EFFECT OF DESIGN FACTORS ON THERMAL FATIGUE CRACKING OF DIE CASTING DIES

Final Technical Report

David Schwam John F. Wallace

Sebastian Birceanu

Department of Materials Science Case Western Reserve University

Cleveland, Ohio

Work Performed Under Contract DE-FC07- 00ID138486 US Department of Energy Assistant Secretary for Energy Efficiency and Renewable Energy Washington DC

October 2004

TABLE OF CONTENTS

Content Page

TABLE OF CONTENTS 2

LIST OF TABLES 4

LIST OF FIGURES 5

ACKNOWLEDGEMENTS 8

ABSTRACT 9

1. Introduction 10

1.1. Die Failure Modes 10

1.2. Thermal Fatigue Failure Mechanisms 11

1.3. Thermal Shock and Thermal Fatigue Resistance Evaluation Factors 16

1.4. Thermal and Physical Properties that Affect Thermal Fatigue Resistance 7

1.4.1. Thermal Conductivity 7

1.4.2. Thermal Expansion Coefficient 8

1.5. Mechanical Properties that Affect Thermal Fatigue Resistance 9

1.5.1. Elastic Modulus and Strength 9

1.6. The Effect of Thermal Cycling on Microstructural Stability 10

1.7. The Temperature-Time Effect on the Structure of Martensitic Steel 20

2

Content Page

1.7.1. Martensitic Transformation in Steel - Brief Overview 20

1.7.2. Tempering of Martensite 22

2. Materials and Experimental Procedures 26

2.1. Materials 26

2.2. The Thermal Fatigue Test 27

2.2.1. Specimens and Equipment 27

2.2.2. Thermal Fatigue Cracks Evaluation Procedure 28

2.2.3. Temperature Measurement 29

2.2.4. Microhardness Measurement 30

2.2.5. Scanning Electron Microscopy 30

3. Results and Discussion 39

3.1. Softening During Thermal Cycling and Thermal Fatigue Resistance 39

3.1.1. The Influence of Immersion Time on Softening and Thermal Fatigue dsadasdasCracking 40

3.1.2. The Influence of Cooling Line Diameter on Softening and Thermal sadadadadFatigue Cracking

47

3.2. Stress Analysis at the Specimen Surface and Around the Cooling Line 51

3.3. Microstructure Degradation that Promotes Softening During Thermal sasdsdCycling 52

4. Conclusions 59

5. Bibliography 99

3

LIST OF TABLES

Table Page

2.1. Chemical Composition of Experimental Material - Premium Grade H13 31

2.2. Typical Properties of Premium Grade H13 32

2.3. Characteristics of the Tested Specimens 33

3.1. Measurement Data For Different Immersion Times 44

3.2. Immersion Time Effect on Hardness Variation Across the Surface 45

3.3. Measurement Data For Different Cooling Line Diameters 49

3.4. Cooling Line Diameter Effect on Hardness Variation Across the Surface 50

4

LIST OF FIGURES

Figures Page

1.1. Hysteresis Loop at the Surface of a Material Subjected to Cyclic Heating aaaaand Cooling 15

2.1. CCT Diagram for H13 steel 35

2.2. The Reference Specimen for Thermal Fatigue Test 36

2.3. The Thermal Fatigue Test Equipment 37

2.4. Temperature Measurement 38

3.1. Relationship between Tensile Properties and Hardness 61

3.2. The Effect of Thermal Cycling on Crack Area-Different Immersion Times 63

3.3.The Effect of Thermal Cycling on Crack Length- Different Immersion Times 64

3.4. The Effect of Thermal Cycling on Microhardness Distribution Across the aaaaSurface- Different Immersion Times 65

3.5. The Effect of Temperature on Crack Area-Different Immersion Times 66

3.6. The Effect of Temperature on Crack Length-Different Immersion Times 67

3.7. Effect of Elevated Temperature on Tensile Strength 68

3.8. The Effect of Hardness Recovery on Thermal Fatigue Cracking 70

3.9. Relationship Between Total Crack Area and Average Maximum Crack aaaaLength 71

3.10. The Relationship Between Maximum Crack Length and Microhardness at aaaaaMaximum Crack Length 72

3.11. The Effect of Temperature on Microhardness-Different Immersion Times 73

3.12. The Effect of Microhardness at Average Maximum Crack Length on Crack aaaaaArea 74

5

Figures

Page3.13. The Effect of Microhardness at Average Maximum Crack Length on Crack aaaaaLength-Different Immersion Times 75

3.14. Microhardness Profile at the Corner of 12 Seconds Immersed Specimen 76

3.15. Tempering Curve for H13 77

3.16. Maximum Temperature Cycle for 1.5" Cooling Line Specimen After 12 aaaaaSeconds Immersion Time 78

3.17. The Effect of Thermal Cycling on Crack Area-Different Cooling Line aaaaaDiameters 79

3.18. The Effect of Thermal Cycling on Crack Length-Different Cooling Line aaaaaDiameters 80

3.19. The Effect of Thermal Cycling on Microhardness Distribution Across the aaaaaSurface-Different Cooling Line Diameters

81

3.20. The Effect of Temperature on Crack Area-Different Cooling Line aaaaaDiameters

82

3.21. The Effect of Temperature on Crack Length-Different Cooling Line aaaaaDiameters

83

3.22. The Effect of Temperature on Microhardness-Different Cooling Line aaaaaDiameters

84

3.23. The Effect of Microhardness at Average Maximum Crack Length on Crack aaaaaArea-Different Cooling Line Diameters

86

3.24. The Effect of Microhardness at Average Maximum Crack Length on Crack aaaaaLength-Different Cooling Line Diameters

87

3.25. The Effect of Immersion Time on Temperature

88

3.26.The Effect of Cooling Line Diameter on Temperature

89

3.27. Cracks at the Corner of H13 Specimen

90

3.28. Crack at the Cooling Line of H13 Specimen

90

3.29. Stress Modeling at the Corner and Cooling Line 91

6

Figures

Page

3.30. Effect of Volume Percent Primary Carbides on the Transverse Charpy V-notch Impact Toughness of H13

92

3.31. Microstructure Sampling at the Corner of 12 Seconds Immersion Specimen

93

3.32 a. The Effect of Temperature on Microstructure – Unaffected

94

3.32 b. The Effect of Temperature on Microstructure – 0.2” from Corner

94

3.32 c. The Effect of Temperature on Microstructure – 0.1” from Corner

94

3.32 d. The Effect of Temperature on Microstructure – 0.06” from Corner

94

3.32 e. The Effect of Temperature on Microstructure – Corner

94

3.33.Temperature Influence on Carbide Size and Distribution-Photomontage

95

3.34. Effect of Austenitizing Temperature on the Weight Percentage of Isolated aaaaaCarbide Residues in H13 Steel

96

3.35. Small Carbide in Softened H13

97

3.36. Large Carbide in Softened H13

98

7

ACKNOWLEDGEMENTS

This research investigation was supported by the Department of Energy, Office of

Industrial Technology through the Cast Metal Coalition program. The Die Materials

Committee of the North American Die Casting Association provided guidance for this

work. The efforts of Mr. Steve Udvardy, Director of Research and Education at NADCA

and the members of the committee are gratefully acknowledged.

This publication was prepared with the support of the U.S. Department of Energy

(DOE), Award No. DE-FC07-00ID138486. However, any opinions, findings, conclusions

or recommendations expressed herein are those of the authors and do not necessarily

reflect the views of the DOE.

8

ABSTRACT

The thermal fatigue of steel die casting dies becomes more severe at higher

operating service temperatures, reducing die life significantly. Consequently, to extend

die life, die design has to address efficient cooling methods. A key issue in this respect is

the size and location of cooling lines relative to the surface of the die. This subject was

studied in detail, to elucidate the effect of die temperature on thermal fatigue cracking.

The investigation correlates the thermal fatigue cracking in an immersion test specimen

with the temperature attained near the surface and the corresponding softening of the

steel. The effect of cooling line location vis-à-vis the surface temperature and the

resulting cracking pattern are shown for various immersion times and different sizes of

cooling lines. Higher temperatures induce faster and deeper softening of the steel leading

to more thermal fatigue damage. Die design with cooling lines close to the surface can

reduce this damage significantly.

Since the thermal fatigue test has previously provided a remarkably accurate

prediction of the relative thermal fatigue cracking, these results should have good

applicability in die casting operations.

9

1. INTRODUCTION

1.1. Die Failure Modes

The durability of materials in molten aluminum is an important consideration in

engineering applications such as die casting, containment of liquid metal and semi-solid

processing [1]. Die casting is the process of choice in many manufacturing industries -

automotive, hardware, electrical and electronics, computers and many others. It provides

high volume and cost effective aluminum, zinc and magnesium components with good

properties. Some of the advantages of this technology over the traditional sand castings

are [2, 3]:

- Die casting is able to provide complex shapes within closer tolerances;

- Higher rates of production with little or no machining required;

- The die casting parts are durable, dimensionally stable, and have a good appearance;

- Die castings are monolithic; they combine many functions in one, complex shaped part;

The main failure modes of aluminum die casting dies are physical erosion

(washing), chemical attack (corrosion), gross cracking (cleavage cracking) and thermal

fatigue cracking (heat checking) [2,3,4,5]. Erosion occurs when a swift flow of melt

exists relative to the surface of the solid, and becomes more severe when there are hard

particles in the melt. Chemical corrosion refers to dissolution of materials by the melt as

well as the formation of interphase layers, when the relative motion between the solid

material and the melt is negligible [1].

10

Gross cracking is usually catastrophic and may result in complete cracking

through the die. Massive fracture of die casting dies occurs when the die material is

stressed beyond its fracture strength. This can occur even when the applied stress is

below the yield stress. It results from a combination of thermal and mechanical stresses

[6]. This type of failure is related to the inherent resistance of the die material to fracture

termed “fracture toughness”.

1.2 Thermal Fatigue Failure Mechanisms

The life of dies used at elevated temperatures is often determined by their thermal

fatigue properties [7]. The fatigue failure produced by fluctuating thermal stresses is

known as thermal fatigue. Thermal stresses occur when the expansion or contraction of a

part as a result of a temperature change is restrained [8]. The constraint may be internal

or external [9]. External constraints produce forces that act on a component that is

alternately heated and cooled. Internal constraints may result from temperature gradients

across the section (simply because heat is not able to flow quickly enough in response to

the external changes), structural anisotropy and different coefficients of expansion in

adjacent phases or grains [10]. Thermal fatigue resulting from the action of internal

constraints can also be defined as thermal cycling damage.

Temperature gradients form as a result of heating and cooling of the surface

during injecting the molten metal, ejection and the lubricant spraying stages of the die

casting cycle [4, 11]. When molten aluminum is injected, the die surface heats up

creating a steep temperature gradient between the surface and the cooler underlying mass

11

of the die. As a result, the surface wants to expand more than the interior. Because the

interior is more massive, it prevents the surface layer from expanding. As a result of this

internal restraint, the surface is placed under compression. As heat is conducted into the

underlying layers, the surface temperature decreases rapidly. When the casting is ejected,

the surface of the die cools down. The spray of die lubricant further decreases the surface

temperature [12]. The surface then cools more rapidly than the interior, the compression

stresses are relieved and tensile stresses may be created.

The temperature gradient and the coefficient of the thermal expansion of the

material determine the magnitude of the surface stress. For the simple case of a bar with

fixed end supports, the thermal stress generated by a temperature change ∆T is:

σ = αE∆T (1-1)

where α is the linear thermal coefficient of expansion and E is the elastic modulus [8].

For a biaxial condition, the stress is given by:

σ=αE∆T /(1-ν) (1-2)

where ν is Poisson’s ratio. The thermal stresses never fully develop because α, E, ν and

the thermal conductivity all change with temperature [13].

The fatigue damage of metallic materials can be subdivided into the stages of

crack initiation, subcritical crack propagation, and final unstable fracture [14]. Crack

initiation usually occurs at stress concentration sites originating from component

geometry, machining irregularities or surface imperfections [15]. During the compressive

part of the cycle, the increase in temperature lowers the yield strength of material, and the

12

compressive strain may become plastic when substrate prevents deformation. During the

tension part of the cycle, the concentrated thermal stress is larger than the yield strength

of the material, and reversed plastic deformation may occur. After sufficient number of

cycles, the localized plastic deformation will cause a fatigue crack. Once a crack is

initiated, propagation occurs along a plane perpendicular to the maximum tensile stress.

Fatigue cracks in steels can deviate noticeably from the expected plane of propagation

when large prior austenite grain boundaries are present or when another crack is in close

proximity [15]. The influence of other cracks on crack propagation is explained by the

interaction of the highly stressed regions at the tips of the cracks. When the crack tips are

close enough, this interaction changes the general state of stress. This provides an

attraction of cracks to each other until the cracks are joined by reversed crack branching

or forking. When the side branches join, they produce the classical craze-cracking pattern

[16]. Several opinions exist about the driving force for crack propagation. One opinion is

that cracks widen and deepen by the wedging action of the molten metal or oxide that is

forced into them in each shot [17]. Another is that crack propagation occurs only by the

thermal fatigue mechanism. In this event, propagation has to occur during the cooling

cycle, since that provides the tensile stress necessary for crack propagation. Once a

fatigue crack is formed, it will continue to grow because of the stress concentration effect

even when the tensile stress is low [18].

The damage caused by thermal cycling can be separated into stages [14]:

- crack initiation at the surface;

- crack linking at the surface;

13

- growing of small cracks in depth direction from the crack net;

- growing of the largest crack to the complete failure.

The generation and evolution of thermal stress have been explained based on the

type of thermal transients that occur during the service of a part subjected to thermal

shock-thermal fatigue [10]. Suppose a thick structure at low temperature is suddenly

brought in contact with a hot fluid. As explained before, the surface tends to expand

against the remaining material and goes into compression, yielding along OQ (Figure

1.1). Because of the heat transfer towards the core, the temperature gradient decreases

and the system expands, taking the surface into tension at R. The residual tension is

responsible for intergranular cracking. When the material is further subjected to rapid

cooling this series of events is reversed. The surface now goes into tension, as it tends to

contract, with a tension peak at point S that promotes transgranular cracking since the

corresponding strain rate is high and the temperature decreases. Later, when the whole

structure cools, the yielded section at A goes into compression. The compression stress at

P (at the original strain level) is not particularly damaging, but when many cycles are

repeated, the system goes into a hysteresis loop at the surface (PQRS), leading to thermal

fatigue cracking due to the reversed deformation.

1.3. Thermal Shock and Thermal Fatigue Resistance Evaluation Factors

A common measure of thermal shock resistance is the maximum sudden increase

in the surface temperature that a material can sustain without cracking [19]. The thermal

shock resistance and thermal fatigue resistance depend on a number of material

14

properties. These include the thermal expansion coefficient α, thermal conductivity k,

thermal diffusivity K, elastic modulus E, fracture toughness KIc, tensile (fracture) strength

σf and upon the additional parameters of heat transfer coefficient h, specimen size and

duration of thermal shock [4,19,20,21].

Figure 1.1. Hysteresis Loop at the Surface of a Material Subjected to Cyclic

Heating and Cooling [10].

15

A commonly used thermal shock - thermal fatigue resistance parameter is the

merit index of R=σf/Eα or R'=kσf/Eα [4,19,21] . For large values of Biot number

(β = bh/k), i.e. large heat transfer coefficient h, radius or thickness r or b, respectively,

and small thermal conductivity k, or when thermal strains are the result of the material

being mechanically constrained, the thermal shock resistance is determined by R. For

very small Biot numbers, i.e. small heat-transfer coefficients, small radius or thickness,

and large conductivity, or when thermal strains are a consequence of thermal gradients

resulting from rapid heating or cooling, the thermal shock resistance is determined by R'.

If we consider the thermal fatigue as a series of repeated thermal shocks, these

parameters can be used to describe the thermal fatigue resistance and for ranking of

materials.

In this respect, the effect of elements incorporated into the chemistry of an alloy

should be considered based on their contribution to [22]:

a) Thermal properties of the material: coefficient of thermal expansion, specific heat, and

thermal conductivity

b) Material strengthening through carbide formation, solid solution strengthening with

consequent increase in the capacity for withstanding repeated strains and with improved

creep performance.

c) Microstructural stability and oxidation resistance

1.4. Thermal and Physical Properties that Affect Thermal Fatigue Resistance

1.4.1 Thermal Conductivity

16

The thermal conductivity is the quantity of heat transmitted, due to unit

temperature gradient, in unit time under steady conditions in the direction of the

temperature gradient. This condition occurs when the heat transfer is dependent only on

the temperature gradient [23]. Under the conditions described above, thermal

conductivity will reduce thermal fatigue by establishing a low thermal gradient between

the surface and the underlying layer. Equations (1-1) and (1-2) indicate that a lower

temperature gradient will decrease the stress in the material. The successful use of a

molybdenum insert in pressure die casting dies partly results from the high values of

thermal conductivity. However, because of its variation with temperature, the influence

of this parameter may be diminished by the operating conditions. For instance, ferritic

steels have generally higher values of thermal conductivity than austenitic steels, but at

high temperature, say at 1073 K, their thermal conductivities become similar [22].

1.4.2. Thermal Expansion Coefficient

The coefficient of linear thermal expansion is the ratio of the change in length per

degree K to the length at 273 K. The coefficient of volume expansion is about three times

the linear coefficient [23]. The combination of the temperature gradient and the

coefficient of thermal expansion determine the magnitude of stress, as shown by the

equation (1-1). Indeed, the amount of expansion in the axial direction of a slit from a

blade-divided subjected to a temperature Tx will be, according to Duhamel's analogy [20]:

ε = αTx (1-3)

17

and the compressive stress induced by bringing the slit back at its initial dimension, will

be:

σ = -EαTx (1-4).

Among metals, refractory metals have the lowest thermal expansion coefficients [22].

1.5. Mechanical Properties that Affect Thermal Fatigue Resistance

1.5.1. Elastic Modulus and Strength

The elastic modulus is a measure of the stiffness of the material. It is defined as

the ratio of the stress and strain in the elastic regime:

E = σ/ε (1-5)

A lower modulus results in lower stress at a given strain level. Some metallic, but

especially structural ceramic materials are susceptible to failure when thermally shocked

due to a high Young's modulus, combined with relatively high thermal expansion

coefficient, low strength and low thermal conductivity [22].

In general, a material with a low Young's modulus and a high value of yield

strength is desirable, as the elastic component of the strain is large and the plastic

component is small during a typical thermal cycle. The best combination of properties is

18

a high strength-high ductility (high toughness) material, but unfortunately a high strength

is often associated with a low value of ductility.

A very important issue related to thermal fatigue resistance of materials is the hot

hardness and the variation of strength with the temperature. At high temperatures, the

surface loses strength and hardness, especially in steels. This fact will lower the thermal

fatigue resistance. The parameters R and R' will therefore change their values with the

temperature not much because of the variation of the Eα product, which is roughly

constant with increasing temperature [10], but due to the sudden drop in strength at a

certain temperature (which in the case of R'), may not be compensated by the rise in

thermal conductivity.

1.6. The Effect of Thermal Cycling on Microstructural Stability

The prolonged exposure to elevated, varying temperatures and repeated stresses

as it takes place in thermal cycling naturally causes changes in the microstructure [22].

These changes in the metallurgical structure may contribute to failure by reducing

strength and they are referred to as instabilities. Sources of instabilities include

transgranular-intergranular fracture transition, recrystallization, aging or overaging, phase

precipitation or decomposition of carbides. Borides or nitrides, intermetallic phase

precipitation, delayed transformation to equilibrium phase, order-disorder transition,

general oxidation, intergranular corrosion, stress-corrosion cracking, slag-enhanced

corrosion, and contamination by some trace elements also cause instabilities [5].

19

1.7. The Temperature-Time Effect on the Structure of Martensitic Steel

1.7.1. Martensitic Transformation in Steel - Brief Overview

The rapid cooling of a plain-carbon eutectoid steel, after it was heated in the

austenite region, in such a manner that it misses the nose of the TTT diagram curve, will

lead to the formation of the martensite [31]. The conversion of an austenitic

microstructure to a martensitic microstructure in many steels takes place continuously

with decreasing temperature during uninterrupted cooling. This is a unique characteristic

of the transformation kinetics of martensite and is referred to as athermal transformation

[24]. As a general definition, a martensitic transformation occurs by nucleation and

growth and involves the coherent formation of a phase from another without a change in

composition, by a diffusionless and homogeneous lattice shear [32].

Martensite in steels is a metastable body-centered tetragonal (or body-centered

cubic, below 0.2% C) supersaturated solid solution of carbon and other alloying elements

in Fe-α, in which the alloying elements remain locked into the position they occupied in

the parent austenite [31,33].

From a morphological point of view, Fe-C martensites can be classified into two

types:

- lath martensite, typical of all low and medium carbon with up to 0.6 %C;

- plate martensite, above about 1.0% C; its formation was found to be favored by

austenite stabilizers, such as N, Ni, Pt or Mn, but prevented by ferrite stabilizers like Si,

20

Cr, W, V, and Mo. Between 0.6 and 1 % C, a mixture of lath and plate martensite occurs

[31,34].

Another important issue for the material properties past transformation is the

grain size of the parent phase, austenite. The austenitic grain size will not affect the

number of martensite nuclei in a certain volume, but the plate size is a function of the

grain size. In larger grain size material the strain associated with the transformation can

cause large residual stresses to build between adjacent grains. This can eventually lead to

grain boundary rupture. Fine grains will not be that susceptible to this phenomenon, due

to self-accommodation, and together with a smaller martensitic plate size, will provide

for a stronger and tougher material [33]. The general trends related to the austenitic grain

size in heat treated products are:

- Hardenability - deeper hardening for coarse-grain austenite, and shallower hardening for

fine-grain; the addition of alloying elements, except for cobalt, will minimize the

difference, because of the increase of hardenability and inhibition of the grain growth

[40].

- Toughness - higher for small-grained material

- Distortion, quench cracking, internal stress - less present or prevalent in fine- grained

structure [39].

An important observation is that increasing the austenitizing temperature will

produce an improvement in the thermal fatigue performance as a result of the higher

tempering resistance [40], despite a larger grain size. Large grains were proved to be

21

detrimental to thermal fatigue resistance [42,43]. This effect is probably the result of the

more effective dissolution of alloy carbides and the consequent increase of alloying

elements in solid solution.

1.7.2. Tempering of Martensite

The martensitic transformation is essential for the hardening of steel and induces a

desirable hardness. It also increases brittleness, which results from factors such as lattice

distortion caused by carbon atoms trapped in the octahedral sites, impurity atom

segregation at austenite grain boundaries, carbide formation during quenching, and

residual stresses. The hardness of martensite will increase with carbon content and/or

alloying elements. In order to improve ductility and toughness (and sometimes even

strength), most of the technological steels must be tempered. During the heating for the

tempering process, a number of solid-state reactions may occur [24,30,33,34,35,37,38]:

• 25-100 °C (77-212 °F) Carbon segregation to dislocations and boundaries or pre-

precipitation clustering (in high-carbon steels), caused by the interaction energy created

between carbon and strain field around dislocations. In low carbon-steel Ms temperature

is higher and can be sufficient time for carbon to segregate or even precipitate as ε

carbide or cementite during quenching.

• 100-200 °C (212-392 °F)- First stage of tempering - Precipitation of transition

carbides - η(Fe2C) or ε(Fe2.4C) - in steels with carbon content above 0.2 %. The

phenomenon is accompanied by a slight increase in hardness. Below 0.2 %C, the atoms

22

prefer to diffuse at the boundaries or dislocation sites during cooling. Consequently, not

much carbon is left in solution to precipitate upon reheating.

• 200-350 °C (392-662 °F)- Second stage of tempering - Decomposition of

austenite retained after quenching especially in low-alloy steels with more than 0.4% C,

into ferritic bainite and carbides. It is associated with tempered martensite embrittlement,

since carbides replace the austenite in the spaces between the laths of martensite.

• 250-350 °C (482-662 °F) - Beginning of the third stage of tempering - Lath-like

orthorombic Fe3C precipitation.

• 350-550 °C (662-1022 °F) - Segregation of impurity and alloying elements, which

is responsible for temper embrittlement. The temper embrittlement has been attributed to

the segregation of impurity atoms such as P, Sb, As or Sn to prior austenite grain

boundaries.

• 400-600 °C (752-1112 °F) - Recovery of dislocation structure; Lath-like Fe3C

agglomerates to form spheroidal Fe3C, but the lath structure is maintained. During

recovery, the cell boundaries and random dislocations contained between them are

annihilated and a fine grain acicular structure is developed.

• 500-700 °C (932-1292 °F) - Formation of alloy carbides, also called the fourth

stage of tempering. Occurs in steels containing sufficient carbide forming elements (Ti,

Cr, Mo, V, Nb or W). Above about 500 °C, substitutional diffusion becomes significant

and alloy carbides replace the less stable cementite which dissolves as a finer alloy

carbide dispersion forms. Two ways exist in which cementite-alloy carbide

transformation can take place:

23

- in situ transformation - the alloy carbide nucleates at several points at the

cementite/ferrite interfaces, and grow until cementite disappears and is replaced by a

alloy carbide dispersion

- by separate nucleation and growth - the alloy carbides nucleate heterogeneously within

the ferrite on dislocations, lath boundaries, and prior austenite grain boundaries. The

carbides then grow at the expense of cementite. The stable carbide forming elements like

V and Mo are hence the promoters of the strengthening reaction that occurs in the

temperature range from 500 to 600 °C. This is known as secondary hardening, induced

by the replacement of the coarse cementite by the finer alloy carbide, as described above.

• 600-700 °C (1112-1292 °F) - Recrystallization and grain growth occur. The

ferrite can recrystallize more readily in low rather than high-carbon steels, because the

grain boundary pinning caused by carbide precipitates inhibits the process. After

recrystallization is complete, growths of carbide particles and of ferrite grains are the

only kinetic processes that continue.

One of the major concerns in die steel selection is the softening that occurs due to

the thermal cycle. Steels for aluminum die casting experience a high temperature that

could reach 1200 °F during the casting thermal cycling [4]. It has been shown in previous

studies [4,11] that the thermal fatigue behavior is better for temper resistant steels.

Alloying elements that help retard the rate of softening during tempering are desirable.

24

The most effective elements in this regard are strong carbide formers such as chromium,

molybdenum and vanadium [24]. The decrease in hardness and strength of carbon steels

during tempering is largely due to the coarsening of Fe3C with increasing temperature.

Under these conditions an element with a greater affinity for carbon like those mentioned

would form alloy carbide with high resistance to coarsening and therefore provide

hardness retention, good creep and thermal fatigue resistance. The favorable influence of

these alloying elements can turn into a deleterious one, when present in steels in too high

of a quantity. Excess alloying elements produces large carbide particles on the grain

boundaries in the quenched and tempered steel and increase the brittleness of the steel,

resulting in gross cracking. A high austenitizing temperature can dissolve the carbides in

the solid solution, but too high of a temperature will lead to a grain coarsening with same

detrimental results.

25

2. MATERIALS AND EXPERIMENTAL PROCEDURES

2.1. Materials

The material chosen for this work was the Premium Grade H13 steel, since this is

the preferred die steel for the aluminum die casting industry. The composition of the steel

is given in Table 2.1. H13 is a chromium hot work steel. It is basically a hypoeutectoid

steel with high hardenability and a good combination of strength, hot hardness, toughness

and ductility. It has good resistance to tempering. Some typical physical and mechanical

properties of H13 are presented in Table 2.2. This steel has limited amount of alloy

segregation, a fine grain size and a structure that has a low inclusion content and low

concentration of sulfur and phosphorus.

The following heat treatment procedure was chosen in order to obtain the strength

and toughness combination required by the aluminum die-casting industry. The

specimens were austenitized at 1875 °F, oil-quenched according to the schematic CCT

diagram in Figure 2.1, and then double tempered at 1100 °F for 2 hours. Such a

procedure with double tempering will tend to eliminate the residual austenite, and lead to

a predominantly tempered martensitic structure with a hardness of 44 - 46 Rc, high

strength and good toughness.

26

2.2. THE THERMAL FATIGUE TEST

2.2.1. Specimens and Equipment

Specimens for the thermal fatigue test were processed to the dimensions shown in

Figure 2.2. The reference specimen is 2”x2”x7”, rectangular in shape with a 1.5”

diameter and 6.5” long hole in the center for internal water-cooling. Three other

specimens were designed with 1.6", 1.7", and 1.8" cooling line diameters. The four

corners of the specimens were designed and fabricated with a radius of 0.010” and the

specimens' surface was hand polished with 240, 320, and 400 grit silicon carbide paper.

The thermal fatigue test equipment is shown in Figure 2.3. The specimens were

alternately cycled (dunked) in a molten aluminum alloy (380 grade) bath, which was

maintained at 1350 °F. A pneumatic system consisting of an cylinder automatically

actuated was used to immerse and withdraw the specimens from the aluminum bath at

different cycle durations consisting of 5, 7, 9 (reference) and 12 seconds immersion and

24 seconds withdrawn. Water flowed through the specimens at a rate of 1.5 gal/min

through the internal cooling line shown in Figure 2.2. The outer surface of the specimen

was sprayed with water just before it entered the molten aluminum bath. The specimens

were turned 90° around their long axis every 1,500 cycles to insure the uniform spraying

of the water. Table 2.3 summarizes the specimens used and their particular

characteristics.

27

2.2.2. Thermal Fatigue Cracks Evaluation Procedure

Specimens were removed from the test system after 5,000, 10,000 and 15,000

cycles and their cracks were measured. Since the temperature fluctuations and

geometrical constraints are the greatest at the corners, cracks form mainly at the corners.

For measuring the cracks, the surface of the specimens is polished with 240, 320

and 400-grit silicon carbide paper. A V-shaped fixture with 400-grit silicon carbide paper

is used to polish the corner. Only cracks on the corners within a 3” central length were

measured, to eliminate the end effect of the top and bottom areas.

Two concepts are used to evaluate the thermal fatigue resistance of the steels,

Average Maximum Crack Length and Total Crack Area [4]. The Average Maximum

Crack Length La is the average length of the longest cracks on the four corners, within the

middle three inches of the corners.

∑=

=4

141

imia LL

where i = 1...4 indicates each of the four corners, and Lmi is the maximum crack length of

i corner.

The crack area of each crack is defined as the square of the crack length. The

Total Crack Area is the sum of the products of the number of cracks in each 100 micron

size range and the square of the midpoint of that range for all the four corners.

28

∑∑= =

=4

1

2

1,

ij

n

jjit LNA

where Lj = 100j-50 µm

j = 1...n, corresponds to different crack length range and Ni,j represents the number of the

cracks of i corner in the crack length range of 100(j-1) to 100j µm. The number and

length of all cracks were measured under an optical microscope attached to a Leitz

microhardness tester.

2.2.3 Temperature Measurement

In order to determine the temperature of the corner, a thermocouple hole was

drilled in a specimen with an initial cooling line diameter of 1.5” (Figure 2.4). The

drilling was performed at an angle from the vertical in order to reach as close as possible

to the corner, at the middle of the specimen. The distance of the thermocouple junction

from the corner was estimated at about 0.06”. After inserting the thermocouple, it had to

be fixed in place in order to minimize the errors given by the eventual displacement of

the tip from the center bottom of the hole. The temperature values for different

immersion times were then recorded on a computer. After the first set of measurement on

the 1.5” diameter, the cooling line diameter was increased by machining to 1.6”, and

subsequently to 1.7” and 1.8”. This procedure ensures excellent relativity, since the

thermocouple and its location were constant.

29

A Hitachi S-4500 Scanning Electron Microscope (SEM) was used to study the

microstructure of the materials. The specimens were polished and then etched in 2 %

Nital solution. The attached Energy Dispersive Spectrometer (EDS) was used to

determine the composition of carbides.

2.2.5. SCANNING ELECTRON MICROSCOPY

The microhardness was taken at the middle of the specimen, starting from the

corner towards the center. The first measurement was made at 0.01" from the edge, then

at 0.02", 0.04" and so on until no further variation in hardness was obtained. A

supplemental set of measurements were performed on the cross section of the 12 seconds

immersion time specimen, as seen in the hardness distribution chart, Figure 3.14. The 12

seconds specimen was chosen due to the severe conditions that it has been subjected

compared to the other specimens.

The microhardness of the specimens was measured before testing and after testing

at 5,000, 10,000, 15,000 cycles. A Buehler Micromet 2100 Microhardness Tester was

used to obtain a profile distribution of hardness from the specimen corner to the center at

both the surface and inside the specimen. A Vickers indenter was used, with a 500 g

indentation load. The Vickers hardness was converted to Rockwell C scale directly by the

tester's scale converter.

2.2.4 Microhardness Measurement

30

Element C Si Mn Cr Mo V Ni P S Fe

Weight % 0.40 1.00 0.40 5.25 1.50 1.00 0.11 0.018 <0.001 bal

Typical Composition of AISI/SAE H13

0.32-0.45 0.80-1.2 0.20-

0.50 4.75-5.50

1.10-1.75

0.80-1.20

max. 0.30

max. 0.025

max 0.005 bal

TABLE 2.1. Chemical Composition of Experimental Material - Premium Grade H13

31

Density lb/in3 (g/cm3)

Coefficient of Thermal Expansion, linear µin/in.°F (µm/m.°C)

Thermal Conductivity BTU.in/ft2.h.F (W/m.K)

Elastic Modulus ksi (GPa)

6.11 (11) 25-95 °C 169 (24.3) 215 °C 6.39 (11.5) 25-205 °C 169.3 (24.4) 350 °C 0.282 (7.8) 6.89 (12.4) 25-540 °C 171.4 (24.7) 605 °C

30,500 (210)

22

Tempering Temperature °F (°C)

Tensile Strength ksi (MPa)

Yield Strength ksi (MPa)

Reduction in Area %

Hardness Rockwell C

Impact Energy ft-lbf (J)

980 (525) 284 (1960) 228 (1570) 46.2 52 12 (16)

1120 (605) 217 (1495) 187 (1290) 54 44 22 (30)

TABLE 2.2. Typical Properties of Premium Grade H13 [44,45]

32

Specimen Cooling Line Diameter Immersion Time

A(*) 1.5" 9 sec

B 1.5" 5 sec

C 1.5" 7 sec

D 1.5" 12 sec

E 1.6" 9 sec

F 1.7" 9 sec

G 1.8" 9 sec

23

33

TABLE 2.3. Characteristics of the Tested Specimens

(*) Reference Specimen

A cooling - oil quench, martensite, no carbides B cooling - air cooling, martensite+carbides

FIGURE 2.1. CCT Diagram for H13 steel [4]

35

(*)

FIGURE 2.2. The Reference Specimen for Thermal Fatigue Test

(*) Three other specimens had the cooling line diameter 1.6", 1.7" and 1.8" respectively

36

FIGURE 2.3 The Thermal Fatigue Test Equipment

37

Φ0.08”

FIGURE 2.4 Temperature Measurement

38

3. RESULTS AND DISCUSSION

3.1. Softening During Thermal Cycling and Thermal Fatigue Resistance

During the aluminum die casting process, some parts of the die are subjected to

very severe conditions of temperature and consequently, stress. Generally, these are thin

sections, fingers and corners, where the heat transfer is two-dimensional and the amount

of energy that the material must absorb is much higher than the average for the rest of the

die. Often these are the sections that fail first. It becomes critical to create conditions for

rapid heat extraction from the surface, dissipation inside the material, or transfer towards

a “heat conveyor” such as a cooling line.

The main mechanism and the most frequent manifestation of die failure is thermal

fatigue cracking. It has been shown in previous investigations [15,40] that the strength is

very important in controlling crack initiation. It is also known that the mechanical

properties are directly related to the hardness of the material (Figure 3.1).

The capability of a steel to preserve good mechanical properties during cycling at

temperatures above the tempering temperature, is essential in establishing a satisfactory

level of performance to be expected. This assertion leads directly to the interest in the

phenomenon of softening during thermal cycling of die casting dies. This subject will be

considered and analyzed in this work. Hot-work tool steels like H13 are used in quenched

and tempered condition.

In this work, the extent of softening and the means to minimize it were evaluated. The

effect of immersion time and diameter of the cooling line on the temperature and

39

softening of the surface were studied. In both cases, there are differences regarding the

heat supply and extraction to and from the surface. In the case of different immersion

times, the amount of heat supplied to the surface of the specimen is limited by the time

spent in contact with molten aluminum. The heat extraction capacity is determined by the

size of the cooling line. If the immersion time is constant, a constant amount of heat is

supplied to the surface. However, a larger cooling line diameter will enhance the capacity

of heat extraction. In production, it is very difficult to control the time spent by the

casting in the die. The cycle length is limited by solidification time, especially for large

parts. Under these conditions, designing the die with cooling lines closer to the surface

may be the only feasible solution. One must be cautious and consider the limits set by the

hoop stresses, which are increasing with the temperature gradient. Nevertheless, varying

immersion time is very useful for the proposed study due to the ability to simulate

extreme conditions that may occur during the die life.

3.1.1. The Influence of Immersion Time on Softening and Thermal Fatigue Cracking

The experiment involved testing of three specimens for which the immersion time

in the molten aluminum was the only variable. The maximum and minimum temperatures

reached at the corner and the temperature distribution inside the specimen (toward the

cooling line) varied as a function of the time spent in the molten metal bath. Thermal

fatigue behavior of the three specimens was compared with the reference 9 seconds

immersion time specimen. The results are presented as the Total Crack Area and the

Average Maximum Crack Length for each immersion time (Figure 3.2 and Figure 3.3).

40

The hardness measured at the corner of the specimens after 15,000 cycles is shown in

Figure 3.4.

There is a clear trend for the thermal fatigue cracking parameters (i.e. Total Crack

Area and Average Maximum Crack Length) to increase with immersion time. This

observation points at the main cause of thermal fatigue damage, which is the temperature

variation during cycling. One of the direct effects of the temperature, in particular the

maximum temperature reached at the corner, is the softening of the steel. The extent of

softening during tempering is generally evaluated by a master parameter, known as

Hollomon-Jaffe parameter. This value represents the combined effect of temperature and

time. Since the temperature and time are interdependent variables in the thermally

activated process of tempering, a trade-off of temperature for time or vice-versa is based

upon a simple equation:

P = T(C + logt) x 10-3 (3-1)

where P is the Hollomon-Jaffe parameter

T is the absolute temperature [K]

t represents the time [hours]

C is a material constant

This equation yields a reasonably accurate prediction of hardness for carbon and alloy

steels containing 0.2-0.85% carbon and less than 5% total alloying elements, irrespective

of initial structure. It is not the scope of this work to investigate the hardness of steel as a

function of temperature-time. The hypothesis is however that the temperature is the main

41

factor that causes hardness loss. The dependency of thermal fatigue cracking on the level

of hardness is investigated in detail.

The dependency of Total Crack Area and Maximum Crack Length on the

maximum temperature at the corner of the specimen for different immersion times is

presented in Figures 3.5 and 3.6. The results demonstrate that the higher the temperature,

the more thermal fatigue damage will occur. As previously discussed, a higher

temperature will produce a more severe and deeper softening of the surface and within

the section of the specimen. Initially, the surface deformation (strain) is within the elastic

capabilities of the die steel. The surface of the specimen has irregularities in the forms of

corrosion pits or surface scratches. These sites serve as stress concentrations. Plastic

deformation can therefore occur at stresses well below the yield strength of the parent

material (it must be also noted that the strength of the material drops at high temperature,

see Figure 3.7), and initiate fatigue cracks. In addition to the stress concentrations caused

by surface imperfections, tempering weakens the surface material. A cumulative fatigue

process occurs in the material, since plastic strain gradually increases during the test as a

result of lower yield strength of the material. The compressive stress will eventually

exceed the elastic limit of the steel and plastic deformation will take place after the initial

elastic strain has occurred [50]. Under these conditions, it is therefore necessary for the

material to drop below of certain strength level characterized by a lower hardness value

in order for the crack to initiate.

It has been experimentally demonstrated that if the strength properties of the

material are reclaimed before the cracks initiated, the thermal fatigue behavior can be

42

markedly improved. The experiment consisted of cycling a H13 steel specimen for 2,500

cycles and then re-heat treating it to the original hardness value. The results compared to

regular 51 HRC and 46 HRC H13 specimens are presented in the Figure 3.8. It is clearly

shown that the re-heat treated specimen to 51 HRC after every 2,500 cycles exhibited

better resistance against heat checking. The cyclic heat treatment reclaimed the strength

of the material and its resistance against cracking, impeding crack initiation, as well as

the propagation of the existent cracks. Based on this evidence, it is believed that for a

certain combination of temperature/stress the crack initiation will occur at a

correspondent value of hardness. Therefore, it is expected that in a specimen subjected to

a higher maximum temperature the hardness will drop faster. The higher drop in strength

during immersion will thus cause cracks to initiate earlier. The cracks have then more

time to grow, and the Average Maximum Crack Length will presumably be higher. The

relationship between Average Maximum Crack Length and Total Crack Area is presented

in Figure 3.9. Longer cracks correspond to a higher value of Total Crack Area. In

addition, more cracks may initiate at the weakened surface, grow faster, and contribute to

a higher Total Crack Area.

At the same time, the behavior of the propagating crack is influenced by the

characteristics of the material at the crack tip, and hence by the ability to resist plastic

deformation. A parameter was chosen, which could provide information about the

properties ahead the crack tip/front, namely the microhardness at a distance equal to the

Average Maximum Crack Length (Figure 3.10). It is asserted that this distance

characterizes well the propagation of cracks inside the specimen. The dependency of this

43

new parameter on the temperature measured at the corner is presented in Figure 3.11 and

appears to have a linear trend. The relationship between the cracking parameters and the

Immersion Time 5 sec 7 sec 9 sec 12 sec

Maximum Temperature [F] 926 991 1087 1147

Minimum Temperature [F] 322 346 399 460

Total Crack Area [x 106 µm2] After 15,000 Cycles 1.97 5.9 108.56 167.72

Average Maximum Crack Length After 15,000 Cycles [x 100 µm] 2.25 3 12.5 15.25

Hardness at the Average Maximum Crack Length [HRC] After 15,000 Cycles 36.9 33.8 31.6 29.2

TABLE 3.1. Measurement Data For Different Immersion Times

44

Distance From the Corner [in] 5 sec 7 sec 9 sec 12 sec

0.01 36.9 33.8 24.5 24.3

0.02 39.4 36.5 27.4 25.1

0.04 42.4 39.1 31.2 27.7

0.06 43.3 42.4 32.3 29.2

0.08 44.1 44.2 34.5 29.2

0.1 44.3 44.3 36.2 30

0.2 44.4 44.5 40.9 34.1

TABLE 3.2: Immersion Time Effect on Hardness Variation Across the Surface

45

microhardness measured at the distance equal to the Average Maximum Crack Length is

described in Figures 3.12 and Figure 3.13.

It can be concluded that a higher maximum temperature will accelerate the loss in

hardness at the corner of the specimen. The crack will extend to a longer distance, as it

will have more time to propagate. The hardness loss will be also more severe further

inside the material. It appears that the longer the crack is, i.e. the higher the temperature

at the surface, the lower the hardness ahead of it. A possible explanation of this

phenomenon is that the thermal stresses decrease from the surface towards the interior as

the temperature gradient drops, mainly due to the decrease in the maximum temperature.

In order for the crack to advance, the strength of the material must decrease even more.

However, in this particular configuration, the cooling line does not allow a very deep

softening and the crack may eventually stop before it attains a critical length that will

lead to instability.

In addition to the surface microhardness evaluation, internal hardness

measurements were taken on a center section of the 1.5" cooling line diameter and 12

seconds immersion time specimen. This specimen was under to the most severe test

conditions among all the specimens used. The sample was sectioned at the center (about

3.5" from the both ends). The Rockwell hardness profiles obtained from these HV

microhardness values are shown in Figure 3.14.

The hardness distribution plots indicate that the softening near the edge is

significantly higher than inside the sample, which is predictable due to the higher

temperature and heat transfer conditions. However, the degree of softening exceeds the

typical tempering curve, shown in the Figure 3.15. The temperature measured at about

46

0.06" from the corner of the specimen is almost 1150 °F. The temperature cycle at this

location has a minimum at 460 °F and a maximum at 1147 °F. If the peak portion of the

thermal cycle is separated (from the cycle presented in Figure 3.16), it will show that the

specimen resided a total of 5.5 sec x 15,000 cycles = 82,500 seconds ~ 23 hours, at a

temperature between 1100 °F and 1150°F. The value of the hardness at 0.06" from the

corner was measured to be 30.6 HRC. According to the tempering curve of the steel

(from Figure 3.4), such a drop in hardness from 45 to around 30 HRC would be produced

after about 23 hours at 1150 °F. It is concluded that another mechanism contributed to the

softening, presumably cyclic stress softening.

3.1.2. The Influence of Cooling Line Diameter on Softening and Thermal Fatigue

Cracking

Frequently, critical sections (usually thin parts or complicated shape sections

subjected to multidirectional heat transfer) occur within a die-casting die. These sections

are under high temperature and severe stress conditions. The importance of the maximum

temperature and its influence on softening, and hence on thermal fatigue cracking was

discussed in the previous section.

In this experiment the maximum and the range of temperature reached at the

corner and the variation inside the specimen was investigated as a function only of the

cooling line diameter. A larger cooling line will actually bring down the maximum

temperature at the surface, and at the same time will keep the temperature range almost

the same, since the minimum temperature drops as well, because of a higher heat

extraction capability.

47

The thermal fatigue behavior of three specimens with different cooling line

diameters, 1.6", 1.7" and 1.8", was compared with the reference 9 seconds immersion

time - 1.5" cooling line diameter specimen. The results are presented as the Total Crack

Area and the Average Maximum Crack Length for each cooling line diameter (Figures

3.17 and 3.18). The evaluation of softening or hardness loss at the corner of the

specimens after 15,000 cycles is shown in Figure 3.19. The same trend as in the previous

experiment was observed for the thermal fatigue cracking parameters. The values of

Total Crack Area and Average Maximum Crack Length decrease with the increase in the

maximum temperature

The effect of different cooling line diameter on Total Crack Area and Average

Maximum Crack Length as a function of maximum temperature at the corner of the

specimen is presented in Figures 3.20 and 3.21. The curve seems to reach a plateau as

the cooling line diameter becomes smaller. If the curves are compared with those

obtained for varying immersion times, it will be noticed that the tendency of the curve to

level around a certain maximum temperature is common for both situations. The

variation of microhardness measured at the Average Maximum Crack Length is shown in

Figure 3.22. The dependency of the thermal fatigue cracking on the microhardness

measured at the Average Maximum Crack Length is presented in Figures 3.23 and

3.24.The relationship between the cracking parameter and the microhardness measured at

the distance equal to the Average Maximum Crack Length follows the temperature trend,

confirming the observation made for different immersion times. The longer the crack is,

because of the higher temperature at the surface, the lower the hardness ahead the crack.

48

Cooling Line Diameter [in] 1.8" 1.7" 1.6" 1.5"

Maximum Temperature [F] 909 939 1002 1087

Minimum Temperature [F] 197 237 326 399

Total Crack Area [x 106 µm2] After 15,000 Cycles 35 62.05 79.44 108.56

Average Maximum Crack Length After 15,000 Cycles [x 100 µm] 8.25 10 12.25 12.5

Hardness at the Average Maximum Crack Length [HRC] After 15,000 Cycles 35.1 34.6 34.1 31.6

TABLE 3.3: Measurement Data For Different Cooling Line Diameters

49

Distance From the Corner [in] 1.8" 1.7" 1.6" 1.5"

0.01 30.6 29.4 29.3 24.3

0.02 33.7 31.3 30.6 25.1

0.04 35.6 34.2 33.9 27.7

0.06 38.6 37.3 36.3 29.2

0.08 40.1 38.2 36.9 29.2

0.1 40.5 39.4 38.3 30

0.2 43.5 42.9 42 34.1

TABLE 3.4: COOLING LINE DIAMETER EFFECT ON HARDNESS VARIATION

ACROSS THE SURFACE

50

3.2. Stress Analysis at the Specimen Surface and Around the Cooling Line

The stresses developed at the surface are responsible for initiation and subsequent

crack propagation. The required hardness loss for crack initiation and propagation varies

function of the level of induced stress. These thermal stresses are generated by the

difference between the maximum and minimum temperature (temperature gradient).

Different testing or production conditions will result in different temperature and stress

distributions. The effect of immersion time and cooling line diameter on maximum,

minimum and range of temperature are presented in Figures 3.25 and 3.26. More severe

conditions (longer immersion time or smaller diameter of the cooling line) will not raise

only the maximum temperature but also the minimum temperature, so that the

temperature gradient (range, in the plots) will not increase too much. Consequently, a

larger cooling line or a shorter immersion time will shift the overall cycle towards lower

values, keeping the stress in about the same range. This observation is extremely

important, because of the implications resulting from the capacity of a larger cooling line

diameter to promote a lower softening-causing maximum temperature without a major

increase in the stress level.

The stresses in the thermal fatigue specimen are complex. The cracks initiate not

only at the corner of the specimen, where softening favors the plastic strain accumulation

(Figure 3.27), but also at the cooling line, due to high hoop tensile stresses created during

immersion. The latter formation is promoted by the existence of severe stress

concentrators caused by cooling water corrosive action. In extreme conditions, the cracks

initiated at the cooling line can cause failure of the specimen, mainly because they initiate

51

and grow faster in the thinnest section of the specimen or die as a consequence of high

tensile hoop stress induced by the extreme temperature gradient (Figure 3.28). The axial

stress range at the corner can be estimated using equation 1-1. For 12 seconds immersion

and 1.5" cooling line diameter specimen:

σ = αE∆T = 6.9 µin/in°F * 30,500 ksi * 687 °F = ~145 ksi.

where α is the coefficient of thermal expansion, E represents the elastic modulus and ∆T

is the temperature gradient. This estimation agrees relatively well with the value of stress

range obtained by computer modeling (Figure 3.29). The computer modeling for the 1.5”

cooling line diameter specimen and 12 seconds immersion time shows that during

immersion in molten aluminum the compressive axial stress at the corner attains a high

value. High compressive stress and low yield strength may generate plastic strain. The

result is a residual tensile stress, which is well below the yield strength, but high enough

to initiate fatigue cracks at stress concentrators. The axial stress at the cooling line is

tensile. Because the temperature at the wall of the cooling line is low, the axial stress is

tolerated. However, the hoop stress developed is markedly higher. In the presence of

stress concentrators like machining marks and corrosion pits, cracks can initiate and

propagate from the cooling line.

3.3. Microstructure Degradation that Promotes Softening During Thermal Cycling

The alloying elements present in the steel affect the hardening, tempering

characteristics and the carbides in steels. As a consequence, they have the ability to

impart certain features to die steel [48]:

52

- Greater strength in large sections because of deeper hardening or increased

hardenability. In steels, strength is virtually proportional to hardness.

- Less distortion in the process of hardening by increased hardenability due to the ability

to harden the steel with a less drastic quench. Less distortion, dimensional change, and

quench cracking are direct results of this lowering of thermal stresses set up by large

temperature gradients.

- Greater resistance to abrasion at the same hardness by promoting the formation of hard,

stable and wear-resistant carbides.

- Alloying elements induce higher toughness in small sections by promoting fine grain

size. They also lower the internal stresses through less drastic quenches, and permit a

greater relief of internal stresses through the use of higher tempering temperature without

much loss of hardness.

As far as the individual effect of particular elements, it is known that

molybdenum is effective in improving the hardenability and high temperature strength. It

retards the softening of martensite at all tempering temperature and reduces susceptibility

to tempering embrittlement. Above 1000 °F, the presence of molybdenum keeps the size

of carbides small. Like molybdenum, chromium also retards the softening of martensite.

By substituting chromium for some of the iron in cementite, the coalescence of carbides

is retarded. However, its effect on the hardenability is less than that of molybdenum.

Vanadium is a stronger carbide forming element than the above two elements. The

vanadium containing carbides are stable at elevated temperature. Thus the steel has to be

53

austenitized at a sufficiently high temperature and for a sufficient length of time to bring

most of the carbides into solution. For instance, when the H13 tool steel is austenitized at

1010 °C (1850°F) for an hour, the molybdenum and chromium carbides are dissolved in

solid solution, but the vanadium carbide (VC or V4C3) does not dissolve [29, 30, 35].

Also, the precipitation of vanadium carbides and carbonitrides in high strength low alloy

(HSLA) steels raises the strength of these steels well above the normally processed mild

steels [24, 25]. Silicon not only has its own potential in increasing hardenability, but also

stabilizes the ε−iron carbide upon tempering to such an extent that it is still present in the

microstructure after tempering at 400°C in steels with 1-2 % Si. Silicon slows down the

nucleation and growth of the carbide and also enters into the carbide structure, delaying

the transformation of ε to Fe3C. However, if in large quantity, silicon precipitates on

grain boundaries, martensite lath boundaries and/or martensite lath/carbide interfaces

during tempering [26, 27], which enhances the embrittlement and lower the toughness of

die steels. A lower content of silicon presumably minimizes the interfacial segregation

and results in higher toughness and thermal fatigue resistance [6].

Thermal fatigue resistance is affected by the combination of primary carbides formed

in as quenched condition and carbides precipitated during tempering. Smaller and fewer

carbides in as quenched conditions make the crack initiation hard, and the well dispersed

carbide precipitation pattern makes it harder for crack to propagate [11]. The influence of

the amount of primary carbides on the impact toughness properties of H13 is shown in

Figure 3.30.

54

The effect of temperature and hardness on thermal fatigue behavior of quenched and

tempered H13 steel is supported by the observations made on the microstructure of

quenched and tempered H13 steel at different distances from the corner, in the cross-

section of the 12 seconds immersion time specimen (Figures 3.31). As the temperature

increases, the coarsening of the carbides becomes more severe (Figures 3.32, a – d).

The photomontage in Figure 3.33 illustrates the distribution of carbide in this cross-

section. As the temperature decreases from the surface towards the cooling line, the

carbides become finer. Figure 3.34 (a) shows the microstructure of H13 sample in the

unaffected material. The base material shows agglomeration of carbides from the original

temper and some larger carbides probably remained undissolved during the austenitizing.

These carbides are concentrated primarily at the austenite grain boundaries and between

the lathes of martensite [15].

Figure 3.30. Effect of Volume Percent Primary Carbides on the Transverse Charpy V-notch Impact Toughness of H13 [41]

55

02468

1012141618202224

0 0.005 0.01 0.015 0.02 0.025 0.03

Avg. Vol. Primary Carbides

Tran

sver

se C

VN (f

t - lb

s.)

1875 oF - 25 min., Oil quenched1135 oF - 2hrs. Air Cooled

- Surface - Center

Tracking the carbides in a complex alloy steel like 5CrMoV(H13) back to the

annealed structure, it is found that the total weight percentage of carbides present is about

4.4 %. After austenitizing at 1850 °F, only 2.3 % of weight represents undissolved

carbides (Figure 3.34), of which most is vanadium and some molybdenum. Further

increase in austenitizing temperature results in extensive dissolution of Mo and V. By

1950 – 2000 °F, most of the molybdenum is in solution and the weight percentage of

undissolved carbides is about 1.5 % and most of this is vanadium carbide [49]. Previous

studies on the effect of austenitizing temperature on the amount of carbides present in the

quenched microstructure of other tool steels have shown the same dependency [3]. The

important characteristic of H13 steel is the presence of a higher content of vanadium than

in other tool steels. Vanadium carbide tends to be more stable at higher austenitizing

temperatures. The hardness after quenching is a good indicator of the effect of dissolution

of vanadium carbide. A jump from about 59 to 61 HRC has been observed by increasing

the austenitizing temperature from 1850-1950 °F range to 1950-2100 °F.

Alloying elements also affect the softening resistance during tempering. They restrain

the coarsening of cementite in the range 400 – 700 °C (Si, Cr, Mo, W) either by entering

into the cementite structure or by segregating at the carbide-ferrite interfaces. Secondly,

in alloy steels such as H13, a number of alloying elements form fine carbides that are

thermodynamically more stable than cementite. The alloying elements Cr, Mo, V, W and

Ti form carbides with substantially higher enthalpies of formation [47]. When strong

carbide forming elements are present in steel in sufficient concentration, their carbides

will be formed in preference to cementite. However, during tempering of alloy steels,

56

alloy carbides do not form until 500 – 600 °F. Below this temperature range the metallic

alloying elements cannot diffuse fast enough to allow alloy carbides to nucleate. The

metallic elements diffuse substitutionally, in contrast to carbon that diffuses interstitially.

Hence, the diffusivity of carbon is several orders of magnitude greater in iron than those

of the metallic alloying elements.

The coarsening of carbides in steel influences markedly the mechanical properties.

The strengthening theories show that the yield strength of a dispersed alloy, controlled by

the capacity of dislocations to move around spherical particles, varies inversely with the

spacing between particles. If the carbide dispersion is coarsened by further heat

treatment, the hardness and strength of the alloy falls [47]. The theory for coarsening of a

dispersion shows that the coarsening rate is dependent on the diffusion coefficient of the

solute:

rt3 – r0

3 = (k/RT) Vm2 D σ t (3-2)

where

r0 = the initial mean particle radius

rt = the mean particle radius at time t

D = diffusion coefficient of solute in matrix

σ = interfacial energy of particle/matrix interface per unit area

V= molar volume of precipitate

k = constant

57

Under any given temperature, cementite will coarsen at a higher rate than any of the

alloy carbides. This is typical in alloy steels in which cementite and an alloy carbide

coexist, where cementite dispersion is always much coarser.

A basic Energy Dispersive Spectrometry (EDS) analysis of carbides in the over

tempered structure of the H13 near the corner of the specimen after 15,000 cycles, has

shown that the largest carbides in the microstructure are Cr-rich. The smaller carbides are

Mo-rich carbide (Figures 3.35 and 3.36). Chromium diffuses more rapidly in ferrite than

most metallic alloying elements, with the result that in chromium steels Cr7C3 is detected

during tempering at temperature as low as 500 °C, and it coarsens rapidly compared to

molybdenum or vanadium carbides [47]. Thus, in chromium steel, continuous softening

will normally occur during tempering between 500 – 700 °C, although the addition of

other elements, such as Mo, can reduce the rate of coarsening of Cr7C3. Also, previous

works have shown that during tempering at 1200 °F of 5CrMoV steel, an iron-rich

chromium carbide forms, (CrFe)7C3 [16]. The small carbide appears to be Mo-rich, in the

form of M2C or eventually M6C, explainable by the fact that molybdenum carbide is less

sensitive to growth. Vanadium-rich carbides were not detected EDS, even though some

vanadium was found to be present in Cr-rich carbides. This is due to the ability of

vanadium to maintain a very fine carbide (VC or V3C4) dispersion even at temperatures

approaching 700 °C. The detection of vanadium-rich carbides by EDS method is at best

difficult.

58

4. CONCLUSIONS

1) For a configuration without severe stress concentrators, the softening of the steel is

the most important factor for the crack initiation. Less thermal fatigue damage has been

observed when the conditions promoted lower temperature at the surface, which

preserved the hardness and hence the strength. A high value of yield strength means

higher material resistance to plastic deformation. At the same time, elevated temperature

at the surface will induce a deeper softening. It appears that a condition for the extension

of the thermal fatigue cracking damage is the decrease in strength ahead the crack front.

2) In die-casting applications, the highest maximum temperature will occur in thin

sections where the material capacity to absorb and transfer the heat from the surface is

very different. From another point of view, high temperature - long resident time

conditions are important, because of the similarity with the die casting of large

components, when the die is subjected to elevated temperature for longer periods of time.

The experimental results have shown an important decrease of the cracking when the

cooling line is positioned closer to the surface. Moreover, the experimental data indicates

the existence of a temperature threshold, below which the thermal fatigue damage is

minimal. A cooling line closer to the surface will shift the maximum temperature towards

lower values, and keep at the same time the stresses at a relative constant value. However,

decreasing the maximum temperature at the surface by placing the cooling lines too close

to the surface may be limited by the high level of hoop stresses created at the cooling line.

59

3) The presence of strong carbide-former elements like chromium, molybdenum

and vanadium, will reduce the softening by preserving a fine distribution of carbides.

These elements inhibit the coarsening of cementite in the range 400 – 700 °C. At the

same time, these elements form fine carbides that are thermodynamically more stable

than cementite. Among the three elements, chromium-rich carbide is the most susceptible

to growth, but the presence of molybdenum and vanadium inhibits it to certain measure.

60

61

FIGURE 3.1. Relationship Between Tensile Properties and Hardness for H13 Steel [45, Reprinted with permission of American Society of Materials]

0

20

40

60

80

100

120

140

160

180

5000 10000 15000

Number of Cycles

Tota

l Cra

ck A

rea

[x 1

06 µm2 ]

5 sec 7 sec

9 sec 12 sec

all below 0.2 5 sec

12 sec

9 sec

7 sec

The Effect of Immersion Time

FIGURE 3.2. The Effect of Thermal Cycling on Crack Area

- Different Immersion Times -

63

0

2

4

6

8

10

12

14

16

18

20

5000 10000 15000

Number of Cycles

Ave

rage

Max

imum

Cra

ck L

engt

h [x

100

µm]

5 sec 7 sec

9 sec 12 sec

The Effect of Immersion Time

5 sec7 sec

9 sec

12 sec

FIGURE 3.3. The Effect of Thermal Cycling on Crack Length - Different Immersion Times -

64

0.01 0.02 0.04 0.06 0.08 0.1 0.220

25

30

35

40

45

Har

dnes

s H

RC

Distance from the Corner [in]

12 sec9 sec

7 sec

5 sec

The Effect of Immersion Time

15,000 cycles

FIGURE 3.4. The Effect of Thermal Cycling on Microhardness Distribution Across the Surface

- Different Immersion Times -

65

0

20

40

60

80

100

120

140

160

180

200

900 950 1000 1050 1100 1150 1200

Temperature [F]

Tota

l Cra

ck A

rea

[x10

6 µm

2 ]

5 sec

9 sec

12 sec

7 secThe Effect of Immersion Time

15, 000 cycles

FIGURE 3.5. The Effect of Temperature on Crack Area

- Different Immersion Times -

66

0

2

4

6

8

10

12

14

16

18

900 950 1000 1050 1100 1150 1200

Temperature [F]

Ave

rage

Max

imum

Cra

ck L

engt

h [x

100

µm] 15, 000 cycles

The Effect of Immersion Time

5 sec7 sec

9 sec

12 sec

67

FIGURE 3.6. The Effect of Temperature on Crack Length

- Different Immersion Times -

68

FIGURE 3.7. Effect of Elevated Temperature on Tensile Strength [45, Reprinted with permission of American Society of Materials]

70

0

50

100

150

200

250

300

350

400

450

500

2500 5000 7500 10000 12500 15000

Thermal Cycles

Tota

l Cra

ck A

rea

[x 1

06 µm2 ]

H13 at 46 HRC

H13 at 51 HRC

H13 Re-Heat Treated to 51 HR

FIGURE 3.8. The Effect of Hardness Recovery on Thermal Fatigue Cracking

C After Every 2,500Cycles

2BAR/46HRC

2BAR/51HRC

OIL/51HRC2BAR Quench+Double

Temper to 51HRC

71

FIGURE 3.9 Relationship Between Total Crack Area and Average Maximum Crack Length

FIGURE 3.9 Relationship Between Total Crack Area and Average Maximum Crack Length

0

20

40

60

80

100

120

140

160

180

0 2

Av

Tota

l Cra

ck A

rea

[x 1

06 µm

2 ]

4 6 8 10 12 14 16 18

erage Maximum Crack Length [x 100 µm]

T1 < T2

FIGURE 3.10. The Relationship Between Maximum Crack Length and Microhardness at Maximum Crack Length

72

25

30

35

40

900 950 1000 1050 1100 1150 1200

Temperature [F]

Mic

roha

rdne

ss a

t Max

imum

Ave

rage

Cra

ck L

engt

h [H

RC

]

The Effect of Immersion Time

5 sec

7 sec

9 sec

12 sec

15, 000 cycles

FIGURE 3.11. The Effect of Temperature on Microhardness

- Different Immersion Times-

73

0

20

40

60

80

100

120

140

160

180

2829303132333435363738

Microhardness at the Average Maximum Crack Length [HRC]

Tota

l Cra

ck A

rea

[x 1

06 µm

2 ]

12 sec

9 sec

7 sec5 sec The Effect of Immersion Time

15, 000 cycles

FIGURE 3.12. The Effect of Microhardness at Average Maximum Crack Length on Crack Area

- Different Immersion Times -

74

0

2

4

6

8

10

12

14

16

18

283032343638

Microhardness at the Average Maximum Crack Length [HRC]

Ave

rage

Max

imum

Cra

ck L

engt

h [x

100

µm

]

5 sec7 sec

9 sec

12 sec

The Effect of Immersion Time

15, 000 cycles

75

FIGURE 3.13. The Effect of Microhardness at Average Maximum Crack Length on Crack Length

- Different Immersion Times -

FIGURE 3.14. Microhardness Profile at the Corner of 12 Seconds Immersed Specimen

76

77

FIGURE 3.15. Tempering Curve for

H13

400450500550600650700750800850900950

10001050110011501200

0 5 10 15 20 25 30 35

Time [sec]

Tem

pera

ture

[F]

FIGURE 3.16: Maximum Temperature Cycle for 1.5" Cooling Line Specimen After 12 Seconds Immersion Time

78

0

20

40

60

80

100

120

140

160

180

5000 10000 15000

Number of Cycles

Tota

l Cra

ck A

rea

[x 1

06 µm

2 ]1.8" 1.7"

1.6" 1.5"

The Effect of Cooling Line Diameter

1.5"

1.8"

1.7"

1.6"

all below 0.2

FIGURE 3.17. The Effect of Thermal Cycling on Crack Area - Different Cooling Line Diameters -

79

0

2

4

6

8

10

12

14

16

18

5000 10000 15000

Number of Cycles

Ave

rage

Max

imum

Cra

ck L

engt

h [x

100 µ

m] 1.8" 1.7"

1.6" 1.5"1.5"

1.8"

1.7"

1.6"

The Effect of Cooling Line Diameter

FIGURE 3.18. The Effect of Thermal Cycling on Crack Length

- Different Cooling Line Diameters -

80

0.01 0.02 0.04 0.06 0.08 0.1 0.220

25

30

35

40

45

Har

dnes

s H

RC

Distance from the Corner [in]

1.5"

1.8"1.7"

1.6"

15,000 cycles

The Effect of Cooling Line Diameter

FIGURE 3.19. The Effect of Thermal Cycling on Microhardness Distribution Across the Surface

- Different Cooling Line Diameters -

81

0

20

40

60

80

100

120

900 920 940 960 980 1000 1020 1040 1060 1080 1100

Temperature [F]

Tota

l Cra

ck A

rea

[x10

6 µm

2 ]

1.6"

1.7"

1.8"

1.5"

The Effect of Cooling Line Diameter

15, 000 cycles

FIGURE 3.20. The Effect of Temperature on Crack Area - Different Cooling Line Diameters -

82

0

2

4

6

8

10

12

14

900 950 1000 1050 1100

Temperature [F]

Ave

rage

Max

imum

Cra

ck L

engt

h [x

100

µm

]

1.8"

1.7"

1.6" 1.5"15, 000 cycles

The Effect of Cooling Line Diameter

FIGURE 3.21. The Effect of Temperature on Crack Length - Different Cooling Line Diameters -

83

25

30

35

40

900 950 1000 1050 1100

Temperature [F]

Mic

roha

rdne

ss a

t Ave

rage

Max

imum

Cra

ck L

engt

h [H

RC

]

The Effect of Cooling Line Diameter

1.8"1.7"

1.6"

1.5"

15, 000 cycles

FIGURE 3.22. The Effect of Temperature on Microhardness - Different Cooling Line Diameters -

84

0

20

40

60

80

100

120

313233343536

Microhardness at the Average Maximum Crack Length [HRC]

Tota

l Cra

ck A

rea

[x10

6 µm

2 ]

1.5"

1.6"

1.7"

1.8"

The Effect of Cooling Line Diameter

15, 000 cycles

FIGURE 3.23. The Effect of Microhardness at Average Maximum Crack Length on Crack Area - Different Cooling Line Diameters -

85

0

20

40

60

80

100

120

313233343536

Microhardness at the Average Maximum Crack Length [HRC]

Tota

l Cra

ck A

rea

[x10

6 µm

2 ]

1.5"

1.6"

1.7"

1.8"

The Effect of Cooling Line Diameter

15, 000 cycles

FIGURE 3.23. The Effect of Microhardness at Average Maximum Crack Length on Crack Area - Different Cooling Line Diameters -

86

0

2

4

6

8

10

12

14

16

18

313233343536

Microhardness at the Average Maximum Crack Length [HRC]

Ave

rage

Max

imum

Cra

ck L

engt

h [x

100

µm]

1.8"

1.7"

1.6"1.5"

The Effect of Cooling Line Diameter

15, 000 cycles

FIGURE 3.24. The Effect of Microhardness at Average Maximum Crack Length on Crack Length - Different Cooling Line Diameters -

87

0

200

400

600

800

1000

1200

1400

5 sec 7 sec 9 sec 12 sec

Immersion Time

Tem

pera

ture

[F]

Maximum TemperatureMinimum TemperatureRange

FIGURE 3.25. The Effect of Immersion Time on Temperature

88

0

200

400

600

800

1000

1200

1.51.61.71.8

Cooling Line Diameter

Tem

pera

ture

[F]

Maximum TemperatureMinimum TemperatureRange

89

FIGURE 3.26. The Effect of Cooling Line Diameter on Temperature

FIGURE 3.27. CRACKS AT THE CORNER OF H13 SPE

Coo

FIGURE 3.28. CRACK AT THE COOLING LINE OF H1

90

Corner

X 45

CIMEN

X 200

ling Line

3 SPECIMEN

-100000

-80000

-60000

-40000

-20000

0

20000

40000

60000

80000

100000

0 6 12 18 24 30 36

Time [sec]

Stre

ss [p

si]

Axial Stress Node AAxial Stress Node BHoop Stress Node B

B

A

1.5” Cooling Line, 12 Seconds Immersion

FIGURE 3.29. Stress Modeling at the Corner and Cooling Line

91

Figure 3.30. Effect of Volume Percent Primary Carbides on the Transverse Charpy V-notch Impact Toughness of H13 [41]

02468

1012141618202224

0 0.005 0.01 0.015 0.02 0.025 0.03

Avg. Vol. Primary Carbides

Tran

sver

se C

VN (f

t - lb

s.)

1875 oF - 25 min., Oil quenched1135 oF - 2hrs. Air Cooled

- Surface - Center

92

0.06”

0.2”

0.1”

Corner

FIGURE 3.31. Microstructure Sampling at the Corner of 12 Seconds Immersion Specimen

93

15,000 Cycles

~30 HRC

FIGURE 3.32(D) THE EFFECT OF TEMPERATURE ON MICROSTRUCTURE –

0.06” FROM CORNER

15,000 Cycles ~27 HRC

FIGURE 3.32(E). THE EFFECT OF TEMPERATURE ON MICROSTRUCTURE –

CORNER

94

12 Seconds Immersion Time 1.5” Cooling Line Diameter

15,000 cycles

FIGURE 3.33. Temperature Influence on Carbide Size and Distribution Photomontage

95

FIGURE 3.34. Effect of Austenitizing Temperature on the Weight Percentage of Isolated Carbide Residues in H13 Steel [16]

96

FIGURE 3.36. EDS Analysis of Small Carbide in Softened H13

15,000 cycles Corner of 12 Seconds Immersion Time, 1.5” Cooling Line Specimen

FIGURE 3.35. Small Carbide in Softened H13

97

FIGURE 3.36. EDS Analysis of Small Carbide in Softened H13

15,000 cycles Corner of 12 Seconds Immersion Time, 1.5” Cooling Line Specimen

FIGURE 3.36. Large Carbide in Softened H13

98

5. BIBLIOGRAPHY Advanced Materials & Processes, ASM International, Vol. 159, No. 12, December 2001, p. 83 [45] Advanced Materials & Processes, ASM International, Vol. 159, No. 12, December 2001, pp. 61. [44] Bain, E.C.; Paxton, H.W. Alloying Elements in Steel, American Society for Metals (1961) [35] Benedyk, J.C. Thermal Fatigue Behavior of H13 Die Steels, Ph.D. Thesis, Case Western Reserve University (1969) [40] Bertolo, R.B. Fracture Toughness of Aluminum Die Casting Die Steels, Ph.D. Thesis, Case Western Reserve University, 1976 [6] Bethge, K.; Munz, D.; Neumann, J. “Crack Initiation and Propagation Under Thermal Cyclic Loading”, High Temperature Technology Vol. 8 No. 2 (1990) [14] Brick, R.M.; Pense, A.W.; Gordon, R.B. Structure and Properties of Engineering Materials, Fourth Edition, McGraw-Hill (1977), p. 307 [29] Brick, R.M.; Pense, A.W.; Gordon, R.B. Structure and Properties of Engineering Materials, Fourth Edition, McGraw-Hill (1977), pp. 152-163 [38] Brooks, C.A. Principles of the Heat Treatment of Plain Carbon and Low Alloy Steels, ASM International (1996) [36] Campbell, I.E. (Editor- in-Chief) High Temperature Technology, John Wiley and Sons (1957), pp. 460-476 [21] Contractor, G.P.; Schempp, E.G.; Morgan, W.A. “A Study of Carbide Composition and Microstructure During Quenching and Tempering of a 5% CrMoV Steel”, Trans. ASM, Vol. 54 (1961) pp. 208-219 [49] Das, S. K. Effect of Heat Treatment on the Thermal Fatigue Behavior and Fatigue Toughness of H13 Steel for Aluminum Die Casting Dies, Ph.D. Thesis, Case Western Reserve University (1981) [16] Dieter, G. Mechanical Metallurgy, McGraw-Hill (1986), pp. 430 [8] Engineering Properties of Steel, P.D. Harvey-editor, American Society for Metals (1982),

99

pp. 460-461 [46] Failure Analysis and Prevention, ASM Handbook Vol. 11, 9th Edition, American Society for Materials Publications, pp. 266-267 [5] Garrison, W.M., Jr. “Influence of Silicon on Strength and Toughness of 5 wt. % Cr Secondary Hardening Steel”, Materials Science and Technology, No 3(1987) [26] Graham, R.R. Thermal Fatigue Mechanisms in Aluminum Die Casting Die Steel, Ph.D. Thesis, Case Western Reserve University (1978) [15] Heat-Resistant Materials-Elevated-Temperature Mechanical Properties of Stainless Steel, ASM Handbook (1997), edited by J.R. Davis, pp. 125-127 [28] Honeycombe, R.W.K.; Bhadeshia, H.K.D.H. Steel, Microstructure and Properties, Second Edition (1982) pp. 181-197 [47] Kaye, A. Die Casting Metallurgy, Butterworth Scientific (1978) [3] Krauss, G. Steels: Heat Treatment and Processing Principles, ASM International (1990), [24] Lu, T.J.; Fleck, N.A. “The Thermal Shock Resistance of Solids”, Acta Materialia, Vol. 46, No. 13 (1998) [19] Manson, S.S. Thermal Stress and Low Cycle Fatigue, McGraw-Hill (1966), pp. 21-22, 303 [20] Martensite, E.R. Petty-ed., Longmans (1970), p. 5 [32] Muscatell, et al. “Thermal Shock Resistance of High Temperature Alloys”, Proc. Amer. Soc. Test. Mat., Vol. 57 (1957), p. 947 [42] Noesen, S.J.; Williams, H.A. “The Thermal Fatigue of Die Casting Dies”, Modern Casting, Vol. 51, No. 6, American Foundrymen’s Society (1967), pp. 119-132 [50] Park, J.H. et al. “The Effects of Alloying Elements on Thermal Fatigue and Thermal Shock Resistance of the HSLA Cast Steels”, ISIJ International, Vol. 40 No. 11 (2000) [25] Porter, D.A.; Easterling, K.E. Phase Transformations in Metals and Alloys, Chapman&Hall (1992), pp. 417-437 [33] Roberts, G.A.; Carry, R.A. Tool Steels, 4th Edition, American Society for Metals (1992),

100

pp 223-224 [48] Roberts, G.A.; Grobe, A.H. “Service Failure of Aluminum Die Casting Dies”, Metal Progress, 69 (1956) [18] Schmidt, M.L. “What Really Influences the Mechanical Properties of Premium H-13 Tool Steel?”, Paper G-T89-011, NADCA 15th Die Casting Congress (1989) [41] Seth, B.B. “A Review of Die Casting Dies”, Die Casting Engineer 16 (1972) [13] Sharp, H.J. “Aluminum Pressure Die Casting Dies-Their Failure by Surface Cracking”, Metal Industry 51 (1953) [17] Sinha, A.K. Ferrous Physical Metallurgy, Butterworth Publishers (1989), pp. 523-564 [37] Skelton, R.P. “Introduction to Thermal Shock”, High Temperature Technology 8 No. 2 (1990) [10] Smith, W.F. Structure and Properties of Engineering Alloys, McGraw-Hill (1981), pp. 25-35 [31] Smith, W.F. Structure and Properties of Engineering Alloys, McGraw-Hill (1981), pp. 382-386 [30] Speich, G.R.; Leslie, W.C. “Tempering of Steel”, Metallurgical Transactions 3 (1971), pp. 1043-1054 [34] STERN, M. DIE-CASTING PRACTICE, MCGRAW-HILL BOOK COMPANY (1930) [2]

Swindeman, R.W.; Douglas, D.A. “The Failure of Structural Metals Subjected to Strain Cycling Conditions”, J. Basic Eng. 81 (1959), p. 203 [43] The Making, Shaping and Treating of Steel, Ninth Edition, H.E. McGannon-editor, United States Steel (1971) [39] Toyoda, H. et al. “Thermal Fatigue Properties and Their Evaluation”, Transactions ISIJ 25 (1985) [7] Ule, B.; Vodopivec, F.; Pristavec, M; Gresovnok, F. “Temper Embrittlement of Hot Work Die Steel”, Materials Science and Technology 6 (1990) [27] Wallace, J.F. “Thermal Conditions in the Die”, Foundry 96 (1968) [12]

101

Wang, Y. Effect of Composition and Processing on the Thermal Fatigue and Toughness of High Performance Die Steels, Ph.D. Thesis, Case Western Reserve University (1997) [4] Weronski, A.; Hejwowski, T. Thermal Fatigue of Metals, Marcel Dekker, Inc. (1991), pp.118-124 [22] Winter, M. WebElementsTM Periodic Table [23] Yan, M.; Fan, Z. “Durability of Materials in Molten Aluminum Alloys”, Journal of Materials Science 36 (2001) [1] Zhou, Q. Master Thesis, Case Western Reserve University (2001) [11] Zuchovski, R. “Analysis of the Thermal Fatigue Process”, Journal of Materials Testing and Technology 106 (2000) [9]

102