5 0.6
55

250 --- -
200-
CD
CO)
100-
0)0
0
0
50-
0 1
-160 -140 -120 -100 -80 -60 temperature (0C)
-40 -20
Figure 4-14: Temperature dependence of yield strength in PVC
0 20
56

2
1.8-
1.6-
1.4-
0
1.2-
E
) 0.8-
0.6-
0.4-
0.2-
0
-160 -140 -120 -100 -80 -60 temperature (0C)
-40 -20
Figure 4-15: Temperature dependence of elastic modulus in PVC
0 20
57

2200
2000-
1800-
1600
1400
1200
2-
0 1000
-/
800
-'
600
400 i
200
0
0 1
I
2
I
3
I
4 5 displacement (m)
6
I
Figure 4-16: Unmodified plots of fracture toughness tests on PVC
7
-
-----
--
31000
-300C
300C
-500C
500C
-
-700C
-700C
-700C
-700C
-11 000
0
14500
8 x 10-3
9
58

2200
2000
1800 I
16001
1400
-Z 1200
0 1000
-
II
800
600
400
2001
-(
-/
I1OI
-
-~ -1-
/-300C
-300C
~~ -500C
-50*C
- -700C
-700C
-
'700C
-70OC
-1100C
145G-
0
0 1 2 3 4 displacement (m)
5 6
Figure 4-17: Modified plots of fracture toughness tests on PVC
7 x 10-3
8
59

10
7-
E a. 6--
CO,
5-
0
C,
2-
3-
9-
8-
2-
1
0
2epeaur
I I
(C
-160 -140 -120 -100 -80 -60 temperature (00)
-40 -20
Figure 4-18: Temperature dependence of the critical stress intensity factor in PVC
0 20
60

300
250-
200a-
150-
I
100-
50 -
0
0 0.1 0.2 0.3 true strain (mm/mm)
5 200C
X -1 04C
+ -30'C
-500C..
A -700C
V -110oC
-145
0
C
1
0.4
Figure 4-19: Compression tests on ULTEM at various temperatures
1
0.5 0.6
61

250
200a.
150-
,5100--
0 1
50--
00
0 0
-160 -140 -120 -100 -80 -60 temperature (4C)
-40 -20
Figure 4-20: Temperature dependence of yield strength in ULTEM
0 20
62

2
1.8-
1.6
1.4-
0
E u0.8-
0.6
0.4-
0.2-
0
-160 -140 -120 -100 -80 -60 temperature (0C)
-40 -20
Figure 4-21: Temperature dependence of elastic modulus in ULTEM
0 20
63

1600
1 4001
12001
10001
U)
0
0
8001
600
-.
-
-
-
-10
0
C
-O 0
4C
-300C
-30
0
C
-500C
~- 500C
-700C
-700C
-700C
-1100C
14500 / 7
/
"A
/
~
/*
,-~
400
/ /
200
0
0 0.2 0.4 0.6 0.8
1 displacement (m)
1.2 1.4
Figure 4-22: Unmodified plots of fracture foughness tests on ULTEM
1.6 1.8
x 10-3
-
64

1500
10001
/
/
~~/7
/ ///
/
4'
/
"
~
/
/
/
/
1 j*j i~
~
H
>NI
II ii
I
K
I--,
0
7
500
0
0
/
0.2
/
//
/
//
0.4
1/
I
0.6
7
I
0.8 1 displacement (in)
1.2
-
Figure 4-23: Modified plots of fracture toughness tests on ULTEM
-L;"
-100C
-30*C
-300
-0C
-500C
-504C
-700C
-700C
-700C
_700C
-11 000
1-4500 1.8 x 10-3
2
65

1 5 I I------
C 10-
E
0z
CL to
0)
C
5
0
00
20 50
0
-160 -140 -120 -100 -80 -60 temperature (SC)
-40 -20 0
Figure 4-24: Temperature dependence of the critical stress intensity factor in ULTEM
20
66

1
0.6-
E o
0
.5
0.4-
0.3-
0.2-
0.9-
0.8
0.7-
I I I 1 1
0.1 --
00
-160 -140 -120 -100 -80 temperature (SC)
-60 -40
Figure 4-25: CTOD as a function of temperature in Ultem
-20
0
0
67

0.5
0.45-
0.4-
0.35-
E
E 0.3-
N
0
N0
0.25 -
Cu)0.2-0
0.15-
0.1 -
0.05 --
0
-160 -140 -120 -100 -80 temperature (4C)
-60 -40
Figure 4-26: Plastic zone size as a function of temperature in Ultem
-20
-
0
68

160
140-
120-
100c 0-)-
60--
40--
20 --
0 20"C
-10C
+-300C
-500C
S-700C
V-11 00C
45 S0.1 0.2 0.3 true strain (mm/mm)
0.4
Figure 4-27: Compression tests on polypropylene at various temperatures
.6
69

C')
CO)
U)
0
U)
100-
50-
00
-160 -140 -120 -100 -80 -60 temperature (0C)
-40 -20
Figure 4-28: Temperature dependence of yield strength in Polypropylene
0 20
70

2
1.8-
1.6-
1.4a- 1.2-
0
0
0
0.6-
~0.2-8
0.4-
0.2-
I I I
-160 -140 -120 -100 -80 -60 temperature (OC)
-40 -20
Figure 4-29: Temperature dependence of elastic modulus in Polypropylene
0 20
71

1000-1
900-
800
700-
10*C
-300C
-300C
-
-- -500C
500C
-
- -700C
6700C
-700C
110 C
1
600\
Q 500
2
400
300
-
/-
-
200-~
100
0 0.5 1 1.5 2 displacement (m)
2.5
Figure 4-30: Unmodified plots of fracture toughness tests on Polypropylene
3
-
3.5
X 10-3
72

1100
1000-
900
800
700
600 -/-
3 500 -
400--
300 -
200
100
0
0
/
/..
9000 r
\
I-
-
-1
0
0 C
-300C
-300C
500C
~50'C
~~~-704C
-700C
-700C
-700C
-1100C
14
0.5 1 1.5 2 displacement (m)
2.5 3
Figure 4-31: Modified plots of fracture toughness tests on Polypropylene
-
-
3.5
X 10-3
4
73

10
E
7-
Cz
CO,
0
63--
2-
3-
2-
9-
8-
1
I I
-160
474
-140 -120 -100 -80 -60 temperature (00)
-40 -20
Figure 4-32: Temperature dependence of the critical stress intensity factor in
Polypropylene
0 20
74

1
0.9 _
0.8-
0.7-
E 0.6 -
0
N
(0.5-
C)
0.3-
0.2-
0
0
-160 -140 -120 -100 -80 temperature (0C)
-60 -40
Figure 4-33: Plastic zone size as a function of temperature in Polypropylene
-20 0
75

0.6-
E o 0.5-
0.4 --
1
0.9-
0.8-
0.7-
0
0.1 -
0
0
-160 -140 -120 -100 -80 temperature (CC)
-60 -40
Figure 4-34: CTOD as a function of temperature in Polypropylene
-20 0
76

5.1 Specimen Preparation
A LEO 438VP environmental scanning electron microscope was used to study the fracture surfaces. The specimens were machined to fit in the SEM and the sputter coater. 200
A of gold/palladium was vapor deposited on the fracture surfaces of the samples enhancing the image contrast and eliminating charge buildup on the surface. The samples were pressed down on copper tape and mounted on metal studs.
Silver paint was along the sides and edges of the samples improved conduction from the gold/palladium coating to the stud draining the charge build-up. A computer interface was used to produce digital images at several locations on the fracture surfaces.
5.2 HDPE
5.2.1
The fracture surfaces of HDPE had three well defined regions very similar to that found from impact loading notched standard HDPE. The first was a small milky crazelike region that was more readily observed in the higher temperature tests where crack bifurcation did not occur. The second was a 'mountainous' region that had ridges
77

extending out from a central point delineating an origin of crack growth. The third was a very smooth usually curved surface.
The first region was of particular interest beacuse it was only observed clearly in the samples where the crack did not bifurcate. Under closer inspection the topography of the region resembled indpendent sites of cavitation (Fig. 5-1 and Fig. 5-2), a process that usually accompanies ductile fracture. Most cavities had centeral foci that indicated sources. The cavities were smaller in size and greater in number near the crack front of the razor tapped crack yet greater in size and less in number away from the crack front. Therefore, the earlier cavities were nucleation controlled as opposed to the latter ones being growth controlled. This is due to the greater tensile stresses experienced by those nearer to the crack front. The boundaries of these cavities were bowed out towards the crack. In ductile fracture, cavities occur just beyond the crack front and grow towards the crack causing the crack to travel forward. If the plastic zone size is considered, an interesting correlation is found.
Using the Irwin approximation: r, =
37
( a,,
)2
(5.1) where r= plastic zone ahead of the crack
K
1
= stress intensity factor o= yield stress
A rough estimate of the plastic zone size can be calculated. Using the values obtained at -10 0 C (K
1
= 4.7 MPaV-\Y and -o =- 47 MPa) the plastic zone size is found to be 1.05 mm. The cavitated region measures about 1.1 mm (Fig. 5-1).
Here the initial form of fracture was by ductile cavitation. At -30
0
C, the plastic zone size is approximately 700 Mm. Since the cavities are, on average, 200pam in diameter the plastic zone size would have about 2-3 cavities. Figure 5-3 shows this to be true.
There is virtually no evidence of cavitation at -70
0
C or below.
A fracture surface feature that appeared at the lower temperatures involved lo-
78

calized shear. This can be seen in Fig. 5-5 which is a fractograph of the crack front on the specimen tested at -50
0
C.
Between the first and second regions is a very short intermediate stage (Fig. 5-6) that appears to be the result of a temperature rise-induced softening due to plastic cavitation prior to the unstable crack growth. This transitional softening is better seen in Fig. 5-7. Saatkamp and Hartwig [291 demonstrated that adiabatic heating increased crack tip plasticity in fracture tests using chevron shaped CT specimens.
They concluded that the increased level of plastic flow raised the fracture toughness.
Kausch [4] also noted this in impact loaded notched HDPE. Local heating at a crack tip can be induced by friction, chain scission, or other dissipative processes. This effect is especially strong at low temperatures where the specific heat is small, and even low heat pulses raise the temperature drastically. Adiabatic conditions exist at unstable crack propagation where the rate of heat generation is lower than for its removal by thermal conductivity. Since the crack undergoes a significant velocity increase between the region of ductile stable crack growth and unstable brittle crack growth, adiabatic heating would have occured. Hartwig [9] showed that the ratio between thermal relaxation time and heat generaion time is 10 500, meaning that fully adiabatic conditions exist for unstable crack propagation for polymers in low temperatures.
Marshall et. al. [241 showed that the abrupt rise in crack velocity during the transition from stable crack gowth to unstable crack growth is due to an adiabatic/isothermal transition. Below the transitional crack velocity an isothermal state of heat dissipation accompanies the stable crack growth whereas above the transitional crack velocity there is adiabatic heating at the crack tip and the material in that region is thermally softened.
79

Marshall et. al. [241 showed that adiabatic conditions can occur at relatively low crack velocities because of relatively low thermal conductivities using: c (ATad)
(60
2 PCk
K;e) 2 where
= unstable crack velocity
ATad = adiabatic temperature rise T To
T = temperature at the crack tip
T= test temperature
(5.2) p = density c specific heat k = thermal conductivity
Kj*c= value of K
1 e at the instability
Saatkamp and Hartwig [29] argued that the temperature rise in front of the crack tip can be estimated using the fracture energy as the upper limit for the heat source.
About 60% of GI, is consumed in the plastic zone and thus converted into heat (from studies by Weichert and Schonert
Under these assumptions one obtains, for the adiabatic case:
A\Tad
<
0.6 Gjc pcrp
(5.3)
This reasoning showed that ATad can range from 50 K to 100 K. This temperature rise can develop added yielding at the crack tip. In some cases the crack tip can be experiencing above glass transition temperatures while the bulk is below glass transition [19] [29] [9].
80

The SEM micrograph in Fig. 5-1 shows characteristic river markings of brittle fracture surfaces. Figure 5-8 is a sketch of the region examined on the SEM image. The fracture surface is divided into four regions:
1. The flaw is the defect that acts as the cracks source. The fracture origin is normally due to machining, impact, or material defects such as pores or inclusions.
The material defects can be inherent defects that exist naturally in a material or are introduced in processing.
2. The mirror region is an area with a smooth, glossy appearance that surrounds the initiation region. This region marks the beginnings of unstable crack propagation. The crack velocity is relatively slow in this region but it has begun to accelerate.
3. The mist region has been described as having a 'matte appearance'. This region is substantially rougher than the mirror region because the crack travels faster creating parabolic markings giving it the 'matte appearance'.
4. The hackle region appears to be the roughest area due to highest crack velocity.
Macroscopic crack branching initiates at the end of the hackle region.
This sequence of processes is observed in Fig. 5-3
Cavity formation is driven by the hydrostatic tensile components of stress and this intrinsically causes the crack to propagate in a plane perpendicular to the maximum tensile stress. This typically results in slow cavity expansions followed by crack extension through the cavitated regions. At the higher temperatures where there was a well developed plastic zone and cavitation occurred, the crack did not deviate from its expected plane. However, at the lower temperatures the crack was more affected
by shear yielding which occurs at a 450 ± 8' angle as is explained in Sect. 5.2.4 below.
81

Local Stress Intensities on Bifurcating Cracks
Clarification of the stress distribution around the crack tip is important to understand crack behavior. Suresh [10] showed that crack bifurcation and kinking can be understood by the measure of the angle of deviation and local stress intensity factors:
K
1
= a
11
(oz)K
1
+ a
12
(a)KII
K
2
= a
2 1
(OZ)Kj + a
22
(a)Kr-
(5.4)
(5.5) where K, and KI, denote the mode I and mode II stress intensity factors for the main crack in the absence of any bifurcation and a is the angle between the two new crack fronts as demonstrated in Fig. 5-4. K
1 and K
2 are the local stress intensity factors for the bifurcated cracks. They are, respectively, transverse and along the path of bifurcated crack propagation. To a first approximation in a, the dimensionless factors for a kinked crack are:
I a al (a) = (3cos + cos
3a\
3 /8, a12(a)
-
4 a si + sin
32\
1 / a .3a\ a21(ce) = stn + sin all (o) = (cos + 3cos 2a
With only a mode I applied to the crack in these experiments Eq.s 5.4 and 5.5 simplify to:
(5.6)
(5.7)
(5.8)
(5.9)
K
1
= a
11
(c)KI
K
2
= a
21
(a)Kr
(5.10)
(5.11)
The tapped cracks are hardly, if ever, in the same plane as the machined crack.
This is due to the tipping of the blade during the tapping process. The tipping angle
82

depends on the geometry of the specimen (specifically the machined groove) and the blade dimensions. In these experiments the blade tapping produced a crack at a 50 angle away from the crack plane. Applying equations 5.4 and 5.5 to this crack:
al ~ 0.997 a
2
1
~ 0.044
The angle subscribed by the two crack fronts is initially 30±50. As the cracks grow apart the angle increases to 90' with the average being about 57" ± 5'. With similar derivations for finite kinked cracks, Suresh and Shih [31] showed that k
2 vanishes at 2a = 32' whereas k, reaches a maximum while still remaining smaller than K
1
.
Therefore, the phenomena of crack forking improves the fracture toughness. This improvement, however, is not significant as demonstrated in 4.1.3
83

Crack Tip Stress Distributions
Although the initial 50 angle caused by the blade tapping did not have a significant affect on the macroscopic stress intensity it does alter the stress distribution around the crack tip. Considering :
0 = angle from crack plane r = distance from crack tip
For Mode I
o-r =
0 =
-rO =
K
1
5 0 1 301 cos-
-
-cos-
V2 2 4 2
K
1
3 0 1 301
cos- + -cos -
/2irr 4 2 4 2
I
K
1
1 . 0 1 . 30
-sn- + -smn-
/2irr 4 2 4 2
For Mode II
-rr = rO =
K
2
V/2 ir
[ 5 . 0
-- sin-
4 2
3 . 0
-- sin-
4 2
-
3 301
+ -si-I
4 2
3
.301
4 2
K
2 r
1 0 3
.301
-cos- + -sin-
4 2 4 2
(5.12)
(5.13)
(5.14)
(5.15)
(5.16)
(5.17)
84

Superposition gives:
K
1
-os-
0 -
-s-+
V/2 4 2 4 2
K2
/2Fr
5 .
42
3 30]
4 2
K
1
UOs =
00
27rr
K
01-ro =4sn-
1
3 0 1 30~
-Cos- + -Cos-
4 2 4 2
K
2
3 .
3 .30
sin - - -sinr 4 2 4 2
1 0 1. 30 K
2
+ -sin- +-Cos--
2 2 4 2 /2
1 0 3 30
+ -sin--
4 2 4 2 l with
K
1
= 0.997Kr
K
2
= 0.044K
1 from Sect. 5.2.4.
Differentiating to get the angles of maximum tensile and shear stress:
(5.18)
(5.19)
(5.20) d 3 0 1 301 5 .0 3 301
0 = d (0.997 [4 cos2 + 4 cos 2 + 0.44 [-4sn2 + 4 sin 2
0 = d
0 1.31
0.997 -sin- + -sin + 0.044 cos + 3sin
(5.21)
(5.22)
The maximum tensile stress is alligned 50 away from the crack plane while the maximum shear stress is 430 away from the crack plane. This places the plane of maximum shear stress along the same plane as the shear bands.
Local stress intensities and stress distributions explain the negative temperature dependence of fracture toughness but fail to explain the cause of crack bifurcation.
Considering the other polymers did not experience bifurcation the difference must lie in the difference of structures between HDPE and the remaining polymers or the unquantifiable environmental effects that could affect HDPE and not the other polymers. HDPE is significantly more crystalline than PVC or PP and Parrish and Brown
114] showed that PE becomes markedly crack sensitive in liquid nitrogen environments at low temperatures.
"Changes in crack path are generally induced by such factors as multiaxial far-field
85

streses, interaction of the crack tip with microstructural inhomogeneities, abrupt load excursions, or the embrittling effect of an aggressive environment."[31] Therefore the crack may have bifurcated due crystalographic structure and environmental effects.
The local stress intensties coupled with the crack's velocity result in the higher K
1 c's with continued bifurcation.
Much like HDPE, PVC had 3 well defined fracture surface regions that varried in size with temperature (Fig. 5-9). The PVC fracture surfaces had a clear correspondence between the behavior observed on the fracture toughness plots and the fracture surface topography of the specimens.
The first region was a dull flat surface that was featureless to the naked eye. This region marks the stages of ductile stable crack growth defined by the linear portion and of the curve and the onset of non-linearity. As the temperature was reduced this region became less visible. At -145oC the region is absent and the corresponding curve in Fig. 4-17 exhibits minimal displacement.
The second stress whitened region (from grey to milky grey) is defined by ridges that resemble ribs. The surface was rougher than the previous one yet still flat. This is evidence of a transition to brittle fracture. The plot bounds this region between the point the curve reaches the maximum load and the point of a markedly precipitous drop which indicates that the crack has transitioned from stable ductile fracture to unstable brittle fracture. The scanning electron fractograph (Fig. 5-10) shows this transition with a trench that spans the width of the speciman at the crack front (Fig.
5-11). It resembles the crazes reported by Lee in her study of deformation mechanisms in PVC.
86

Dowling f5] used a transition crack length, at, as an approximate crack length above which the strength is expected to be limited by brittle fracture.
at =(5.23)
When comparing this calculated transition crack length to the surface measurement it is found that the transitional crack length was half of the measured length at best. Since Dowling is using this transition crack length as an industrial criterion for the employment of fracture mechanics a factor of safety of two three is probably implicitly used in Eq. 5.23. If this is true the results are reasonable.
Another interesting characteristic is the region immediately following the transition (Fig. 5-12). Many voids accompany this region which is better seen in Fig.
5-13.
The final stage of crack growth is marked by the sudden drop in the load-displacement curve which marks complete fracture. The region returns the color of the PVC from the milky-gray appearance of the former region. The SEM image in Fig. 5-14 shows it as an extremely 'mountainous' region. This is the region in which the crack may deviate from its plane due to increased crack velocities. At the low temperatures the entire surface was composed of this region only less 'mountainous'.
5.4 Ultem
As alluded to in Sect. 4.3 the surface of the fractured Ultem CT specimens primarily consisted of a wave-like pattern that bowed out away from the crack. These waves could be the result of the interference of the crack front with the elastic waves released by the fracture itself. The waves become less frequent and have greater wavelength along the crack's path of propagation. This is the outcome of a crack that is accelerating.
The fracture surface is dominated by hackle marks originating from the origin of
87

the brittle crack propagation (Fig. 5-15). The flaw, mirror region, and hackle marks are distinguishable. A prior form of crazing can be seen in Fig.s 5-15 and 5-16. The lower fracture toughness of Ultem at the lower temperatures is attributed to smaller craze regions since craze initiation is an energy absorbing process.
Fig. 5-17 illustrates the transition from the hackle marks to a wavelike pattern.
5.5 Polypropylene
The macroscopic fracture surface of polypropylene was deceptively brittle in appearance. However, the polypropylene surface contained much microvoiding. Microscopically the failure was ductile on a small scale with the surface consisting of a series of cusps with highly deformed fibrils between. The cusps form from the microvoids which appear to nucleate from within the material. Fig. 5-18 shows a fibrillated form of void growth at the blade tapped crack front. Voiding was prevalent throughout the entire surface (Fig. 5-19).
Unlike Ultem, polypropylene cavitated more readily at the lower temperatures.
The active liquid nitrogen environment has been known to cause materials to tend to craze. The larger craze regions at the lower temperatures absorb more enrgy causing the fracture toughness to rise. Another factor to be considered is adiabatic heating. Adiabatic heating, as discussed above in Sect. 5.2.2, is a function of crack velocity. At the lower temperatures the crack is travelling at greater velocities causing increased adiabatic heating. The heating could very well have raised the crack tip temperature above the glass transition since the glass transition is well within the testing temperatures. The material would thus exhibit more ductility at the crack tip.
88

Figure 5-1: The cavitated region and the brittle crack origin in HDPE
89

I - i FeEw . -, -I . - -- ---
Figure 5-2: Closeup of the cavitations and the tapped crack front in HDPE
90

.- I -- - --
- =-- - -- .
-F-F4- I - - I I , -m
Figure 5-3: Brittle crack origin and hackle marks in HDPE at lower temperatures
91

do
-.-
I, rp
'p
I
~-fl
>
I
/
/
Figure 5-4: A schematic demonstration of (a) kinked and (b) forked crack geometries and the associated nomenclature
92

. I --- , - - - I . -- -- --- - aw - - 2
I
Figure 5-5: Shear yielding in HDPE
93

DATA 001170 1 1 0 1 1024 768
Figure 5-6: The intermediate region during the cracks transition from ductile to brittle fracture in HDPE
94

Figure 5-7: A closeup of the intermediate region in HDPE
95

Flaw
Mist Region
Mirror Region
Hackle Region
Figure 5-8: Typical brittle fracture surface
96

lk goIV
U:
Figure 5-9: PVC fracture surface photograph
97

Figure 5-10: Stable to unstable crack growth in PVC
98

Figure 5-11: A craze in PVC
99

Figure 5-12: The initial crack transition in PVC
100

Figure 5-13: A closeup of the initial crack transition in PVC
101

i ............
Figure 5-14: The mountainuos region that defines the final stages of the high speed brittle crack
102

Figure 5-15: Brittle crack origin in ULTEM
103

Figure 5-16: Crazing in ULTEM
104

Figure 5-17: Transition from hackle marks to wave pattern in HDPE
105

Figure 5-18: Multiple crazing at tapped crack front in PP
106

OEM
19119REArn k ''I --- , ,, -
Figure 5-19: Microvoiding in PP
107

Fracture toughness of crystalline thermoplastics generally rises with decreasing temperature. This fracture toughness is diminished and ceases to be temperature dependent below the glass transition for polymers like HDPE. Some polymers, like PVC, with above room temperature glass transitions may still exhibit this transition at secondary glas transitions below room temperature. While others yet, do not have a brittleness transition like PP. Glassy polymers like Ultem, however, have a positive temperature dependence.
The higher KI,'s at lower temperatures is the result of many factores. Higher yield stresses with constant CTOD predicts this negative dependence. Though less crazing and cavitation, which are energy absorbing processes may be apparent at the lower temperatures, higher crack velocities accompanied by adiabatic heating could increase the energy absorbing capacity of the plastic zone.
Crazing, cavitation, and shear yielding may be present in cryogenic fracture testing. Failure by crazing or cavitation occurs at lower Kin's than failure by shear yielding. Crazing is the dominant mechanism of fracture in glassy polymers. Temperature affects the mechanisms of deformation. More importantly environmental factors have an unestablished effect on them such as causing HDPE to fracture in compression tests where it would not otherwise.
108

In cryogenic fracture testing crack velocities are of such magnitude that crack tip heating is significant. This adiabatic heating can cause the temperature in the crack tip plastic zone to rise above the glass transition inducing greater ductility.
Bifurcation of the crack is introduced by a slight variance in the crack plane
(usually introduced by initiating through blade tapping) in the absence of energy absorbing plastic zone micromechanims of deformation. However this is not a sufficient reason. Crystal structure and environment are probably the major causes of crack branching.
6.2 Suggestions
More tests are required at a variety of temperatures to confirm the trends established in this thesis. At least three tests should be conducted at each temperature. Testing temperatures in the glass transition regime would be of interest. More polymers should also be considered in order to increase the available literature on cryogenic testing of polymers.
Thicker specimens should be considered to reduce the degree of plasticity in the higher cryogenic temperatures. In retrospect, chevron shaped specimens would be better suited for testing. Chevron specimens have many added benefits. For example energy release rates can be determined through direct methods because crack arrest will occur. These crack arrests will also allow for utilization of a COD gage from which crack velocities can be better quantified using crack tip opening displacement rates. J-integral concepts should also be utilized.
More fractography can be implemented. Extensive studies of each of the stages of crack propagation can be studied with special attention to transitonal points (eg: point of crack forking and ductile to brittle transitions). Using the chevron CT fracture experiments, unstable crack growth could be inhibited or arrested at desired test points. These points could be studied by cleaving the remainder of the specimen at high rates exposing and marking those regions. Side views of the fracture surfaces would contain much information too. A determination of the spherulite size and
109

crystal structure can be benefitial to understanding the cavitated region in the HDPE and the cause of crack bifurcation. Measures of surface roughness would also aid the understanding of the fracture process.
The effect of the temperature chamber environment should be considered since it is well known that polymer behavior is drastically affected by active environments.
110

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