THE EFFECTS OF PRIOR MICROSTRUCURE ON …MS-2009).pdfTHE EFFECTS OF PRIOR MICROSTRUCURE ON SPHEROIDIZING . KINETICS AND COLD WORKABILITY IN BAR STEELS . by . ... Two different hot-rolled

THE EFFECTS OF PRIOR MICROSTRUCURE ON SPHEROIDIZING

KINETICS AND COLD WORKABILITY IN BAR STEELS

by

R. Allen Schaneman Jr.

A thesis submitted to the Faculty and Board of Trustees of the Colorado School of Mines

in partial fulfillment of the requirements for the degree of Master of Science (Metallurgical and

Materials Engineering).

Golden, Colorado Date ______________

Signed: __________________________________ R. Allen Schaneman Jr.

Signed: __________________________________ Dr. Chester J. Van Tyne

Thesis Advisor

Golden, Colorado Date ______________

Signed: __________________________________

Dr. John J. Moore Professor and Head

Department of Metallurgical and Materials Engineering

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ABSTRACT

Spheroidizing heat treatments are used to soften steel prior of cold forming. Many

automotive parts such as gears, hubs, and universal joint crosses are cold formed utilizing

spheroidization. However, spheroidization heat treatments can last several hours to days and

represent a large investment in time and energy. The time to spheroidize could be shortened by an

understanding of the effects of starting microstructure on spheroidization kinetics and the

resulting cold formability. Two different hot-rolled (bainitic and pearlitic) and one normalized bar

15MnCr5 steel microstructures were heat treated at 692 °C (1277 °F), underwent microstructural

characterization with image analysis software, and were subjected to tensile tests. The bainitic

microstructure spheroidized the fastest, followed by the hot rolled pearlitic and then the

normalized steel. The workability was evaluated with reduction in area values from tensile tests.

Even though it had the lowest percentages of spheroidization, the normalized steel had the highest

reduction in area prior to six hours of heat treatment. The two hot rolled steels had higher

reductions in area after six hours due to higher percentages of spheroidization. The starting

microstructure has a dominant effect on reduction in area, UTS, and upper yield strength

regardless of percent spheroidization at low percentages of spheroidization. However, at high

percentages of spheroidization these properties converge to single values regardless of prior

microstructure. The lower temperature rolled pearlitic structure seems to have the best

combination of heat treatment time to spheroidized and resulting workability.

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TABLE OF CONTENTS

ABSTRACT . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .iii

LIST OF FIGURES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .vii

LIST OF TABLES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .xii

ACKNOWLEDGEMENTS. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .xiii

1. INTRODUCTION. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1

2. LITERATURE REVIEW. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

2.1 Spheroidization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

2.2 Heat Treatments for Spheroidization. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4

2.2.1 Subcritical Heat Treatment. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .4

2.2.2 Intercritical Heat Treatment. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5

2.2.3 Other Heat Treatments. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .6

2.3 Mechanisms. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .8

2.3.1 Raleigh’s Perturbation Theory. . . . . . . . . . . . . . . . . . . . . . . . . . . 8

2.3.2 Mullins and Nichols Modified Perturbation Theory. . . . . . . . . . 9

2.3.3 Thermal Groove Theory. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .10

2.3.4 Fault Migration Theory. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

2.3.5 Multiple Mechanisms Theory. . . . . . . . . . . . . . . . . . . . . . . . . . .11

2.3.6 Ostwald Ripening. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .12

2.4 Kinetic Factors. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

2.4.1 Mechanical Work. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .16

2.4.2 Prior Microstructure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .16

2.4.3 Vacancy Concentration. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .17

2.4.4 Diffusion. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .18

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2.4.5 Other Defects. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19

2.4.6 Kinetic Equations. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

2.5 Workability. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .22

2.5.1 Tension Testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .22

2.5.2 Torsion Testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23

2.5.3 Upset Testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .24

2.5.4 Bend Testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

2.6 Summary. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .27

3. EXPERIMENTAL PROCEDURES. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29

3.1 Material. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .29

3.2 Heat Treatments. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31

3.3 Metallography. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .31

3.4 Microhardness. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .33

3.5 Macrohardness. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 34

3.6 Image J Analysis. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .34

3.7 Compression Testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36

3.8 Tension Testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .37

4. RESULTS . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39

4.1 Heat Treatment. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .39

4.1.1 HR. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39

4.1.2 CR. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 40

4.1.3 Norm. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 46

4.2 Image Analysis. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .47

4.2.1 Particle Area. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47

4.2.2 Percent Spheroidization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50

4.3 Microhardness – Carbide-Rich Regions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .58

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4.4 Microhardness – Ferrite Regions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62

4.5 Macrohardness. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63

4.6 Compression Testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66

4.7 Tension Testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .68

4.7.1 Reduction in Area. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

4.7.2 Uniform Elongation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .71

4.7.3 Total Elongation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72

4.7.4 Ultimate Tensile Strength. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75

4.7.5 Yield Strength. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77

5. DISCUSSION. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 79

5.1 Reduction in Area. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .79

5.2 Total Elongation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .80

5.3 Ultimate Tensile Strength. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .84

5.4 Yield Strength. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .86

5.5 Effects of Initial Microstructure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .87

5.6 Industrial Relevance. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .89

6. SUMMARY. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .91

7. FUTURE WORK . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93

REFERENCES. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .95

APPENDIX A INTERCRICIAL ANNEALING . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 97

APPENDIX B LOGNORMAL STATISTICS. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .99

APPENDIX C COMMERCIALLY SPHEROIDIZED 16MNCR5. . . . . . . . . . . . . . . . . . . .101

APPENDIX D UNIFORM ELONGATION MEASUREMENT. . . . . . . . . . . . . . . . . . . . . 103

APPENDIX E “U” SAMPLE TESING. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .105

APPENDIX F CHARPY TESTING . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .107

APPENDIX G COCKCROFT AND LATHAM FRACTURE CRITERION. . . . . . . . . . . .109

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LIST OF FIGURES

Figure 1.1 A cylindrical steel blank and two different kinds of cold forged universal joint

crosses. The cold forged crosses have close dimensional tolerances and require minimal machining. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .1

Figure 2.1 SEM micrograph of an AISI 4037steel subcritically annealed at 704 °C (1299 °F)

after (a) 4 hours and (b) 12 hours holding. [3]. . . . . . . . . . . . . . . . . . . . . . . . . . . . 5 Figure 2.2 SEM micrograph of an intercritically annealed AISI 4037steel after (a) 4 hours

and (b) 12 hours holding. [3]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6 Figure 2.3 Fracture strain (solid lines) and microhardness (dashed lines) data for a pearlitic

AISI 4037 steel. intercritically annealed () subcritically annealed at 704°C (). [3]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

Figure 2.4 Hardness evolution of time for supercooled austenite at -650°C -700°C and subcritically heat treated Δ-700°C AISI 1045 steel. [4] . . . . . . . . . . . . . . . . . . . . 8 Figure 2.5 Schematic of Raleigh’s perturbation theory for various cylinder lengths.

[7]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12 Figure 2.6 Schematic of Mullins and Nichols modified perturbation theory. (a) carbide plate (b) edges of flat plate thicken due to difference in chemical potential, (c) thickened outer rim develops sinusoidal perturbations, (d) ring breaks up into smaller particles [7]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13 Figure 2.7 Schematic of the thermal groove theory of spheroidization breakup. [7]. . . . . . 14 Figure 2.8 Schematic of the fault migration theory of spheroidization breakup. [7]. . . . . . .15 Figure 2.9 Cylindrical cementite showing characteristics of Raleigh’s perturbation theory. high purity eutectoid steel spheroidized at 700°C for 100 hours. [7]. . . . . . . . . .15 Figure 2.10 Different dynamic strain rates effect on the kinetics of spheroidization for fine pearlite eutectoid steel. ( 700°C 650°C) [8]. . . . . . . . . . . . . . . . . . . . . . 17 Figure 2.11 Spheroidization times for fine, medium, and coarse lamellar spacing in a nearly

eutectoid plain carbon steel annealed at 700°C. [14]. . . . . . . . . . . . . . . . . . . . . . 18 Figure 2.12 The growth of holes and fissures in a cementite plate of a high purity eutectic steel alloy showing a preferred crystallographic orientation. [7]. . . . . . . . . . . . .21 Figure 2.13 Different strain paths possible during upset testing due to different friction conditions and different sample geometery. [20]. . . . . . . . . . . . . . . . . . . . . . . . . 25 Figure 2.14 Figure 2.14 Schematic of the possible sample geometries for upset testing.

(a) cylindrical, (b) tapered, and (c) flanged. [20]. . . . . . . . . . . . . . . . . . . . . . . . 25

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Figure 2.15 Fracture limit diagram for 2024 aluminum alloy with T351 temper. Tests

performed at room temperature and 250°C (480°F) [20]. . . . . . . . . . . . . . . . . . .26 Figure 2.16 Forming limit diagram for materials A (low ductility) and B (high ductility) with

plotted strain paths (a) (high friction) and (b) (low friction) for bolt heading operation. [20]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .27

Figure 3.1 Micrographs of as received material 16MnCr5. (a) HR light optical image,

(b) HR SEM image, (c) CR light optical image, (d) CR SEM image, (e) Norm light optical image, and (f) Norm SEM image (a),(c), and (e) light optical micrograph, picral etch. (b), (d), and (f) SEM images of carbide rich regions, picral etch.. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .32

Figure 3.2 CCT diagrams for the 15MnCr5 steel. (a) austenitized at 870 °C (1600 °F) and

(b) austenitized at 1050 °C (1922 °F). [22] . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33 Figure 3.3 (a) SEM image at 10,000X of CR material after 4 hrs of heat treatment, picral

etch, (b) SEM image after the contrast and threshold have been adjusted in Image J, and (c) the numbered outlines of the analyzed particles analyzed.. . . . . . . . . .36

Figure 3.4 A schematic of the 0.252 in diameter tensile samples used. The samples conform

to ASTM E 8 specifications. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38 Figure 4.1 Micrographs of 16MnCr5 HR conditioned steel after various times at 692 °C

(1277 °F) (light optical micrograghs, picral etch). (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours. . . . . . . . . . . . . . . . . . . . . . . . 42

Figure 4.2 Micrographs of carbide-rich regions in 16MnCr5 HR conditioned steel after

various times at 692 °C (1277 °F) (SEM micrograghs, picral etch). (a) 10 sec, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours. . . . . . . . . . . . .43

Figure 4.3 Micrographs of 16MnCr5 CR conditioned steel after various times at 692 °C


Figure 4.4 Micrographs of carbide-rich regions in 16MnCr5 CR conditioned steel after


Figure 4.5 Micrographs of 16MnCr5 Norm conditioned steel after various times at 692 °C


Figure 4.6 Micrographs of carbide-rich regions in 16MnCr5 Norm conditioned steel after


Figure 4.7 Histograms of the particle area for various heat treatment times for the 15MnCr5

steel in the HR condition. (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .54

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steel in the CR condition. (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .55


steel in the Norm condition. (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .56

Figure 4.10 Changes in average spheroidized particle area during the 692 °C (1277 °F) heat

treatment for the 16MnCr5 steel. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . .57

Figure 4.11 Changes in area percent spheroidized during the 692 °C (1277 °F) heat treatment

for the 16MnCr5 steel. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .59

Figure 4.12 ln(Vc/Vu) with respect to time for the 15MnCr5 steels subcritically spheroidized

at 692 °C (1277 °F). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .60 Figure 4.13 Microhardness in the carbide rich regions for the 16MnCr5 steel after the 692 °C

(1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .61

Figure 4.14 Microhardness in the ferrite regions for the 16MnCr5 steel after the 692 °C


Figure 4.15 Macrohardness of the 16MnCr5 steel after the 692 °C (1277 °F) heat treatment.

(a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .65

Figure 4.16 Compression samples utilizing different stresses, strain rates, sample geometries

and frictional conditions.(a) Compressed to 20 kip, (b) compressed to 60 kip, (c) compressed to limit (0.075 in/min), (d) compressed to limit with roughened ends, (0.075 in/min), (e) compressed to limit (50 in/min), (f) 0.3 in diameter compressed to limit (0.045 in/min), (g) 0.2 in diameter compressed to 54 kip (0.045 in/min), (h) compressed to limit with constrained ends and no talcum powder (0.075 in/min). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67

Figure 4.17 Typical engineering stress-strain curves for the 15MnCr5 steels at various heat

treatment conditions. (a) HR steel as-received, 6 hours, and 20 hours. (b) CR steel as-received, 6 hours, and 20 hours. (c) Norm steel as-received, 6 hours, and 20 hours. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .70

Figure 4.18 Average reduction in area after tensile testing for the 16MnCr5 steel after the

692 °C (1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . .71

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Figure 4.19 Average uniform elongation during a tensile test for the 16MnCr5 steel after the 692 °C (1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . .73

Figure 4.20 Average total elongation for the 16MnCr5 steel after the 692 °C (1277 °F) heat

treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .74

Figure 4.21 Average ultimate tensile strength for the 16MnCr5 steel after the 692 °C


Figure 4.22 Average upper yield strength after a tensile test for the 16MnCr5 steel after the

692 °C (1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . .78

Figure 5.1 Reduction in area and the corresponding percentage of spheroidization for the

16MnCr5 steel for various prior microstructures. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, (d) all microstructures. . . . . . . . . . . . 81

Figure 5.2 Total elongation and the corresponding ferrite microhardness for the 16MnCr5

steel for various prior microstructures. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, (d) all microstructures. . . . . . . . . . . . 83

Figure 5.3 Ultimate tensile strength and the corresponding percent spheroidization for the


Figure 5.4 (a) The relationship between ultimate tensile strength and total elongation for the

15MnCr5 steel. (b) The relationship between ultimate tensile strength and reduction in area for the 15MnCr5 steel. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .86

Figure 5.5 Upper yield strength and the corresponding percent spheroidization for the


Figure 5.6 (a) The relationship between yield strength and total elongation for the 15MnCr5

steel. (b) The relationship between yield strength and reduction in area for the 15MnCr5 steel. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89

Figure A1 SEM micrographs of the 15MnCr5 steel after spheroidization heat treatments.

(a) intercritically anneled HR steel, (b) intercritically annealed CR steel (c) intercritically annealed Norm steel, and (d) subcritically annealed CR steel after 20 hours. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 98

Figure B1 Probability plots for the CR 15MnCr5 steel after one hour of heat treatment at 692 °C (1277 °F) (a) normal distribution, (b) lognormal distribution. . . . . . . . 100

Figure C1 Micrographs of the commercially spheroidized 16MnCr5 Steel, picral etch (a)

light optical micrograph (b) SEM micrograph. . . . . . . . . . . . . . . . . . . . . . . . . .102

x

Figure E1 Photographs of the region of maximum bending for the compressed “U” samples 15MnCr5 steel. (a) compressed at room temperature, (b) compressed at 0 °C (32 °F), and (c) magnified photograph of the room temperature sample showing microcracks. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .106

Figure F1 Photograph of the broken full-size and sub-size charpy samples for the

commercially spheroidized 16MnCr5 steel. . . . . . . . . . . . . . . . . . . . . . . . . . . . 107 Figure F2 Photograph of the fracture surface of the full-size charpy sample of commercially

spheroidized 16MnCr5 steel. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .108 Figure G1 Cockcroft Latham coefficient for the 16MnCR5 steel heat treated at 692 °C

(1277 °F). (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . . . . . . . . .110

xi

LIST OF TABLES

Table 2.1 Activation Energies for Iron and Carbon Diffusion. [13] . . . . . . . . . . . . . . . . . . 19 Table 3.1 Composition, in wt%, of Received 16MnCr5 Steel. . . . . . . . . . . . . . . . . . . . . . . 29 Table 4.1 Average Spheroidized Particle Area (in µm2) and the Experimental Uncertainty

for the 16MnCr5 Steel Heat Treated at 692 °C (1277 °F). . . . . . . . . . . . . . . . . . 58 Table 4.2 Hardness and Experimental Uncertainty in the Carbide-Rich Regions for the

15MnCr5 Steel After Various Heat Treatment Times at 692 °C (1277 °F). . . . .62 Table 4.3 Hardness in the Ferrite Regions for the 15MnCr5 Steel After Various Heat

Treatment Times at 692 °C (1277 °F). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 64 Table 4.4 Macrohardness for the 15MnCr5 Steel After Various Heat Treatment Times at

692 °C (1277 °F). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .66

Table 4.5 Increase in Temperature from Adiabatic Heating during the Tension Test on the 16MnCr5 Steel Tested at a Crosshead Velocity of 495 mm/min (19.5 in/min). .68

Table 4.6 Uniform and Non-Uniform Elongation Values for the 15MnCr5 Steel at Various

Heat Treatment Times at 692 °C (1277 °F). . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75 Table D1 Uniform elongation values for the as-received 16MnCr5 steel using a nominal

load method and Considère’s construction. . . . . . . . . . . . . . . . . . . . . . . . . . . . .103

Table C1 Chemical Composition in wt % of the Commercially Spheroidized 16MnCr5 steel. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .101

Table C2 Image Analysis Results and Tensile Test Data for the Commercially Spheroidized 16MnCr5 Steel. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 102

xii

xiii

ACKNOWLEDGEMENTS

I would like to thank the ASPPRC, all its sponsors, and the Forging Industry Educational

and Research Foundation (FIERF) for funding this research.

I would like to thank Prof. Van Tyne, Prof. Findley, and Dr. Mataya for serving on my

committee and all the guidance they have provided me. I would also like to acknowledge Prof.

Dong-Su Bae and Josh Southworth for thier help in initiating this work. I would also like to thank

Prof. Matlock for his advice and support during this project.

I would like to thank my industrial mentor, Bob Cryderman, and Gerdau MACSTEEL for

supplying me with material, valuable input, and a tour of the Gerdau MACSTEEL bar mill in

Monroe, MI. I would like to thank Mike Burnett form Timken for information on workability and

Dr. Krauss for information on mechanical properties of spheroidized steels.

I would like to acknowledge Dr. Chandler and Gary Zito for their assistance in operating

the SEM. I would like to thank Alex Hudgins for his help with operating the hydraulic press, Joe

Ronevich and Kimani Partin for their help in setting up the tensile testing apparatus, and Kester

Clarke for his help with the Image J analysis software.

I would like to thank my fellow ASPPRC students, especially the night shift, for all their

friendship, support, and good humor. I want to acknowledge Elaine Sutton for her friendship and

aid during my graduate career.

I would also like to give a special thanks to my parents Rodney and Cheryl for all their

love and support, my sister Katryna, for her advice on “Spherical Pineapples,” and my brother

Warnar for his input on how to fracture steel.

A final thank you to all my friends new and old and the countless other people who have

made a difference in my life.

CHAPTER 1

INTRODUCTION

Cold-forming is utilized to produce automotive steel parts such as gears, hubs, bearing

components, and crosses for universal joints. This process produces parts with close dimensional

tolerances and the cold-formed parts require minimal machining after forming. Figure 1.1 shows

a blank and two different types of cold forged crosses for universal joints.

Figure 1.1 A cylindrical steel blank and two different kinds of cold forged universal joint crosses. The cold forged crosses have close dimensional tolerances and require minimal machining after forging.

To increase the ductility of materials prior to cold forming, a spheroidizing heat treatment

is often performed on the steel. During heat treatment the cementite phase in the steel acquires a

spherical morphology. The spherical carbides allow the steel to plastically deform without

cracking. Additionally, the spheroidization heat treatment is used to reduce the applied forces

during forming allowing heavier deformations to be realized. Spheroidizing heat treatments can

last for several hours up to days at temperatures around 700 °C (1300 °F). This process can be

costly in terms of time and energy. The microstructure prior to the heat treatment can affect

1

several factors such as 1) the time needed for spheroidization to occur, 2) cementite morphology,

and 3) the properties of the final spheroidized product. Some producers of spheroidized product

normalize the steel prior to spheroidization which adds extra time and energy consumption to the

manufacturing process. Understanding the effects of prior microstructure and heat treatment

variables on final carbide morphology can increase the efficiency of heat treatment operations

and produce steel of optimal cold formability.

The present study quantifies the changes that develop in three different starting

microstructures after undergoing a spheroidization heat treatment. Two of the starting

microstructures are hot rolled steels and the third steel is in a normalized state. These steels are

subjected to spheroidization heat treatments of varying times to determine, carbide morphology,

extent of spheroidization, and resulting mechanical properties during the spheroidization heat

treatment. These steels were examined with scanning electron microscopy (SEM) and the

cementite was analyzed with image analysis software. Hardness and tensile testing were

performed to determine the resulting mechanical properties.

The potential benefits of this project are:

1. The elimination of the normalizing heat treatment

2. Overall reduction in time for spheroidizing heat treatments

These benefits can only be realized by better understanding the effects of the prior microstructure

on spheroidizing heat treatment and final mechanical properties.

The next chapter will discuss the relevant literature on the mechanisms and kinetics of

spheroidization heat treatments. The different tests used to determine formability will also be

explained. The following chapters will discuss 1) the experimental material and procedures,

2) the major results and findings, 3) a discussion of these results, 4) the summary of the major

findings, and 5) the possible areas to be explored in future work.

2

CHAPTER 2

LITERATURE REVIEW

The spheroidization heat treatment is a common treatment to increase the formability of

steels and decrease the forming forces necessary to shape small components by cold forming.

These parts include gears, bearings, and hubs. There is considerable interest in lowering the costs

involved in spheroidization heat treatments by decreasing the time and energy inputs to heat treat

these steels. Spheroidization can be accomplished through several different heat treatment paths.

This chapter reviews the research which has been done to understand the specific

mechanisms that take place during spheroidization and the kinetics of the spheroidization process.

Decreasing the time to perform spheroidization heat treatments can be aided by understanding the

kinetic factors that enhance diffusion: vacancy concentrations, variation in strain rates, heat

treatment parameters, and prior microstructures. The effectiveness of spheroidization treatments

is often evaluated by one of four different workability tests: the tension test, the torsion test, the

bend test, and the upset test. All of these concepts will be reviewed in this chapter. While many

researchers have evaluated the heat treatment parameters and strain rate effects of spheroidization

heat treatments, few researchers have evaluated the effects of different prior microstructures and

the associated differences in workability.

2.1 Spheroidization

Spheroidization heat treatments decrease the strength and increase the ductility in bar

steels by changing the morphology of the carbide phase to spherical particles. This heat treatment

gives the steel a continuous ferrite matrix, which makes spheroidized steel the most ductile

microstructure possible. [1] The change in carbide morphology is thermodynamically driven by

the decrease in ferrite/carbide interfacial energy. [1]

3

There are two steps to the spheroidization process. In the first step any carbides with high

aspect ratios (such as cementite lamellae in pearlite) are broken into many small, spherical

carbides. The spherical particles have lower surface area to volume ratios than the elongated

structures. These small spherical particles are then coarsened by Ostwald ripening into large

particles thus further decreasing the total surface area to volume ratio. [2] The kinetics of both

stages of spheroidization is controlled by the diffusion of carbon and other alloying elements

through the ferrite or austenite matrix.

2.2 Heat Treatments for Spheroidization

There are four different heat treatment processes for spheroidization. The two most

common and most basic are the subcritical and the intercritical. The other two heat treatments are

variations on the intercritical treatment. [3], [4] More information on each heat treatment is

presented in the following sections.

2.2.1 Subcritical Heat Treatment

To perform subcritical heat treatments, steels are heated to just below their A1

temperature and held for a period of time usually lasting several hours and then cooled to room

temperature. [3], [5] Steels do not undergo the transformation back into austenite and it is

possible to retain elements of the prior microstructure. A fine pearlite structure could thus be

maintained to decrease the diffusion distance and increase the spheroidization kinetics. Structures

of subcritically treated steel usually contain numerous fine spherical carbides inside a ferrite grain

thus providing a high driving force for Ostwald ripening. The final carbide size can then be

carefully controlled by adjusting the heat treatment times and temperatures. Figure 2.1(a) shows

that spheroidization is nearly complete in four hours with numerous small spherical particles in a

pearlitic steel subcritically heat treated at 704 °C (1299 °F). Figure 2.1(b) illustrates how the

particles have coarsened after a 12 hour holding time. [3], [5]

4

(a) (b) Figure 2.1: SEM micrograph of an AISI 4037steel subcritically annealed at 704 °C

(1299 °F) after (a) 4 hours and (b) 12 hours holding. [3] 2.2.2 Intercritical Heat Treatment

In the intercritical heat treatment, steel is heated above its A1 temperature for two to three

hours and then slowly cooled to just below the A1 temperature and held for several hours before

cooling to room temperature. [5] In the intercritical treatment, transformation to austenite will put

the carbon into solution and then slowly cooling the material may nucleate cementite particles

that will coarsen into spheroidized particles. [3] These structures usually feature large cementite

particles precipitated at the grain boundaries. Larger particle sizes can provide enhanced

nucleation sites for micro voids, thus lowering the fracture strength of the steel. [6] Upon cooling

from austenite, the material may also transform into coarse pearlite. Thus, the subcritical holding

time is spent breaking up the coarse pearlite instead of coarsening spherical carbides. [3], [5] The

coarse pearlite structure formed during an intercritical heat treatment is still present in an AISI

4037 steel after 4 hours of heat treatment as shown in Figure 2.2(a). Coarse pearlite has a greater

diffusion distance than fine pearlite and will spheroidize more slowly. Figure 2.2(b) shows the

final spheroidized microstructure after a 12 hour intercritical heat treatment. The intercritical heat

treatment yields a larger spheroidized particle size (Figure 2.2(b)) than the subcritical heat

treatment (Figure 2.1(b)). However, the intercritical treatment may be necessary to obtain

5

spheroidized microstructures in hypereutectic steels in order to put the proeutectoid cementite

into solution. [3]

(a) (b) Figure 2.2: SEM micrograph of an intercritically annealed AISI 4037steel after (a) 4 hours

and (b) 12 hours holding. [3]

The subcritical treatment has also been shown to yield higher fracture strains in hole

expansion tests than the intercritical treatment. [3] Figure 2.3 shows the higher fracture strains

associated with a subcritical treatment. In Figure 2.3 the fracture strain (solid lines) for the

subcritical treatment are almost at a saturation point after two hours. The fracture strain for the

intercritical treatment does not reach the saturation point until approximately 12 hours. Figure 2.3

also shows that that even though the hardness (dashed lines) of the intercritical treatment is lower

than the subcritical treatment the fracture strain of the subcritical steel is still higher. The

difference between the hardness results and the fractures strain results shows hardness does not

seem to directly correlate to the formability of a particular steel.

2.2.3 Other Heat Treatments

In addition to the intercritical and subcritical heat treatments, a spheroidization treatment

involving the decomposition of supercooled austenite has been studied. This treatment starts with

austenite that is quenched to about 330 °C for a few seconds and then reheated to below A1. This

introduces a large number of defects into the lattice. Spherical particles can then form within a

6

few minutes. [4] The resulting microstructure is a very fine network of spherical carbides.

However, the increased number of lattice defects increases the hardness of the steel and decreases

the formability; often making the supercooled austenite process commercially undesirable. Figure

2.4 shows the hardness of two different supercooled austenite treatments compared to a

subcritical treatment. The subcritical treatment has a lower hardness than either of the

supercooled austenite treatments due to the lower concentration of lattice defects. The higher

hardness of the supercooled austenite treatment could lead to undesirable performance during

cold-forming operations.

Figure 2.3 Fracture strain (solid lines) and microhardness (dashed lines) data for a pearlitic AISI 4037 steel, intercritically annealed () subcritically annealed at 704°C (). [3]

The cyclical intercritical heat treatment is another spheroidization heat treatment

involving cycling above and below the A1 temperature. [4] The cyclical intercritical annealing

treatments can also form coarse pearlite which causes spherical particles to form very slowly. [3],

[5]

7

Figure 2.4 Hardness evolution of time for supercooled austenite at -650°C -700°C and subcritically heat treated Δ-700°C AISI 1045 steel. [4] 2.3 Mechanisms

There are four commonly accepted theories for the mechanism of pearlite breakup in

spheroidization: 1) Rayleigh’s perturbation theory, 2) Mullins and Nichols modified perturbation

theory, 3) thermal groove theory, and 4) fault migration theory. [7] However, Raleigh’s

perturbation theory and the thermal groove theory are proposed to have a smaller significance

than the other two theories. [7] Once breakup has occurred and spherical particles exist,

coarsening due to Ostwald ripening becomes dominant. In Ostwald ripening large radius particles

grow at the expense of smaller radius particles. [2], [8]

2.3.1 Raleigh’s Perturbation Theory

Raleigh’s perturbation theory assumes cylindrical-shaped carbide will develop a

sinusoidal perturbation over time as a result of capillary-induced perturbation. These

perturbations will get more and more severe over time if the maximum wavelength, λmax, is

greater than some critical wavelength, λc. The values of these wavelengths depend upon the rate

8

controlling mass transport mechanism. With continual increase of the perturbations, the cylinder

eventually breaks up into a row of spheres with distance equal to that of λmax. [7] Figure 2.5

shows a schematic of the Raleigh’s perturbation theory. Time increases from left to right in

Figure 2.5. At the earliest time on the far left a straight cylindrical particle is represented. As time

progresses, the cylinder develops perturbations with a wavelength of λmax. The cylinder then

breaks into small spherical particles at the far right. Figure 2.5 also shows the effect of the height-

to-diameter ratio on the perturbations. Figure 2.5(a) shows and infinitely long fiber developing

perturbation over the entire length. Figure 2.5(b) shows a particle with a height-to-diameter ratio

under 7.2 breaking up into two spherical particles. Cylinders with a height-to-diameter ratio

greater than 7.2 will develop a perturbation at one end and slowly break up into spherical

particles gradually. The difficulty with Raleigh’s theory in the spheroidization of pearlite is the

assumption of a cylindrical morphology which differs from the plate morphology of actual

pearlite. The plate morphology of pearlite has a large flat surface that is stable against capillary-

induced perturbation. [7] Because the initial carbide morphology of pearlite is plate-like,

Raleigh’s perturbation theory was later modified by W. W. Mullins and F. A. Nichols to integrate

the concept of plate-like morphologies. [7]

2.3.2 Mullins and Nichols Modified Perturbation Theory

Mullins and Nichols modified perturbation theory is sometimes referred to as “edge

spheroidization.” [9]. Mullins and Nichols realized the flat edges of plates are very stable

structures as opposed to the length of a cylinder. The edges of a plate, however, are curved

surfaces which have a higher chemical potential than the flat sides of the plate. [7] The effect of

curvature on chemical potential is shown by the Gibbs-Thompson equation, which is

RTrv

aa m

e

c γ2ln = (2.1)

9

where ac is the activity of the particle/matrix interface, ae is the equilibrium activity, γ is the

interfacial energy, vm is the molar volume of the particle, R is the gas constant, T is temperature,

and r is the particle radius. [10] This difference in chemical potential provides a thermodynamic

driving force for diffusion. The plate will develop a thickened ring around the outer edge of the

plate as a result of the diffusion from the round edge to the outer flat surface of the plate. [8] The

ring can then be assumed to be a curved form of a cylinder which is susceptible to capillary-

induced perturbations. Figure 2.6 shows the breakup of a plate utilizing the Mullins and Nichols

modified perturbation theory. The initial plate-like morphology is shown in Figure 2.6(a). In

Figure 2.6(b) the edges of the flat plate have thickened due to the potential gradient between the

flat and curved surface at the edge. Figure 2.6(c) shows the perturbations developing around the

circumference of the plate. Lastly, the perturbations break into smaller, more spherical particles in

Figure 2.6(d).

2.3.3 Thermal Groove Theory

The thermal groove theory, also called boundary splitting, speculates that cementite

plates break up by a diffusion mechanism along sub-boundaries within the cementite plate. [9]

These sub-boundaries are introduced into the plates in the phase transformation from austenite to

pearlite. These sub-boundaries form a triple point junction between the ferrite and the cementite

boundary. The equilibrium of the surface tension of the triple point junction will form a curved

grain boundary groove in the cementite plate. [7] The curvature in this groove will create a

difference in chemical potential according to the Gibbs-Thompson equation. [10] The difference

in chemical potential causes the diffusion of material out of the groove thus widening the groove.

The groove then widens until the plate breaks up into small particles. Figure 2.7 shows the

breakup of a plate according to the thermal groove theory. [7] The schematic on the far left of

Figure 2.7 shows a plate of cementite with several sub-boundaries within. As time moves forward

to the right the triple junctions of the sub-boundaries start to thicken and separate. The far right

10

schematic in Figure 2.7 shows the final spheroidized carbides resulting from the thickening and

break-up of the sub-boundaries. The creation of boundaries is enhanced by cold working of the

material prior to or during the spheroidization treatment. Regularly spaced dislocations created

during deformation provide short-circuit paths for diffusion. Regularly spaced interface

dislocations have been observed by Chattopadhyay, however, Chattopadhyay observed no

internal dislocation structure in deformed cementite plates to validate sub-boundary breakup. [8]

Alternatively, Tian and Kraft have observed various fringes and structural striations, such as

sequence or stacking faults in the cementite plates. These faults may provide sub-boundaries for

the thermal groove mechanism to occur. [7]

2.3.4 Fault Migration Theory

Fault migration theory assumes a series of staggered plates are considered at the same

time instead of a single plate. In fault migration theory, also called “termination migration,” the

curved end of one plate lies adjacent to the flat surface of another and a chemical potential

gradient is created, thus creating a thermodynamic driving force for diffusion. [9] The curved end

of one plate will recede and the flat side of the other will consequently thicken. Figure 2.8 shows

a schematic of the fault migration theory. [7] In Figure 2.8, two cementite plates are seen next to

each other. The plate on the right terminates before the plate on the right. The flatter surface of

the plate on the left thickens at the expense of the curved end of the plate on the right due to the

difference in chemical potential. Tian and Kraft proposed to extend the theory of fault migration

to not only adjacent lamellae but also to other defects in the cementite lamellae. [7] These defects

are discussed in Section 2.4.6.

2.3.5 Multiple Mechanisms Theory

Tian and Kraft theorized that in the early stages of spheroidization the thermal groove

theory breaks up cracked plates, in the majority of the spheroidization process is dominated by

11

plates thickening and diminishing by holes and fissures in the fault migration model, and at the

end of spheroidization there exist cylindrical rods which breakup due to Raleigh’s perturbation

model. [7] Figure 2.9 shows evidence of these cylindrical rods in high purity eutectic steel.

However, Chattopadhyay suggests that the entire process is controlled by Mullin’s and Nichol’s

modified perturbation theory. [8]

Figure 2.5 Schematic of Raleigh’s perturbation theory for various cylinder lengths. [7] 2.3.6 Ostwald Ripening

After pearlite lamellae have been broken up, the process of Ostwald ripening begins.

Researchers have found Ostwald ripening to be the dominant process occurring after sixty percent

of the cementite had broken up into spheres. [8], [11] Ostwald ripening or coarsening is the

process of growing large spheres at the expense of smaller spheres. The driving force for Ostwald

ripening is the reduction of the total surface energy of a system while maintaining equilibrium

volume fraction of carbides. [12] Small particles have a large surface area to volume ratio

whereas larger particles have smaller ratios. The effect of particle size on chemical activity can

12

again be shown by the Gibbs-Thompson equation given as Equation (2.1). One can use this

equation to calculate the critical particle size where particles will neither coarsen nor dissolve.

Assuming concentration (Ci) is approximately equal to activity (ai) and that the concentration at

the interface Cc is equal to the concentration in the matrix Cm substituting these values into the

Gibbs-Thompson equation yields

crit

m

e

m

RTrv

CC γ2

ln = (2.2)

Solving Equation (2.2) for the critical radius yields

1

ln2

−

⎟⎟⎠

⎞⎜⎜⎝

⎛⎟⎠⎞

⎜⎝⎛=

e

mmcrit C

CRT

vr

γ (2.3)

As temperature increases the equilibrium concentration Ce will increase thereby increasing the

size of the critical radius. [10]

Figure 2.6 Schematic of Mullins and Nichols modified perturbation theory. (a) carbide plate (b) edges of flat plate thicken due to difference in chemical potential, (c) thickened outer rim develops sinusoidal perturbations, (d) ring breaks up into smaller particles [7]

13

Figure 2.7 Schematic of the thermal groove theory of spheroidization breakup. [7]

If the rate of coarsening is considered to be controlled by the reaction at the cementite-

ferrite interface, the Lifshitz-Wagner-Greenwood theory predicts a radius squared dependency as

given by

tRTDkC

rr eo *

9822 Ω

=−γ

(2.4)

where r is the particle radius, ro is the initial particle radius, γ is the interfacial energy, Ω is the

atomic volume of the particle, D is the diffusivity, k is the interface reaction constant, Ce is the

equilibrium carbon concentration in ferrite, R is the gas constant, T is temperature, and t is time.

[2] Atasoy et al. propose that this interface reaction is the rate limiting step for Ostwald ripening

during spheroidization. [11]

If the rate of coarsening assumes a diffusion controlled mechanism, the Lifshitz-Wagner-

Greenwood theory predicts a radius cubed dependency represented by [2]

tRTDC

rr eo *

9833 Ω

=−γ

(2.5)

Many researchers consider diffusion to be the rate controlling mechanism in the Ostwald ripening

process during spheroidization. [2], [8], [13]

14

Figure 2.8 Schematic of the fault migration theory of spheroidization breakup. [7]

Figure 2.9 Cylindrical cementite showing characteristics of Raleigh’s perturbation theory. high purity eutectoid steel spheroidized at 700°C for 100 hours. [7] 2.4 Kinetic Factors

Factors that affect the spheroidization process in steels include: mechanical work, prior

microstructure, vacancy concentrations, and other microstructural defects. Changes in these

parameters affect the diffusion occurring within the steel. Since spheroidization is considered a

15

diffusion controlled process, anything that enhances diffusion will decrease spheroidization heat

treatment times.

2.4.1 Mechanical Work

Steel can be spheroidized heat treated with imposed strain at a specific strain rate.

Increasing the strain rate during the heat treatment can increase the breakup of cementite plates.

Figure 2.10 shows the volume percent spheroidized at different times during heat treatments at

700 °C () and 650 °C (). The different lines on the figure represent different imposed strain

rates during the spheroidization heat treatment. The lines on the far right represent static

annealing. Figure 2.10 shows that the time to spheroidize carbides can be cut by 106 seconds with

strain rates in the range of 1.4 s-1. [8] However, the breakup still remains strongly temperature

dependent and is never completely dependent on strain rate as seen by the two separate

temperature lines at a given strain rate in Figure 2.10. [8] This suggests that spheroidization rate

is only enhanced by strain rate not controlled by it. Increasing strain rate only accelerates the rate

controlling mechanism, indicating that spheroidization is a diffusion controlled process. [8]

Mechanical work can help to increase the diffusion kinetics but does not control the

spheroidization process entirely.

2.4.2 Prior Microstructure

It is commonly accepted that fine pearlite will spheroidize more quickly than a coarse

pearlite. [3], [5], [8], [11], [13], [14], [15] The fine pearlite has a shorter diffusion distance and

therefore will spheroidize more quickly than a coarse pearlite structure. Figure 2.11 indicates

spheroidization times for differently spaced pearlite structures. In Figure 2.11 the fine pearlite

spheroidization for eutectoid steel is complete in just over 300 hours whereas it takes nearly 700

hours to spheroidize the coarse pearlite for the eutectic alloy. [14]

16

Karadeniz compared the spheroidization of an AISI 4140 normalized pearlitic prior

microstructure to that of an AISI 4140 martensitic structure. [15] The martensitic steel

precipitated spherical carbides very quickly and had higher values of fracture strain than the

normalized structure. [15] However, the Brinell hardness values for the fully spheroidized

material showed little difference between the martensitic and the pearlitic structures. [15]

2.4.3 Vacancy Concentration

Vacancies play a critical role in the diffusion of iron. [8] Increasing the iron vacancy

concentration by increasing heat treatment temperature can increase spheroidization kinetics.

Mechanical working of materials undergoing a spheroidization anneal will also increase the

concentration of vacancies. Deformation increases the concentration of iron vacancies and can

accelerate the self-diffusion of iron. [8] The acceleration of iron self-diffusion may increase

spheroidizing kinetics as explained in the next section.

Figure 2.10 Different dynamic strain rates effect on the kinetics of spheroidization for fine pearlite eutectoid steel. ( 700°C 650°C) [8]

17

Figure 2.11 Spheroidization times for fine, medium, and coarse lamellar spacing in a nearly eutectoid plain carbon steel annealed at 700°C. [14]

2.4.4 Diffusion

Spheroidization is a diffusion controlled process, however it is still somewhat unclear as

to which diffusion process is rate controlling. Most researchers claim the diffusion of iron in bulk

iron is the rate controlling mechanism. [2], [13] Others assert that iron diffusion at the iron-

carbon interface is the rate controlling step. [11] The commonality between the researchers is

they all agree that the diffusion of iron is somehow the rate controlling step. Table 2.1 shows the

activation energies of various iron rate controlling steps. [13] Tian et al. found experimental

activation energy values in the range of 210-315 kJ/mole (50-75 kcal/mole) concluding volume

diffusion of iron in an iron being the rate limiting step. [13] However, Atasoy found experimental

activation values around 170 kJ/mole (40 kcal/mole) leading to the assumption that boundary

diffusion of iron is the rate limiting step. [11] Chattopadhyay suggested calculating an effective

diffusion coefficient of carbon that involves the diffusion of iron represented by [8]

FeFeccFe

FeCFe

FeFecceffcc

DxDxV

VV

DxDxDx

+⎟⎟⎠

⎞⎜⎜⎝

⎛ −= 2

3 3 (2.6)

18

where xc and xFe are the mole fractions of carbon and iron respectively, Dc and DFe are the

diffusion coefficients of carbon and iron respectively, and VFe3C and VFe are the molar volumes of

cementite and iron respectively. Equation (2.6) takes into account that the cementite plate must

thicken and the ferrite interface must move to accommodate the new volume of cementite. [8]

Table 2.1 Activation Energies for Iron and Carbon Diffusion. [13]

System Activation Energy kJ/mole (kcal/mole)

Volume Diffusion of Fe in Fe 254-268 (60.7-64.0) Grain Boundary Diffusion of Fe 167-174 (40.0-41.5)

Volume Diffusion of C in Fe 80-84 (19.2-20.1) 2.4.5 Other Defects

There are two major defects in cementite that could contribute to the fault migration

breakup kinetics: 1) kinked or curved lamellae and 2) holes and fissures. Kinked or curved

lamellae can occur from a change in growth direction during the transformation from austenite to

pearlite. During growth of pearlite lamellae, the growth planes can change direction to

accommodate thermodynamic perturbations. This change in direction means they are no longer

growing in their lowest energy habit plane and direction. The lamellae will then gradually turn

back to their habit orientation resulting in kinked and curved lamellae. [7] These kinks are not

only curved surfaces creating a chemical potential difference but also act as nucleation sites for

other defects such as holes and fissures. Holes and fissures in cementite are hard to identify by

conventional metallographic methods because they appear as lamellar terminations in normal

micrographs. Only when the ferrite is completely etched away and three-dimensional cementite

plates remain, can holes and fissures be accurately seen. Figure 2.12 shows a typical hole in a

cementite plate. [7] Notice the curvature around the hole that can create a chemical potential

gradient for diffusion and cementite breakup. From Figure 2.12 one can also see that the hole has

preferred [010] and [120] crystallographic orientation in the cementite lattice.

19

2.4.6 Kinetic Equations

Chattopadhyay et al. attempted to quantify the spheroidization rate using metallographic

measurements of the number of spheroidized particles per unit area (Ns), the mean area of

spheroidized carbides in the section ( sS ), and the mean thickness of unspheroidized particles

( x ). The volume fraction of spheroidized particles is represented by [14]

sss SNV = (2.7)

If spheroidization is considered as the formation of new particles of aspect ratio less than a:1,

then the rate of formation for spherical particles per unit area can be written as [14]

2

1

xadtdV

dtdN ss ≅ (2.8)

However, Equation (2.8) does not take into account the effects of coarsening, so the prediction of

number of spheroidized particles is greatly overestimated. [14]

Atasoy proposed an exponential equation for spheroidization based on the work of

Chattopadhyay and Sellars. Atasoy proposed that since spheroidization involves the pinching off

of lamellae, an additional term be added. The term (C) is defined as the rate of lamellae of aspect

ratio a pinching off per unit time per unit unspheroidized area. C then can be thought to be a

variable that changes with kinetic factors such as faults in the lamellae or pearlite spacing. With

the use of this new term one can define the rate of change in the number of spheroidized particles

per unit area as [11]

us CV

dtdN

= (2.9)

Atasoy combined Equation (2.8) with Equation (2.9) to form [11]

2xaCV

dtdV

us = (2.10)

20

Assuming the rate of increase in volume fraction of spheroidized carbides is inversely

proportional to the rate of disappearance of unspheroidized carbides Vs is replaced with -Vu and

Equation (2.10) is integrated to yield [11]

)exp(2txCaAVV cu −= (2.11)

A represents an integration constant and is dependent on the initial spheroidized volume fraction.

Vc is the volume fraction of total carbides equal to a constant for a given steel alloy and is equal to

the sum of the volume fractions of spheroidized and unspheroidized carbides. [11] Equation

(2.11) can be rearranged in terms of the volume fraction of total carbides over the volume fraction

of unspheroidized carbides as shown by [11]

( )txCaAV

V

u

c 2exp1

= (2.12)

Figure 2.12 The growth of holes and fissures in a cementite plate of a high purity eutectic steel alloy showing a preferred crystallographic orientation. [7]

Atasoy defined the spheroidization rate k as equal to txCa2

. Atasoy plotted values of

ln(Vc/Vu) versus time and measured diffusion activation energy values from k. The activation

21

energy values obtained matched closely with grain boundary diffusion of iron atoms in iron. [11]

However, Atasoy’s model much like that of Chattopadhyay and Sellars does not account for the

coarsening of particles that occurs during spheroidization.

2.5 Workability

Materials undergoing deformation processes such as extrusion or drawing can often crack

due to the limits of the material and the nature of the process used to deform them. These

processes are usually designed by trial and error and operator experience. [16], [17] In order to

better design deformation processes, formability testing has been developed to predict material

behavior in these processes. Formability is defined as “the degree of deformation that can be

achieved in a particular metalworking process without creating an undesirable condition.” [18]

From this definition of formability one can see that formability is not simply a material property

but rather is a function of both the material and the process parameters used. [16], [17], [18] Since

workability depends on process parameters as much as it does material, a variety of different tests

have been developed for formability to simulate different processing conditions. Tension testing,

torsion testing, bending testing, and upset testing are all common tests to evaluate workability.

[16], [17]

2.5.1 Tension Testing

Tension testing is one of the most common types of mechanical tests used to characterize

material. [16] Uniform elongation, total elongation, and reduction in area are common parameters

used to measure the ductility of a material. [16], [17] However, the tension test suffers from many

limitations in the measurement of formability. The amount of deformation that can be imparted in

a tension test is limited by necking instability. [16], [17] Most deformation processes impart far

more deformation than is possible in a uniaxial tension test. [17] In addition, complex biaxial and

triaxial stress states occur during most deformation processes that are not reproduced in a tension

22

test. [16], [17] Necking also makes the true strain rate hard to control in a tension test. [16], [17]

The reduced cross sections of tensile specimens require the surface microstructure to be machined

away. In operations were surface cracking is common, maintaining the surface microstructure is

vital in determining formability. [17]

In addition to measuring the reduction in area and total elongation, Cockcroft and Latham

proposed an additional measure of workability. Cockcroft and Latham measured the plastic strain

energy density and found a material will fracture at a certain plastic strain energy value, C. The

plastic strain energy density is calculated from a tensile test using

∫ =f

Cdε

εσ0

* (2.13)

Where εf is the fracture strain, σ* is the maximum normal stress that is operating. For a tensile

test, σ* is the stress acting at the centerline where the fracture is initiated. [19] This criterion for

fracture has been used successfully, however, the complete true stress-true strain must be known.

This often requires the use of the Bridgman correction. The value of σ* can also be difficult to

calculate but the use of finite element analysis software greatly aides calculations of σ*.

2.5.2 Torsion Testing

Torsion tests avoid many of the problems of the tension test. Torsion tests do not undergo

plastic instability and therefore constant strain rates can be maintained. [16], [17] High strain

rates can be achieved because strain rate is proportional to rotational speed. [16], [17] Torsion

tests can also simulate more complex stress states than a tension test. [17] Torsion tests are also

not limited by friction as in a compression test. [16]

Torsion tests do, however, have complications that make them difficult to use for

formability results. The stresses are not uniform over the entire cross section and vary with the

radius. [17] However, this problem has been overcome with the use of tubular specimens. [16],

[17] The greatest limitation of the torsion test is that the strains that occur during the test are

23

unlike those during most metal working processes. [17] The strains in most deformation

processes are coaxial to the principal stress components; however in torsion testing the principal

stress components are at a 45° angle with respect to the torsion axis. [17] This difference in the

directions of the principal stress and strain components causes the material to develop texture

during the torsion tests. [17] Therefore, the results from a torsion test can often be misleading.

[16], [17] The torsion test also requires the machining of a reduced section and removal of the

surface microstructure. [17]

2.5.3 Upset Testing

Upset testing is the closest thing to a standard workability test. [20] Upset testing

involves the compression of a cylindrical sample and measuring the resulting axial and

circumferential strains. Compression testing does not undergo plastic instability and usually does

not have substantial microstructure reorientation during the test. It has a similar stress state to

most metal forming processes and large amounts of deformation can be achieved before failure.

[16] The stress and strain distributions during an upset test can be changed by altering the friction

conditions and the height-to-diameter ratio of the samples. Increasing the friction or decreasing

the height-to-diameter ratio will increase the non-uniformity of the strain path. [17] Figure 2.13

shows the different stain paths possible by changing friction conditions and sample geometry. In

Figure 2.13 more tensile (circumferential) strain is imparted per unit compressive (axial) strain

with rougher die conditions and lower height-to-diameter ratios. [21]

If a series of upset tests are run at different strain paths to failure a fracture limit diagram

can be formed for a material and process combination. [17] Strain paths can be altered by

changing the friction conditions, the height-to-diameter ratio, or the specimen geometry. Figure

2.14 shows various specimens that can be used: cylindrical, tapered, or flanged. Tapered and

flanged specimens are used to get very high tensile values at small levels of compressive strain.

[21] Care should be used with the flanged and tapered samples however because the tapered and

24

flanged samples do not always follow the same behavior as a cylindrical specimen. Figure 2.15

shows a fracture limit diagram for a 2024 aluminum alloy. [21] Cylindrical, tapered, and flanged

specimens were used to make this forming limit line. Notice how the forming limit line in the

250°C line has a higher slope at the low levels of axial strain. If cylindrical samples had not been

used an incorrect fracture limit curve would have been produced.

Figure 2.13 Different strain paths possible during upset testing due to different friction conditions and different sample geometries. [21]

Figure 2.14 Schematic of the possible sample geometries for upset testing. (a) cylindrical, (b) tapered, and (c) flanged. [21]

The fracture limit diagram can be a useful tool in determining material and process

parameters in a deformation process. If a strain path for a particular process is plotted on a

fracture limit diagram, failures can be predicted. [16], [17], [18], [21] If the plotted strain path lies

below the fracture limit line there will be no failures. If the strain path crosses the fracture limit

line a failure is likely. [16], [17] In Figure 2.16, a fracture limit diagram for materials A (lower

ductility) and B (higher ductility) are shown. [21] Strain paths (a) (high friction) and (b) (low

25

friction) for a bolt heading operation are also plotted. [21] Figure 2.16 predicts that no failures

will occur if material B is deformed with either strain path or if material A is used with strain path

(b). If material A is deformed with strain path (a) fracture is likely. [21]

Figure 2.15 Fracture limit diagram for 2024 aluminum alloy with T351 temper. Tests performed at room temperature and 250°C (480°F) [21]

Although upset testing provides a reasonable approach to workability problems, it does

have one major disadvantage. The load during compression increases sharply with greater

amounts of deformation. [16] The high loads encountered in upset testing limits the axial strains

and sample sizes that can be used.

2.5.4 Bend Testing

Bend tests are useful in situations where the desired sample geometry cannot be obtained

in an upset test or if the material is not cylindrical. [17] Bend testing does not suffer from plastic

instability or microstructural reorientation. [17] The stress and strain states on the outside surface

26

are similar to those in an upset test and can be altered by adjusting the width-to-thickness ratio of

the material. [16], [17] The maximum tensile strain (εθ) in a bend test varies with punch radius

(R) and specimen thickness (t). [17] The maximum tensile strain can be calculated by [17]

( ) ( )[ ]2//ln tRtR ++=θε (2.14)

When performing bend tests it is important to choose a radius of punch and a specimen thickness

so that the fracture strain is less than the maximum possible tensile strain. [17]

Figure 2.16 Forming limit diagram for materials A (low ductility) and B (high ductility) with plotted strain paths (a) (high friction) and (b) (low friction) for bolt heading operation. [21]

2.6 Summary

The relevant literature pertaining to spheroidization mechanisms and kinetics as well as

to workability testing has been reviewed in this chapter. However, the vast majority of the

literature does not combine an analysis of spheroidization with adequate workability

27

characterization. The next chapter will discuss the material and techniques used to quantify

spheroidization as well as cold formability that were used in this present study.

28

CHAPTER 3

EXPERIMENTAL PROCEDURES

This chapter will introduce the starting material as well as its microstructure and

experimental heat treatments. The methods for preparing samples for microscopy and the steps

for image analysis will be detailed. The methods for measuring the mechanical properties i.e.,

hardness and tensile testing, will also be explained in this chapter.

3.1 Material

MACSTEEL supplied twenty-one bars from one heat of 16MnCr5 steel for this project.

15MnCr5 has a similar composition to AISI 5120 steel. Table 3.1 gives the composition of the

steel. The bars were approximately 37.6 mm (1.5 in) in diameter and 915 mm (36 in) in length.

The bars were received in three different starting conditions. Six bars were received in the hot

rolled (HR) condition. These bars were heated to 1125 °C (2057 °F) and finished at 1018 °C

(1977 °F). The bars were then air cooled. Six more bars were received in a hot rolled condition

but were rolled at a lower temperature. This lower temperature condition will be referred to as

“Colder Rolled” (CR) throughout the rest of this thesis. The CR bars where heated to 1080 °C

(1865 °F) and finished rolling at 886 °C (1627 °F). The bars where then air cooled. Lastly, nine

bars were received in the normalized (Norm) condition. These bars were heated to 927 °C

(1700 °F) for two hours and were air cooled at a reduced air cooling rate.

Table 3.1 Composition, in wt%, of As-Received 16MnCr5 Steel.

C Mn P S Si Ni Cr Mo Cu Al 0.18 1.13 0.015 0.025 0.21 0.10 1.05 0.04 0.16 0.029

Figure 3.1 shows the as-received microstructures of the steel in the three conditions. The

HR condition Figure 3.1(a) consists of a bainitic structure with 16% (±2%) proeutectoid ferrite

29

and has some small regions of fine pearlite. The CR condition Figure 3.1(c) is composed of 48%

(±4%) proeutectoid ferrite and fine pearlite. The interlamellar spacing is 0.14 µm (± 0.03 µm).

The Norm condition Figure 3.1(e) consists of 47% (±4%) proeutectoid ferrite and pearlite with a

coarser lamellar spacing than the CR condition; its lamellar spacing is 0.17 µm (± 0.03 µm).

Figure 3.1 also reveals more detail about the carbide regions of the three as received

microstructures. The inner structure of the bainite regions is presented in Figure 3.1(b). The

cementite particles are smaller and have much lower aspect ratios compared to the pearlite

structures of the CR and Norm. Figure 3.1(d) shows that the CR pearlite has a fine interlamellar

spacing. Figure 3.1(f) shows a micrograph of pearlite in the Norm starting condition. The Norm

condition appears to have a coarser interlamellar spacing compared with the CR condition.

The cause for the marked microstructural differences in the HR and CR hot rolled steels

is due to the rolling conditions. Since the HR steel was rolled at a higher temperature, the prior

austenite grain size was larger approximately 130 µm2. The CR steel had a prior austenite grain

size of approximately 30 µm2. This decrease in prior austenite grain size shifted the ferrite and

pearlite transformations forward leading to a ferrite pearlite microstructure in the CR steel. Figure

3.2 shows a continuous cooling transformation (CCT) diagram for the 16MnCr5 steel austenitized

at 870 °C (1600 °F) and 1050 °C (1922 °F). [22] Figure 3.2(a) shows the CCT diagram for the

16MnCr5 steel austenitized at 870 °C. Figure 3.2(b) shows the CCT diagram for the 16MnCr5

steel austenitized at 1050 °C. The prior austenite grain size for the steel austenitized at 1050 °C

diagram can be considered to be larger than the steel austenitized at 870 °C. As the austenitizing

temperature increases (and thus the prior austenite grain size), the ferrite and pearlite

transformations are delayed to longer cooling times and the bainite transformation is accelerated

to shorter transformation times. Thus the HR condition with the larger prior austenite grain size

will transform to ferrite and bainite and the CR steel with the smaller prior austenite grain size

will transform into ferrite and pearlite.

30

3.2 Heat Treatments

19 mm (0.75 in) sections of bar from all three microstructures were subjected to

subcritical heat treatments for varying amounts of time in a Carbolite CWF 12/13 box furnace.

The temperature used was 692 °C (1277 °F), because that temperature is 20 °C (36 °F) below the

Ae1 temperature for this steel, which was calculated using ThermoCalc software. The 20 °C

(36 °F) decrease provides a buffer to keep the steel from transforming to austenite, thus insuring

an entirely subcritical heat treatment. The steel specimens were heated to temperature using a

3.2 °C/min (5.7 °F/min) heating rate. Using this heating rate, the samples reached the holding

temperature in 3.5 hours. The specimens were held for 10 s, 1, 2, 4, 6, 10, and 20 hours at

temperature, then air cooled.

An intercritical heat treatment was also performed on all three prior microstructures to

evaluate the effects of a different heat treatment. The details of the intercritical heat treatment are

discussed in Appendix A.

3.3 Metallography

Cross sections of bar were cut with a LECO CM-24 Model 811-400 liquid-cooled cut-off

saw. These specimens were then heat treated and ground perpendicular to the bar axis on a 180

grit belt grinder to get flat surfaces. The metallographic specimen were ground successively on

240, 320, 400, and 600 grit SiC papers washing with water in-between grinding grits. The

samples were polished with 6 µm and 1 µm diamond suspension. The samples were etched in a

2% picral solution for between 15-25 seconds to reveal the carbide structure.

31

(a) (b)

(c) (d)

(e) (f)

Figure 3.1 Micrographs of as received material 16MnCr5. (a) HR light optical image, (b) HR SEM image, (c) CR light optical image, (d) CR SEM image, (e) Norm light optical image, and (f) Norm SEM image (a),(c), and (e) light optical micrograph, picral etch. (b), (d), and (f) SEM images of carbide rich regions, picral etch.

32

Figure 3.2 CCT diagrams for the 15MnCr5 steel. (a) austenitized at 870 °C (1600 °F) and (b) austenitized at 1050 °C (1922 °F). [22]

3.4 Microhardness

Both the carbide regions and the ferrite regions were microhardness tested using a

Vickers indenter. The carbide regions of the HR and CR were subjected to a 50 g load; however,

the Norm samples had to be tested with a 25 g load due to the smaller size of the carbide regions.

The ferrite regions were tested with a 25 g load due to the softness of the phase and the small

33

size. All microhardness tests performed had a 10 s holding time. Each measurement is the result

of twelve hardness measurements. The two highest and two lowest measurements were

eliminated making each measurement the average of eight hardness measurements.

3.5 Macrohardness

Macrohardness performed on each sample was performed using the Rockwell B scale

with a 1/16 in ball indenter. Macrohardness measurements were taken at six different locations

around the mid-radius. The highest and lowest measurements were eliminated making the

measurement an average of four hardness indentations.

3.6 Image J Analysis

SEM micrographs were analyzed utilizing the Image J image analysis program. Image J

can gather information about particle size and morphology by utilizing the color difference in the

pixels of an image. This program was used to get information on the area and aspect ratio of the

cementite particles. The following steps were taken to analyze the images:

1. Black and white micrographs are opened in an eight bit format with the Image J program.

2. The contrast on the image was adjusted to enhance the color difference between the

carbides and the ferrite.

3. The “Smooth” command was used to eliminate any background static in the image.

4. The “Threshold” command was used to highlight the cementite particles in a red color.

All red areas were measured as particles.

5. The image scales was set by drawing a line of known length on the image and converting

its length to a pixel length in the “Set Scale” dialogue box. For example, the scale on the

10,000X images used for the majority of the work is 75 pixels per micrometer.

6. In the “Set Measurements” dialogue box, the measurements to be taken were set. The

“Area” and the “Fit Ellipse” were selected.

34

7. The micrograph being analyzed was compared side-by-side with the original image to

make sure no two particles had been combined. If particles were joined black lines were

drawn between the particles to separate them.

8. A check was made to insure the entire area of the particle was filled with the red pixels. If

not, white pixels were filled over the necessary areas of cementite.

9. When all the particles were separated and filled, a rectangle was drawn around the area to

be analyzed.

10. The “analyze particles” dialogue box was activated and the size range of particles to be

analyzed was entered. For the 10,000X micrographs analyzed, all particles between

0.001 µm2 and infinity were included. This means a particle was at least 3 pixels for it to

be included in the analysis. A three pixel particle can be seen on the micrographs. The 3

pixel area corresponds to a 0.040 µm minimum particle diameter. The purpose of this

step was to exclude any background noise in the analysis.

11. The “Outlines” option was chosen in the “Show” pull-down menu to provide a numbered

map of all the particles analyzed. The “Exclude Edge Particles” option was also chosen

so any partial particles would be ignored.

Figure 3.2 shows the progression of a micrograph through the image analysis software. Figure

3.2(a) is the starting micrograph. Figure 3.2(b) shows the micrograph after the contrast has been

adjusted and the threshold has been added. Figure 3.2(c) is the numbered outline of all the

particles analyzed. For most of the CR material, images were analyzed until 10,000 particles

were measured. Due to the long time span of the analysis, only 5,000 particles were analyzed for

the HR and Norm materials. Particles were considered spheroidized when an aspect ratio of 3:1 or

less was reached. [5] All particle counts were then subjected to lognormal statistical analysis.

Appendix B explains reasoning behind the choice of lognormal statistical analysis.

35

(a) (b)

(c) Figure 3.3 (a) SEM image at 10,000X of CR material after 4 hrs of heat treatment, picral etch, (b) SEM image after the contrast and threshold have been adjusted in Image J, and (c) the numbered outlines of the analyzed particles analyzed.

3.7 Compression Testing

Compression testing was chosen because it is the most common test for evaluating cold

workability. [18] Cylindrical compression samples were machined down from a 35.5 mm (1.4 in)

bar of commercially spheroidized 16MnCr5 steel provided by MACSTEEL. More information on

the commercially spheroidized 16MnCr5 steel is provided in Appendix C. The samples were

machined to a variety of diameters from 5.1 mm (0.2 in) to 12.7 mm (0.5 in). The samples all had

a height-to-diameter ratio of 1.5:1 and had #16 surface finish. These samples were compressed on

a hydraulic press with a 445 kN (100 kip) limit. Samples were compressed with a variety of

36

friction conditions. Samples were compressed with talcum powder lubricant, roughened sample

ends, and constrained sample ends representing a sticking friction condition. The sample

geometry before and after testing was measured and the presence of surface cracks was evaluated.

Most tests were performed at an engineering strain rate of 0.1 min-1; however, to enhance the

probability of cracking, some samples were tested at a crosshead velocity of 1270 mm/min

(50 in/min).

3.8 Tension Testing

Tension testing was chosen to evaluate the cold workability due to the lack of failures

achieved by compression and bend testing. Samples for tension testing were taken from the mid

radius of the bar and machined according to ASTM E 8. [23] The samples had a 6.4 mm

(0.252 in) diameter and a 31.8 mm (1.25 in) gauge length. Figure 3.3 shows the complete sample

geometry for tensile testing. Testing was performed on an Instru-met A30-33 frame with an

Instron 89,000 N (20,000 lb) load cell. Testing was performed at a rate of 500 mm/min

(19.5 in/min). A 25.4 mm (1.0 in) Shepic extensometer with 12.7 mm (0.5 in) of extension was

used to measure the elongation in each sample. The yield strength and ultimate tensile strength

were all obtained from the load cell and extensometer data. Uniform elongation was measured as

the strain coinciding with the peak stress. Comparisons between the values measured utilizing this

method and by using the Considère’s construction are shown in Appendix D. The differences

were negligible so the former method was chosen. Reduction in area was calculated by measuring

the gauge diameter before and after breaking with a pair of dial calipers accurate within

0.025 mm (0.001 in). The total elongation was measured utilizing the 25.4 mm (1.0 in)

indentations left by the extensometer using a dial caliper. According to ASTM E 8, percent

elongation and reduction in area values should only be measured on samples in which break in

the middle-half of the gauge section. [23] It should be noted that the majority of the samples used

37

in this testing broke just outside of the middle-half of the gauge section and were still included in

the percent elongation and reduction in area measurements.

Figure 3.4 A schematic of the 0.252 in diameter tensile samples used. The samples conform to ASTM E 8 specifications.

38

CHAPTER 4

RESULTS

This chapter presents the results of the spheroidized microstructure analysis. The

microhardness and macrohardness results are presented and discussed. This chapter also provides

the results of the mechanical testing.

4.1 Heat Treatment

In order to evaluate the progression of spheroidization, samples of each initial

microstructure were heat treated for a variety of times ranging from ten seconds to twenty hours.

These heat treatments were performed at 692 °C (1277 °F).

4.1.1 HR

Before heat treatment, the HR steel was comprised mostly of proeutectoid ferrite and

bainite with some small regions of pearlite. Figure 4.1 shows the photomicrographs that monitor

the progression of spheroidization as a function of time for the HR material. Figure 4.1(a) shows

the microstructure for the HR steel after ten seconds of holding time at 692 °C (1277 °F). The

carbide-rich regions show very little evidence of spheroidization taking place and there is a clear

delineation between the carbide regions and the bainitic ferrite. Figure 4.1(b) shows the HR steel

after one hour at temperature. The carbides are starting to show a spheroidized structure and the

morphology of the bainitic ferrite can still be seen. Figure 4.1(c) shows the structure of the HR

steel after two hours of heat treatment. The structure appears very similar to that of the one hour

sample (Figure 4.1(b)). Figure 4.1 (d) shows the microstructure of the HR four hour sample. The

spherical carbides have a larger spacing than in previous times and the boundaries around the

bainitic ferrite are becoming harder to discern. Carbides are growing between the boundaries of

39

the proeutectoid ferrite. Figure 4.1(e) shows the microstructure of the HR steel after six hours of

heat treatment. The carbide size has increased and so has the spacing between carbides. The

boundaries of the bainitic ferrite have almost disappeared. Carbides are still growing between the

proeutectoid ferrite. Figure 4.1(f) shows the microstructure of the HR steel after twenty hours of

heat treatment. The spherical carbides have grown even larger and the space between particles

has increased even more. The boundaries of the bainitic ferrite are only faintly visible. Grain

boundary carbides are also visible between the proeutectoid ferrite boundaries.

The carbides of the HR steel consisted of short particles with low aspect ratios. The HR

steel did not have to undergo much heat treatment time to reach a large distribution of spherical

carbides. Figure 4.2 shows SEM micrographs of the carbide morphology during the heat

treatment. Figure 4.2(a) shows the HR carbides after ten seconds of heat treatment. The carbides

are small and are a mixture of spherical carbides and elongated carbides. Figure 4.2(b) shows the

HR carbides after one hour of heat treatment. The carbides have grown larger and more of the

carbides have become spherical even though some elongated particles remain. Figure 4.2(c)

shows the carbide morphology after two hours. Almost all the carbides have attained a spherical

shape and have grown larger. Figure 4.2(d), (e), and (f) show the HR carbide morphology after

four, six, and twenty hours of heat treatment respectively. The spherical carbides have grown

larger and the number density of carbides has decreased due to Ostwald ripening.

4.1.2 CR

Before heat treatment, the CR steel consisted of fine pearlite and proeutectoid ferrite.

Figure 4.3 shows the microstructural evolution of the CR material during the heat treatment.

Figure 4.3(a) shows the CR steel after ten seconds of holding time at 692 °C (1277 °F). Dense

areas of spherical carbides can be seen in the former pearlite colonies. Figure 4.3(b) shows the

CR steel after one hour of heat treatment. The regions of spherical carbides have become less

dense. Figure 4.3(c) shows the CR steel microstructure after two hours of heat treatment. The

40

microstructure appears very similar to that of the one hour treatment (Figure 4.3(b)). Figure

4.3(d) shows the microstructure of the CR steel after four hours of heat treatment. The spherical

particles have coarsened and are less dense in the former pearlite colonies. Grain boundary

carbides are seen growing along the proeutectoid ferrite grain boundaries. Figure 4.3(e) shows the

microstructure of the CR steel after six hours of heat treatment. The spherical carbides have

grown larger and are less dense than before. The grain boundary carbides continued to grow and

spread in between the proeutectoid ferrite. Figure 4.2(f) shows the microstructure of the CR steel

after twenty hours of heat treatment. The spherical carbides have further coarsened and the

carbides are more widely spaced in the former pearlite colonies. The grain boundary carbides

surround the proeutectoid ferrite.

The initial carbide structure of the CR steel was fine pearlite. The spheroidization heat

treatment breaks up the cementite lamellae into spherical particles. Figure 4.4(a) shows the

carbide morphology of the CR steel after ten seconds of heat treatment. The former pearlite is a

mixture of small spheres and elongated carbides. Figure 4.4(b) shows the carbides of the CR steel

after one hour of heat treatment. Spheroidization break up has taken place and the spherical

particles have coarsened. However, elongated carbides still remain. Figure 4.4(c) shows the CR

steel after two hours of heat treatment. Most of the elongated carbides have broken into spherical

particles; however, a small amount of elongated carbides remain. Figure 4.4(d) shows the carbide

morphology of the CR steel after four hours of heat treatment. The spherical carbides have

coarsened and the density of carbides has decreased. The morphology of the carbides is almost

completely spherical. Figure 4.4(e) shows the carbide morphology of the CR steel after six hours

of heat treatment. The carbides are spherical and the space between them increased. The size has

also increased compared to the previous times. Figure 4.4(f) shows the carbide morphology of

the CR steel after twenty hours of heat treatment. The spherical carbides have coarsened to a

larger size and the density of the carbides has decreased.

41

42

(a) (b)

(c) (d)

(e) (f)

Figure 4.1 Micrographs of 16MnCr5 HR conditioned steel after various times at 692 °C (1277 °F) (light optical micrographs, picral etch). (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours.

(a) (b)

(c) (d)

(e) (f) Figure 4.2 Micrographs of carbide-rich regions in 16MnCr5 HR conditioned steel after

various times at 692 °C (1277 °F) (SEM micrographs, picral etch). (a) 10 sec, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours.

43

(a) (b)

(c) (d)

(e) (f)

Figure 4.3 Micrographs of 16MnCr5 CR conditioned steel after various times at 692 °C (1277 °F) (light optical micrographs, picral etch). (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours.

44

45

(a) (b)

(c) (d)

(e) (f) Figure 4.4 Micrographs of carbide-rich regions in 16MnCr5 CR conditioned steel after


4.1.3 Norm

The Norm steel had an initial microstructure of coarse pearlite and proeutectoid ferrite.

Figure 4.5(a) shows the microstructure of the Norm steel after ten seconds. Little evidence of

spheroidization can be seen in the carbide regions but some small grain boundary carbides can be

seen in grain boundaries of the proeutectoid ferrite. Figure 4.5(b) shows the microstructure of the

Norm steel after one hour of heat treatment. The former pearlite colony regions show evidence of

the carbides beginning to spheroidize. The grain boundary carbides between the proeutectoid

ferrite have increased in size. Figure 4.5(c) shows the microstructure of the Norm steel after two

hours of heat treatment. The microstructure appears similar to that of the one hour sample (Figure

4.5(b)); however, carbides are beginning to grow between the proeutectoid ferrite grains. Figure

4.5(d) shows the microstructure of the Norm steel after four hours of heat treatment. The

spheroidized nature of the carbides has become more evident. Figure 4.5(e) shows the Norm steel

after six hours of heat treatment. The spheroidized carbides have increased in size and the space

between carbides has increased. The grain boundary carbides between the proeutectoid ferrite

grains have also grown larger. Figure 4.5(f) shows the Norm steel after twenty hours of heat

treatment. The space between the carbides has increased after twenty hours. The grain boundary

carbides between the proeutectoid ferrite have also coarsened.

The carbides in the Norm steel consisted of coarsely spaced pearlite lamellae prior to heat

treatment. Figure 4.6(a) shows the carbide morphology after ten seconds of heat treatment. The

majority of the carbides still have an elongated morphology; however, some of the original

lamellae have broken up into small spherical carbides. Figure 4.6(b) shows the carbide

morphology of the Norm steel after a one hour heat treatment. The lamellae have broken into

smaller carbides but these carbides still have aspect ratios too high to classify them as spherical.

Figure 4.6(c) shows the carbide morphology of the Norm steel after two hours of heat treatment.

Some of the carbides have reached the spherical morphology but still many remain as elongated

particles. Figure 4.6(d) shows the carbide morphology of the Norm steel after four hours of heat

46

treatment. The majority of the carbides have reached the spherical morphology but elongated

carbides are still present in the microstructure. Figure 4.6(e) shows the carbide morphology of the

Norm steel after six hours of heat treatment. The spherical carbides have coarsened but some

elongated carbides still remain. Figure 4.6(f) shows the carbide morphology after twenty hours of

heat treatment. The spherical carbides have coarsened further and the distance between carbides

has increased. The elongated carbides still have not completely disappeared after twenty hours of

heat treatment in the Norm steel.

4.2 Image Analysis

The SEM images of the carbide regions were analyzed using the ImageJ image analysis

software. This software was used to gather information on carbide area and shape. Images were

analyzed until at least 5,000 carbides were analyzed for each heat treatment time and

microstructural condition.

4.2.1 Particle Area

Particle area can be used to look at the size of the initial spherical particles formed during

spheroidization and to quantify the effects of coarsening after the formation of spherical carbides.

A spherical particle is considered t have an aspect ratio of less than 3:1. Figure 4.7 shows the

distribution of spheroidized particle areas during the heat treatment time for the HR steel. Figure

4.7(a) shows the distribution of spheroidized particle sizes after ten seconds of heat treatment for

the HR steel. The distribution has a lognormal shape and the largest group of spheroidized

particles is in the range of 0.010 µm2. Figure 4.7(b) shows the distribution of spheroidized

particle sizes for the HR steel after one hour of heat treatment. The spheroidized particle size has

increased and the largest group is approximately 0.015-0.020 µm2. Figure 4.7(c) shows the

distribution of spheroidized particle sizes in the HR steel after two hours of heat treatment. The

distribution of particle sizes has increased thus increasing the average particle size to over

47

0.020 µm2. Figure 4.7(d), (e), and (f) show the spheroidized particle area distribution for the HR

steel after four, six, and twenty hours respectively. The frequency of larger particle sizes is

increasing at the expense of the smaller particle sizes. This is the result of coarsening. The

distribution in Figure 4.7(f) shows the wide variety of particle sizes that can commonly be seen

throughout the microstructure. This type of distribution causes a concern for the use of average

particle size as a quantitative characteristic for evaluating spheroidization.

Figure 4.8 shows the distributions of spheroidized particle area for the CR steel at various

times during heat treatment. Figure 4.8(a) shows the spheroidized particle area distribution for the

CR steel after ten seconds of heat treatment. Most of the particles are very small and are between

0.005-0.010 µm2. The distribution also has the lognormal trend to the data. Figure 4.8(b) shows

the spheroidized particle area distribution for the CR steel after one hour of heat treatment. The

average particle size has grown to approximately 0.010 µm2. Figure 4.8(c) shows the

spheroidized particle area distribution for the CR steel after two hours of heat treatment. The

distribution has not changed significantly from the one hour distribution. This may be due to new

spheroidized particles being formed at the same rate that previous particles are coarsening. Figure

4.8(d) shows the spheroidized particle area distribution for the CR steel after four hours of heat

treatment. The range of sizes has increased and the average size has increased to approximately

0.020 µm2 due to the coarsening of particles. Figure 4.8(e) and (f) show the spheroidized particle

distribution for the CR steel after six hours and twenty hours of heat treatment. The number of

large particles has increased and the number of small particles has decreased due to coarsening.

Figure 4.9 shows the distribution of spheroidized particle areas for the Norm steel at

various times during heat treatment. The ten second heat treatment for the Norm steel could not

be analyzed due to the predominant long cementite lamellae occurring in the microstructure.

Figure 4.9(a) shows the distribution of spheroidized particle areas for the Norm steel after one

hour of heat treatment. The average particle size is between 0.010-0.015 µm2 with approximately

30% of the particles being in this range. Figure 4.9(b) shows the spheroidized particle area

48

distribution in the Norm steel after two hours of heat treatment. The average particle size is still

between 0.010-0.015 µm2 however the percentage of the particles in this range has increased to

approximately 40%. This reflects the increase in newly spheroidized particles from the breakup of

cementite lamellae occurring in between the first and second hour. Figure 4.9(c) shows the

distribution of spheroidized particle areas for the Norm steel after four hours of heat treatment.

The distribution looks similar to the two hour distribution showing that particles are coarsening

and new spheroidized particles are being formed. Figures 4.9(d), (e), and (f) show the distribution

of spheroidized particle areas for the Norm steel after six, ten, and twenty hours respectively. The

number of particles in the range of 0.010-0.015 µm2 is decreasing and the frequency of larger

particles is increasing showing evidence of coarsening becoming a dominant process.

In order to better visualize the changes in spheroidized particle size, the average

spheroidized particle size was calculated using lognormal statistical analysis. Figure 4.10 shows

the changes in average spheroidized particle area for each of the prior microstructures. Figure

4.10(a) shows the increasing particle size for the HR steel. The average spheroidized particle size

was 0.015 µm2 after ten seconds; however, after ten hours the average particle size grew to

0.033 µm2. This growth was the result of coarsening taking place during the spheroidization heat

treatment. After ten hours the average particle size appears to reach its maximum size and the

same average particle size is seen after twenty hours. Figure 4.10(b) shows the average

spheroidized particle size for the CR steel. The average particle size was 0.008 µm2 at ten seconds

and coarsened to 0.027 µm2 after twenty hours. Figure 4.10(c) shows the average spheroidized

particle size for the Norm steel. After one hour, the average particle size is 0.015 µm2 and

coarsens to 0.020 µm2 after twenty hours.

Figure 4.10(d) shows the average spheroidized particle area for all three steels. The CR

steel has a larger diameter than the Norm steel after twenty hours because the finer interlamellar

spacing decreases the carbon diffusion distance making larger spheroidized particle sizes

possible. The HR steel had the largest spheroidized particle size. The reason is probably because

49

HR steel did not have to undergo the extensive spheroidization breakup that was needed for the

pearlitic steels. Thus, the HR steel had more time to coarsen. The rate of increase in particle size

was also affected by the spheroidization breakup. The HR material had to undergo little breakup

and had the fastest increase in average spheroidized particle size. The CR and Norm steels had to

breakup into spheres before the spheroidized particles could increase in size.

The experimental uncertainty, calculated using lognormal statistics, on the spheroidized

particle size measurements was rather large. This was due to the variety of particle sizes resulting

from the simultaneous processes of breakup, spheroidization, and coarsening. These processes

created a wide range of particle sizes and made an average size not only difficult to measure but

also a rather poor quantitative measure of the spheroidization process. Table 4.1 shows the

average particle size measure for each of the three steels and the experimental uncertainty. The

uncertainty grows larger for the longer times because coarsening has taken place and created an

even broader range of particle sizes.

4.2.2 Percent Spheroidization

The image analysis software not only analyzes the size of the particles but also the shape.

Hosford et al. define a spheroidized particle as one that has an aspect ratio of less than 3:1. [5]

Using this criterion an area percentage of spheroidization can be calculated.

100*%T

ST

AAA

edSpheroidiz−

= (4.1)

where AT is the total area of all particles and AS is the area of all spheroidized particles. Figure

4.11 shows the area percent spheroidized for the three steels. The area percent spheroidized

follow a logarithmic trend, increasing rapidly then equilibrating around 90% spheroidization.

Figure 4.11(a) shows the area percent spheroidized for the HR steel. The HR steel reaches 90%

spheroidization in approximately two hours of heat treatment. The steel remains around 90%

spheroidized for the rest of the heat treatment times. Figure 4.11(b) shows the area percent

50

spheroidized for the CR steel. The CR steel takes approximately ten hours to reach 90%

spheroidization and is approximately 93% spheroidized after twenty hours of heat treatment.

Figure 4.1(c) shows the area percent spheroidization for the Norm Steel. The Norm steel is has

only reached 84% spheroidization after twenty hours. Figure 4.11(d) compares the area percent

spheroidized for all the microstructures. The structures all reach relatively high percentages of

spheroidization after ten seconds of holding time; however, this time includes a 3.5 hour heating

time and an air cooling time from 692 °C (1277 °F). The relatively larger lamellar spacing of the

Norm steel makes it slower to spheroidize than the coarse spacing in the CR steel. The small

aspect ratios of the bainitic carbides give the HR steel the fastest spheroidization response.

Because the HR and CR steels are both hot rolled steels, the carbides in them likely have more

defects such as cracks and kinks to enhance the spheroidization kinetics.

The kinetics of the spheroidization process can be quantified utilizing Atasoy’s

exponential equation for spheroidization (Equation 2.12). If the values of ln(Vc/\Vu) are plotted

with respect to time, the slope of the resulting line is the spheroidization rate k. This model was

shown to overestimate spheroidization since it does not account for Ostwald ripening. This model

may be used to estimate the spheroidization rates at shorter times when Ostwald ripening is a less

dominant process. Figure 4.12 shows the plot of ln(Vc/Vu) versus time for all three steel

microstructures. The k values for the HR and CR steels are similar: 3.9*10-3 s-1 and 3.5*10-3 s-1

respectively. The spheroidization rate for the Norm steel was 1.9*10-3s-1. These values for

spheroidization rate agree well with trends shown in Figure 4.11. The HR steel spheroidized at

the highest rate followed closely by the CR steel. Both the HR and the CR steel are hot rolled

steels and have carbide structures with many defects such as cracks and kinks that will accelerate

the spheroidization process. The Norm steel spheroidized more slowly because it spheroidized

from coarse pearlite that probably did not contain as many carbide defects to enhance the

spheroidization rate.

51

(a) (b)

(c) (d)

(e) (f)

Figure 4.5 Micrographs of 16MnCr5 Norm conditioned steel after various times at 692 °C (1277 °F) (light optical micrographs, picral etch). (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours.

52

53

(a) (b)

(c) (d)

(e) (f) Figure 4.6 Micrographs of carbide-rich regions in 16MnCr5 Norm conditioned steel after


0 0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4Spheroidized Particle Area (µm2)

0

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0

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2000F

tal (

%)

pheroidized Particle Area (µm2)

Perc

enta

ge o

f To

requ

ency

(e) (f) Figure 4.7 Histograms of the particle area for various heat treatment times for the 15MnCr5

steel in the HR condition. (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours.

54


0

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steel in the CR condition. (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours.

55


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steel in the Norm condition. (a) 10 seconds, (b) 1 hour, (c) 2 hours, (d) 4 hours, (e) 6 hours, and (f) 20 hours.

56

0 4 8 12 16 20Holding Time (Hours)

0

0.01

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roid

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roid

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idiz

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2 )

MicrostructuresHRCRNorm

(c) (d)

Figure 4. 10 Changes in average spheroidized particle area during the 692 °C (1277 °F) heat treatment for the 16MnCr5 steel. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

57

Table 4.1 Average Spheroidized Particle Area (in µm2) and the Experimental Uncertainty for the 16MnCr5 Steel Heat Treated at 692 °C (1277 °F).

HR CR Norm Holding

Time (Hours)

Avg. Area -(µm2) +(µm2) Avg.

Area -(µm2) +(µm2) Avg. Area -(µm2) +(µm2)

0.003 0.015 0.009 0.024 0.008 0.005 0.012 1 0.014 0.005 0.040 0.012 0.008 0.022 0.015 0.009 0.020 2 0.023 0.015 0.047 0.010 0.006 0.015 0.011 0.006 0.012 4 0.018 0.013 0.040 0.019 0.013 0.044 0.012 0.007 0.017 6 0.020 0.013 0.043 0.017 0.011 0.033 0.016 0.010 0.029

10 0.033 0.022 0.061 0.018 0.012 0.034 0.014 0.009 0.022 20 0.034 0.022 0.066 0.027 0.019 0.059 0.020 0.013 0.038

4.3 Microhardness – Carbide-Rich Regions

The microhardness was tested in the carbide-rich regions after the various heat treatment

times. The decrease in hardness in the carbide regions is related to the spheroidization

phenomenon taking place. Figure 4.13 shows the microhardness in the carbide-rich regions for

the three steels. Figure 4.13(a) shows the carbide region microhardness for the HR steel. The HR

steel decreases sharply for the first two hours then steadily decreases in hardness until reaching

the minimum hardness of 175 HV at twenty hours. Figure 4.13(b) shows the carbide-rich region

microhardness for the CR steel. The microhardness of the CR steel decreases at a steady rate

before reaching the minimum hardness of 175 HV at approximately ten hours. Figure 4.13(c)

shows the carbide-rich region microhardness for the Norm steel. The microhardness of the Norm

steel decreases steadily until it reaches the minimum hardness of 175 HV in approximately five

hours. Figure 4.13(d) shows the carbide-rich region microhardness for all the microstructures.

The Norm structure has the highest starting hardness; which corresponds to the almost completely

lamellar shape of the carbides after ten seconds of heat treatment compared to the partially

spheroidized structures of the HR and CR steels. However the Norm steel reaches the minimum

hardness before the CR and HR steels which does not correspond well to the spheroidization

58

results. Also the HR material retains high hardness even though it has the largest spheroidized

carbides and the largest area percentage spheroidized. This indicates that something other than the

spheroidization is controlling the carbide region microhardness. This phenomenon is not fully

understood; however, it may be related to the varying size of the carbide regions among the prior

microstructures.


0

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0Holding Time (Hours)

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(c) (d)

Figure 4. 11 Changes in area percent spheroidized during the 692 °C (1277 °F) heat treatment for the 16MnCr5 steel. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

59

0 400 800 1200Holding Time (Seconds)

0

1

2

3ln

(Vc/

Vu)

MicrostructureHRCRNorm

Figure 4.12 ln(Vc/Vu) with respect to time for the 15MnCr5 steels subcritically spheroidized at 692 °C (1277 °F)

60

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Ca

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RB

)

(a) (b)

160

180

200

220

240

260

Ca

0 4 8 12 16 20

Holding Ti

160

180

200

220

240

260

Ca

me (Hours)

rbid

e M

icro

hard

ness

(VH

N)

75

80

85

90

95

Har

dnes

s (H

RB

)

MicrostructuresHR

0 4 8 12 16 20

Holding Ti

75

80

85

90

95

Ha

CR

me (Hours)

rbid

e M

icro

hard

ness

(VH

N)

Norm

rdne

ss (H

RB

)

(c) (d) Figure 4. 13 Microhardness in the carbide rich regions for the 16MnCr5 steel after the

692 °C (1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

Table 4.2 shows the measured values for the carbide microhardness and the experimental

uncertainty. The average standard deviation in the measurements was ±7 HV.

61

Table 4.2 Hardness and Experimental Uncertainty in the Carbide-Rich Regions for the 15MnCr5 Steel After Various Heat Treatment Times at 692 °C (1277 °F).

HR CR Norm Holding

Time (Hours)

Hardness (HV)

Std Dev (HV)

Hardness (HV)

Std Dev (HV)

Hardness (HV)

Std Dev (HV)

0.003 231 4 232 10 254 10 1 206 7 218 12 213 7 2 191 5 223 7 211 13 4 190 5 206 10 206 7 6 189 5 192 8 175 6

10 182 5 175 8 180 3 20 175 4 173 6 177 2

4.4 Microhardness – Ferrite Regions

During spheroidization, changes not only take place in the carbide rich regions but also in

the ferrite regions. Figure 4.14 shows the microhardness in the ferrite regions after different heat

treatment times for the three steels. Figure 4.14(a) shows the microhardness in the ferrite regions

for the HR steel. The hardness is decreasing for the first six hours. The hardness then increases

after ten hours and then again decreases at twenty hours. Figure 4.14(b) shows the microhardness

in the ferrite regions for the CR steel. The hardness decreases for the first ten hours then increases

after twenty hours. Figure 4.14(c) shows the microhardness in the ferrite regions for the Norm

steel. The hardness decreases for the first two hours then increases until ten hours. The hardness

then decreases at twenty hours. Figure 4.14(d) shows the microhardness in the ferrite regions for

all the steels. All the steels show the trend of decreasing then increasing in hardness. The reason

for this increase in hardness may be related to the growth of the grain boundary carbides between

the proeutectoid ferrite grains. The growth of these carbides can best be seen in Figure 4.5. In

Figure 4.5(a)-(c) the Norm steel has few grain boundary carbides. In Figure 4.5(d)-(f) The growth

of the grain boundary carbides start to outline the ferrite grains and may lead to the increase local

hardness of the ferrite regions.

62

The fluctuations in hardness could be due to experimental uncertainty. Table 4.3 shows

the experimental data for the ferrite microhardness. The standard deviation in ferrite hardness is

between 5-10 HV. However, the fluctuations in hardness between the maximum and minimum

are on the order of 15 HV. This suggests the hardness fluctuations measured in the ferrite are an

actual phenomenon.

4.5 Macrohardness

The changes in overall hardness were measured using a Rockwell B test. Figure 4.15

shows the hardness results for each of the three steels. Figure 4.15(a) shows the hardness of the

HR steel at various heat treatment times. The hardness decreases until six hours and again

increases at ten hours. The HR steel then decreases after twenty hours similar to the behavior of

the microhardness in the ferrite regions. Figure 4.15(b) shows the hardness of the CR steel at

various heat treatment times. The material decreases in hardness for the first four hours; then it

then increases in hardness at six hours. The CR steel then decreases again after ten hours and

stabilizes. Figure 4.15(c) shows the hardness of the Norm steel at various heat treatment times.

The hardness decreases in the first six hours, increases after ten hours, then decreases at twenty

hours. Figure 4.15(d) shows the hardness of the all the steels at the various heat treatment times.

All the microstructures undergo a similar period of decreasing hardness then a rise in hardness

followed by a decrease. The cause of the initial decrease is the spheroidization of carbides and the

softening of the ferrite. The rise in hardness may be attributed to the growth of grain boundary

carbides in the proeutectoid ferrite. The contributions of the cementite and ferrite regions may

cause the difference in peaks and valleys between the micro and marcrohardness. The hardness

could also be affected by the precipitation of chromium carbides; however, no evidence of

chromium carbides was observed.

63

120

130

Fe

140

150

160

170


rrite

Mic

roha

rdne

ss (H

V)

76

80

84

Har

dnes

s (H

RB

)

0 4 8 12 16 20


120

130

Fe

140

150

160

170

rrite

Mic

roha

rdne

ss (H

V)

76

80

84

Har

dnes

s (H

RB

)

(a) (b)

120

130

140

150

160

170

Fe

120

130

140

150

160

170

Fe


rrite

Mic

roha

rdne

ss (H

V)

76

80

84

Har

dnes

s (H

RB

)

0 4 8 12 16 20


rrite

Mic

roha

rdne

ss (H

V)

76

80

84

Har

dnes

s (H

RB

)


(c) (d)

Figure 4. 14 Microhardness in the ferrite regions for the 16MnCr5 steel after the 692 °C (1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

Table 4.3 Hardness in the Ferrite Regions for the 15MnCr5 Steel After Various Heat Treatment Times at 692 °C (1277 °F).

HR CR Norm Holding

Time (Hours)

Hardness (HV)

Std Dev (HV)

Hardness (HV)

Std Dev (HV)

Hardness (HV)

Std Dev (HV)

0.003 165 8 155 7 127 4 1 146 3 147 6 126 3 2 151 4 148 4 123 5 4 155 9 143 3 139 4 6 140 4 138 2 129 3

10 161 2 137 8 139 3 20 147 3 142 3 131 3

64


65

70

75

80

85

90

Har

dnes

s (H

RB

)


65

70

75

80

85

90

Har

dnes

s (H

RB

)

(a) (b)

65

70

75

80

85

90

Ha


65

70

75

80

85

90

Hard

ness

(HR

B)

0 4 8 12 16 20


rdne

ss (H

RB

)


(c) (d) Figure 4. 15 Macrohardness of the 16MnCr5 steel after the 692 °C (1277 °F) heat

treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

Table 4.4 shows the experimental data and standard deviations for the microhardness

testing. The standard deviation on most the hardness measurements is 1 HRB. The difference in

the maximum and minimum hardness between four and ten hours is between 4-5 HRB. This

shows the increase in hardness is not merely due to experimental uncertainty.

65

Table 4.4 Macrohardness for the 15MnCr5 Steel After Various Heat Treatment Times at 692 °C (1277 °F).

HR CR Norm Holding

Time (Hours)

Hardness (HRB)

Std Dev (HRB)

Hardness (HRB)

Std Dev (HRB)

Hardness (HRB)

Std Dev (HRB)

0.003 87 1 82 1 78 1 1 83 2 79 1 75 1 2 79 1 79 1 75 1 4 79 2 77 1 70 2 6 76 1 79 1 69 1

10 80 1 76 1 75 1 20 79 1 76 1 73 1

4.6 Compression Testing

In order to evaluate the workability of these steels, a variety of compression tests were

performed in order to initiate circumferential cracking. These tests were performed on

commercially spheroidized 15MnCr5 steel. Figure 4.16 shows the variety of samples after

compression. Figure 4.16(a)-(c) all had starting diameters of 12.7 mm (0.5 in) and heights of

19.1 mm (0.75 in). These samples were compressed at a rate of 2 mm/min (0.075 in/min) to

89 kN (20 kip), 267 kN (60 kip), and 445 kN (100 kip) respectively. The samples all used talcum

powder to increase friction between the sample and the dies. No signs of cracking were observed

in any of the samples. Figure 4.16(d) had a starting diameter of 12.7 mm (0.5 in) and height of

19.1 mm (0.75 in). The ends of this sample were roughened with 60 grit grinding paper and

talcum powder was applied to further enhance the friction conditions. The sample was

compressed to 445 kN (100 kip) at a rate of 2 mm/min (0.075 in/min). There were no signs of

cracking on the circumference of sample (d). Figure 4.16(e) had a starting diameter of 12.7 mm

(0.5 in) and height of 19.1 mm (0.75 in). Talcum powder was used to increase friction between

the sample and the dies. Sample (e) was compressed to 445 kN (100 kip) at a rate of

1270 mm/min (50 in/min). No cracking was observed around the circumference of sample (e).

66

Figure 4.16(f) had a starting diameter of 7.7 mm (0.3 in) and heights of 11.4 mm (0.45 in).

Talcum powder was used to increase friction between the sample and the dies. This sample was

compressed to 445 kN (100 kip) at crosshead speed of 1 mm/min (0.045 in/min). No cracking

was observed around the circumference of sample (f). Figure 4.16(g) had a starting diameter of

5.1 mm (0.2 in) and heights of 7.7 mm (0.3 in). Talcum powder was used to increase friction

between the sample and the dies. This sample was compressed to 240 kN (54 kip) at crosshead

speed of 1 mm/min (0.045 in/min). No cracking was observed around the circumference of

sample (g). Figure 4.16(h) had a starting diameter of 12.7 mm (0.5 in) and height of 19.1 mm

(0.75 in). The ends of the sample were constrained with a special die to obtain the maximum

amount of friction. The sample was compressed to 445 kN (100 kip) at a rate of 2 mm/min

(0.075 in/min). Sample (h) was subjected to the highest friction and load the frame would allow.

No cracking was observed around the circumference of sample (h). The total circumferential

strain imparted to sample (h) was 0.864 in/in. Since sample (h) did not fracture at a

circumferential strain of 0.864, the fracture strain for this material must occur at a higher strain.

Since compression testing did not initiate a fracture, other testing methods were used to evaluate

the cold workability. A modified bend test was performed and is discussed in Appendix E. Room

temperature Charpy impact testing was also performed and is discussed in Appendix F.

Figure 4.16 Compression samples utilizing different stresses, strain rates, sample geometries and frictional conditions.(a) Compressed to 20 kip, (b) compressed to 60 kip, (c) compressed to limit (0.075 in/min), (d) compressed to limit with roughened ends, (0.075 in/min), (e) compressed to limit (50 in/min), (f) 0.3 in diameter compressed to limit (0.045 in/min), (g) 0.2 in diameter compressed to 54 kip (0.045 in/min), (h) compressed to limit with constrained ends and no talcum powder (0.075 in/min).

67

4.7 Tension Testing

Since no fractures occurred during the compression testing, tension testing was used to

determine the cold formability of the material. Tension tests were carried out at a crosshead

velocity of 495 mm/min (19.5 in/min) to better simulate forging speeds; however, real forging

speeds are much faster. Adiabatic heating calculations were performed on the three steels using

pCdT

ρεησ

=Δ (4.2)

Where ΔT is the change in temperature, η is the fraction of energy stored in lattice defects

approximated to be 0.95, ρ is the density, and Cp is the heat capacity. [24] Table 4.5 shows the

increase in temperature in the tensile specimen for both the as-received and twenty hour heat

treated samples. The HR and CR steels initially have an approximately 40 ºC (72 °F) raise in

temperature, which increases to approximately 45 °C (81 °F) at twenty hours. The Norm steel

maintains an approximate 45 °C (81 °F) increase in temperature from tensile testing.

Table 4.5 Increase in Temperature from Adiabatic Heating during the Tension Test on the 16MnCr5 Steel Tested at a Crosshead Velocity of 495 mm/min (19.5 in/min).

As-Received 20 Hours

Delta T (°C)

Std Dev (°C)

Delta T (°C)

Std Dev (°C)

CR 41.3 1.0 44.7 2.7 HR 39.8 1.6 43.6 1.2

Norm 45.2 1. 4 44.6 0.7

Tension tests can be used to evaluate workability but care must be taken when applying

the results to forming situations. Reduction in area and percent elongation are common ways

measure of ductility and the ultimate tensile strength can provide estimates of flow strength in

some situations. Figure 4.17 shows engineering stress-strain curves for all three steels in the as-

received state, after six hours of heat treatment, and after twenty hours of heat treatment. Figure

4.17(a) shows engineering stress-strain curves for the HR steel in three heat treatment conditions.

68

In the as-received state the HR steel has an ultimate tensile strength of 724 ±5 MPa (105 ksi) and

a total elongation of 23 ±0.5%. After twenty hours of heat treatment, the ultimate tensile stress

has decreased to 545 ±3 MPa (79 ksi) and the total elongation has increased to 35 ±1%. Figure

4.17(b) shows engineering stress-strain curves for the CR steel in three heat treatment conditions.

In the as-received state the CR steel has an ultimate tensile strength of 675 ±8 MPa (98 ksi) and a

total elongation of 25.3 ±1.7%. After twenty hours of heat treatment, the ultimate tensile stress

has decreased to 537 ±3 MPa (78 ksi) and the total elongation has increased to 36.0 ±2.4%.

Figure 4.17(c) shows engineering stress-strain curves for the Norm steel in three heat treatment

conditions. In the as-received state the Norm steel has an ultimate tensile strength of 593 ±4 MPa

(86 ksi) and a total elongation of 33.0 ±0.8%. After twenty hours of heat treatment, the ultimate

tensile stress has decreased to 524 ±3 MPa (76 ksi) and the total elongation has increased to

35.5 ±1%. The Cockcroft and Latham constant can also be used to measure the workability in

steels. Appendix G discusses the Cockcroft and Latham constants for the three steels.

4.7.1 Reduction in Area

Bailey et al. propose the reduction in area during a tensile test to be the best measure of

workability. Reduction in area measures necking strain and indicates the ability of the material to

resist crack propagation. [18] Figure 4.18 shows the average reduction in area data for each of the

three steels. Figure 4.18(a) shows the average reduction in area for the HR steel. The HR steel has

a reduction in area approximately 65 ±0.5% at ten seconds of heat treatment, but after twenty

hours of spheroidization heat treatment the HR steel has reached over 73 ±0.5% reduction in area.

Figure 4.18(b) shows the average reduction in area for the CR steel. The CR steel has over 68

±0.5% reduction at ten seconds of heat treatment and increases to over 73 ±0.5% after twenty

hours of spheroidization. Figure 4.18(c) shows the average reduction in area for the Norm steel.

After ten seconds, the reduction in area for the Norm steel is over 68 ±0.5%. However, after

twenty hours, the reduction in area is approximately 72 ±0.5%. Figure 4.18(d) shows the

69

reduction in area for all three steels. The Norm steel has the largest initial reduction in area and

the lowest after twenty hours. The HR has the lowest initial reduction in area and the CR and HR

have similar reductions in area after twenty hours.

0 0.1Enginee

0.2 0.3 0.4ring Strain (in/in)

0

200

400

600

Engi

neer

ing

Stre

ss (M

Pa)

0

40

80

Engi

neer

ing

Stre

ss (k

si)

As-Recieved

20 hrs6 hrs

0 0.1 0.2 0.3 0.4Enginee

0

200

400

600

Engi

nee

ring Strain (in/in)

ring

Stre

ss (M

Pa)

0

40

80

En)

gine

erin

g St

ress

(ksi

As-Recieved

20 hrs6 hrs

(a) (b)

0 0.1 0.2 0.3 0.4Enginee

0

200

400

600

Engi

nee

0

40

80

Engi

nee

ess (

MPa

)

ksi)

ress

(

ring Strain (in/in)

ring

Str

ring

St

As-Recieved

20 hrs6 hrs

(c)

Figure 4.17 Typical engineering stress-strain curves for the 15MnCr5 steels at various heat treatment conditions. (a) HR steel as-received, 6 hours, and 20 hours. (b) CR steel as-received, 6 hours, and 20 hours. (c) Norm steel as-received, 6 hours, and 20 hours.

70


64

66

68

70

72

74

Red

uctio

n in

Are

a (%

)

0 4 8 12 16 2Holding Ti

0me (Hours)

64

66

68

70

72

74

Red

uctio

n in

Are

a (%

)

(a) (b)

64

66

68

70

72

74

Re

0 4 8 12 16 20Holding Ti

64

66

68

70

72

74

Re

me (Hours)

duct

ion

in A

rea

(%)


duct

ion

in A

rea

(%)


(c) (d)

Figure 4.18 Average reduction in area after tensile testing for the 16MnCr5 steel after the 692 °C (1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

4.7.2 Uniform Elongation

Figure 4.19 shows the uniform elongation for three steels at various times during heat

treatment. Figure 4.19(a) shows the uniform elongation for the HR steel. The uniform elongation

increases approximately 0.030 in/in in the first four hours and then reaches a saturation point at

71

0.147 ±0.002 in/in. Figure 4.19(b) shows the uniform elongation for the CR steel. The uniform

elongation increases approximately 0.015 in/in in the first four hours and then reaches a

saturation point at 0.157 ±0.002 in/in. Figure 4.19(c) shows the uniform elongation for the Norm

steel. The uniform elongation increases approximately 0.015 in/in in the first four hours and then

reaches a saturation point at approximately 0.173 ±0.004 in/in. Figure 4.19 (d) shows the

uniform elongation for all three steels. The HR consistently has the lowest uniform elongation

and the Norm steel consistently has the highest uniform elongation. The prior microstructure

seems to be controlling the behavior of the uniform elongation. The larger amounts of

proeutectoid ferrite may give the CR and Norm steel higher uniform elongations than the HR

steel. The Norm steel also appears to have a finer grain size than the CR steel giving it more

uniform elongation than the CR steel. It is not fully understood why the uniform elongation of all

the steels saturate after four hours.

4.7.3 Total Elongation

The total elongation is a commonly measured quantity to determine ductility in tension

tests. Figure 4.20 shows the measured total elongation data for the three steels. Figure 4.20(a)

shows the total elongation for the HR steel. The total elongation increases linearly for the first six

hours to 34.3 ±2.3%. The total elongation decreases by 0.8% at ten hours then increases by 1.3%

after twenty hours. Figure 4.20(b) shows the total elongation for the CR steel. The total

elongation increases linearly for the first four hours to 34.8 ±0.5%. The total elongation decreases

by 1.3% at ten hours then increases by 2.5% after twenty hours. Figure 4.20(c) shows the total

elongation for the Norm steel. The total elongation increases linearly for the first four hours to

35.5 ±1.3%. The total elongation decreases by 1.2% at ten hours then increases by 1.2% after

twenty hours. Figure 4.20(d) shows the total elongation for all three steels. The Norm steel has

the highest percent total elongation for the first ten hours but the CR steel has a slightly higher

value after twenty hours. The HR steel has the lowest total elongation throughout.

72


0.1

0.12

0.14

0.16

0.18

Uni

form

Elo

ngat

ion

(in/in

)


0.1

0.12

0.14

0.16

0.18

Uni

form

Elo

ngat

ion

(in/in

)

(a) (b)

0.1

0.12

0.14

0.16

0.18

Uni

fo

0.1

0.12

0.14

0.16

0.18

Uni

fo


rm E

long

atio

n (in

/in)


rm E

long

atio

n (in

/in)


(c) (d)

Figure 4.19 Average uniform elongation during a tensile test for the 16MnCr5 steel after the 692 °C (1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

The total elongation is a combination of uniform elongation and non-uniform elongation.

Table 4.6 shows the uniform and non-uniform elongation values for all three steels at various

times. The uniform elongation data for the CR and Norm steel gradually increase with time. The

non-uniform elongation data decreases and then increases for the three steels. The behavior of the

non-uniform elongation indicates the maximum and minimum values seen in the total elongation

can be attributed to the non-uniform elongation.

73


26

28

30

32

34

36Pe

rcen

t Elo

ngat

ion

(%)

0 4 8 12 16 2Holding Ti

0me (Hours)

26

28

30

32

34

36

Perc

ent E

long

atio

n (%

)(a) (b)

26

28

30

32

34

36

Pe

0 4 8 12 16 20Holding Ti

26

28

30

32

34

36Pe

me (Hours)

rcen

t Elo

ngat

ion

(%)


rcen

t Elo

ngat

ion

(%)


(c) (d) Figure 4.20 Average total elongation for the 16MnCr5 steel after the 692 °C (1277 °F) heat

treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

74

Table 4.6 Uniform and Non-Uniform Elongation Values for the 15MnCr5 Steel at Various Heat Treatment Times at 692 °C (1277 °F).

HR Uniform

Elongation (%)

CR Uniform

Elongation (%)

Norm Uniform

Elongation (%)

CR Non-Uniform

Elongation (%)

HR Non-Uniform

Elongation (%)

Norm Non-

Uniform Elongation

(%) 11.7 14.4 16.0 17.1 15.8 17.0 12.5 15.2 16.5 16.6 16.0 16.0 13.0 15.2 16.3 18.3 17.5 18.2 14.6 15.9 17.5 18.9 18.9 18.0 14.8 15.6 17.6 18.7 19.5 17.1 14.9 15.9 17.6 17.6 18.6 16.7 14.7 15.7 17.3 20.3 20.1 18.2

The total elongation data differs from the reduction in area data significantly. The

reduction in area data shows the CR steel to be the most formable and the total elongation data

show the Norm steel to be the most formable. The reduction in area data are preferred for the

purposes of formability and correlate better to the spheroidization data. For these reasons, the

elongation data may not be well suited to measure cold formability.

4.7.4 Ultimate Tensile Strength

The ultimate tensile strength (UTS) has been used to approximate the flow strength for

hot forging operations. However, the correlation between the flow strength and the UTS becomes

less reliable as the temperature decreases. [18] For cold forging operations the UTS is not an

appropriate approximation of flow strength but can demonstrate the strength differences between

the three steels. Figure 4.21 shows the average UTS for the three different steels at the various

heat treatment times. Figure 4.21 (a) shows the average UTS for the HR steel at various heat

treatment times. The UTS after ten seconds of heat treatment is 634 ±4 MPa (92 ksi) and declines

logarithmically to 545 ±3 MPa (79 ksi). Figure 4.21(b) shows the average UTS for the CR steel at

various heat treatment times. The UTS after ten seconds of heat treatment is 620 ±3 MPa (90 ksi)

and declines logarithmically to 537 ±3 MPa (78 ksi). Figure 4.21(c) shows the average UTS for

75

the Norm steel at various heat treatment times. The UTS after ten seconds of heat treatment is

586 ±2 MPa (85 ksi) and declines logarithmically to 524 ±3 MPa (76 ksi). Figure 4.21 (d) shows

the UTS for all three microstructures. The HR and CR have almost identical trends in decreasing

UTS. The Norm steel has lower UTS for all times but the UTS decreases at a similar rate

compared to the HR and CR steels.


520

560

600

640

Ulti

mat

e Te

nsile

Stre

ngth

(MPa

)

76

80

84

88

92

Ulti

mat

e Te

nsile

Stre

ngth

(ksi

)


520

560

600

640

Ulti

mat

e Te

nsile

Stre

ngth

(MPa

)

76

80

84

88

92

Ulti

mat

e Te

nsile

Stre

ngth

(ksi

)

(a) (b)

520

560

600

640

Ul


timat

e Te

nsile

Stre

ngth

(MPa

)

76

80

84

88

92

Ulti

mat

e Te

nsile

Stre

ngth

(ksi

)

520

560

600

640

Ul

0 4 8 12 16 20Holding Ti

76

80

84

88

92

Ul

MicrostructuresHR

me (Hours)

timat

e Te

nsile

Stre

ngth

(MPa

)

timat

e Te

nsile

Stre

ngth

(ksi

)

CRNorm

(c) (d) Figure 4.21 Average ultimate tensile strength for the 16MnCr5 steel after the 692 °C

(1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

76

7.7.5 Yield Strength

The yield strength is another commonly measured property in a tensile test. Figure 4.22

shows the average upper yield strength for all three steels at various temperatures. Upper yield

strength is an easily identifiable value measured from stress-strain data. Figure 4.22(a) shows the

upper yield strength for the HR steel at various times. The yield strength is 483 ±5 MPa (70 ksi)

after ten seconds of heat treatment decreases logarithmically to 407 ±3 MPa (59 ksi) after six

hours and stays constant. Figure 4.22(b) shows the upper yield strength for the CR steel at various

times. The yield strength is 441 ±7 MPa (64 ksi) after ten seconds of heat treatment decreases

logarithmically to 413 ±8 MPa (60 ksi) after four hours and stays constant. Figure 4.22(c) shows

the upper yield strength for the Norm steel at various times. The yield strength is 441 ±3 MPa

(70 ksi) after ten seconds of heat treatment decreases logarithmically to 413 ±1 MPa (60 ksi) after

four hours and stays constant. Figure 4.22(d) shows the yield strength for all three steels for

various heat treatment times. The HR steel has the highest initial yield strength and decreases to

the lowest after six hours. The CR and Norm steels have similar yield strength behavior with the

CR steel having slightly lower yield strength throughout.

77


400

420

Upp

440

460

480

500er

Yie

ld S

treng

th (M

Pa)

60

U

64

68

72

pper

Yie

ld S

treng

th (k

si)


400

420

Upp

440

460

480

500

er Y

ield

Stre

ngth

(MPa

)

60

U

64

68

72

pper

Yie

ld S

treng

th (k

si)

(a) (b)

400

420

440

460

480

500

Upp

er Y

iel


d St

reng

th (M

Pa)

60

64

68

72U

pper

Yie

ld S

treng

th (k

si)

400

420

440

460

480

500

Upp

er Y

iel

0 4 8 12 16 20Holding Ti

60

64

68

72

U

MicrostructuresHRCR

me (Hours)

d St

reng

th (M

Pa) )

Norm ksi

h (

pper

Yie

ld S

treng

t

(c) (d) Figure 4.22 Average upper yield strength after a tensile test for the 16MnCr5 steel after the

692 °C (1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

78

CHAPTER 5

DISCUSSION

This chapter discusses the relationships between the mechanical property data measured

in the tension test to the microstructure developed during the spheroidization process. This

chapter also provides an explanation for the variation in the total elongation as a function of

spheroidization time.

5.1 Reduction in Area

The reduction in area, which is one measure of workability, is affected by the

microstructure of the steel. The variations seen in the reduction in area data can be explained by

the microstructural evolution occurring during the spheroidization heat treatment. The cause of

the high initial reduction in area for the non spheroidized Norm steel is the coarse pearlite

microstructure, which has a more ductile structure than the fine pearlite of the CR and bainite of

the HR. As spheroidization time increases the Norm steel has finer carbides and is not as highly

spheroidized as the HR and CR steels so it has a lower reduction in area. Both the HR and CR

steels have higher percentages of spheroidization for the longer heat treatment times giving them

a greater reduction in area after these longer heat treatment times. However, as the heat

treatments approach twenty hours, all three materials converge to a similar value of reduction in

area. Figure 5.1 shows the relationship between the reduction in area and the percentage of

spheroidization. Figure 5.1(a) shows the relationship between reduction in area and percentage of

spheroidization for the HR steel. The HR steel has a high percentage of spheroidization at ten

seconds (60%) and a low reduction in area (65%). The HR steel increases 13% in reduction in

area with only a 33% change in spheroidization percentage at twenty hours. Figure 5.1(b) shows

the relationship between reduction in area and percentage of spheroidization for the CR steel. The

79

CR steel has 43% spheroidization at ten seconds and a 68% reduction in area. The CR steel

increases 5% in reduction in area with a 51% change in spheroidization percentage at twenty

hours. Figure 5.1(c) shows the relationship between reduction in area and percentage of

spheroidization for the Norm steel. The Norm steel has 46% spheroidization at ten seconds and a

70% reduction in area. The Norm steel increases 2% in reduction in area with a 38% change in

spheroidization percentage at twenty hours. Figure 5.1(d) shows the relationships between

reduction in area and percentage of spheroidization for all the steels. The reduction in area is

dominated by the prior microstructure at low spheroidization percentages; however, each prior

microstructure seems to converge to around 73% reduction in area between 95-100%

spheroidization. Since all three steels have the same composition, at 100% spheroidization the

microstructures for all the steels become ferrite with spherical carbides dispersed throughout.

Therefore, little difference should be seen in the reduction in area at high spheroidization

percentages when the microstructures become similar.

The reduction in area results are the closest measurement of workability provided by the

tension tests. [18] Figure 5.1 shows the Norm steel has the best reduction in area at lower

percentages of spheroidization. Therefore, the Norm steel will have better workability at lower

percentages of spheroidization. However, since the reductions in area converge at high

percentages of spheroidization, the HR or CR steel will have similar workability results to the

Norm at high percentages of spheroidization. In addition, the CR and HR will spheroidize in a

shorter timeframe giving them better reductions in area than the Norm steel after only six hours of

heat treatment.

5.2 Total Elongation

Total elongation is comprised of uniform and non-uniform elongations. Table 4.4 showed

the uniform and non-uniform elongation values for the three steels. The uniform elongation

slightly increased as spheroidization percentage increased. The non-uniform elongation, however,

80

lowered then increased for the three steels. The difference in the behavior between the uniform

and non-uniform elongation suggests the non-uniform elongation is the cause of the maxima and

minima in the total elongation curves.

20 40 60 80 100Area Percent Spheroidized (%)

66

68

70

72

74

Red

uctio

n in

Are

a (%

)


66

68

70

72

74

ruc

tion

in A

Red

ea (%

)

(a) (b)

20 40 60 80 100A

64

66

68

70

72

7

rea Percent Spheroidized (%)

4

Red

uctio

n in

Are

a (%

)

20 40 60 80 100A

64

66

68

70

72

74

rea Percent Spheroidized (%)

Red

uctio

n in

Are

a (%

)


(c) (d) Figure 5.1 Reduction in area and the corresponding percentage of spheroidization for the

16MnCr5 steel for various prior microstructures. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, (d) all microstructures.

81

The total elongation values seen in Chapter 4 exhibited maxima and minima between

four to ten hours of heat treatment. These maximums and minimums on these curves appear

similar to the maximums and minimums on the ferrite microhardness curves. The changes in

ferrite microhardness could be due to the precipitation of chromium carbides. The changes in

ferrite microhardness curves may also be attributed to the presence of grain boundary carbides

between proeutectoid ferrite grains. These carbides might also be cause for the variation in total

elongation that is observed. The grain boundary carbides can act as nucleation sites for

microvoids that would decrease the non-uniform elongation behavior. The maximum on the HR

total elongation occurs at six hours and the minimum occurs at ten hours. These values

correspond well to the values of the minimum and maximum points for the HR ferrite

microhardness. Similarly, the Norm steel undergoes a maximum at four hours and a minimum at

ten hours. The ferrite microhardness goes through a minimum between two and four hours and a

maximum at ten hours. The CR ferrite hardness does not follow the same trend as the HR and

Norm, but because uncertainty on the hardness the CR ferrite is high, the CR ferrite could follow

a similar trend.

Figure 5.2 shows the relationship between the total elongation and the ferrite

microhardness for all three steels. Figure 5.2(a) compares the total elongation and the ferrite

microhardness for the HR steel. The values for the ferrite hardness and the total elongation all

seem to lie in the same region outlined by the dashed ellipse. Figure 5.2(b) compares the total

elongation and the ferrite microhardness for the CR steel. The values for the ferrite hardness and

the total elongation all seem to lie in the same region outlined by the solid line ellipse.

Figure 5.2(c) compares the total elongation and the ferrite microhardness for the Norm steel. The

values for the ferrite hardness and the total elongation all seem to lie in the same region outlined

by the long dashed ellipse. Figure 5.2(d) shows the hardness and total elongation grouping for all

the steels. These ellipses seem to form a linear relationship that can be used to relate the ferrite

hardness with the total elongation. The HR steel does not fit the trend as well as the CR and

82

Norm, but this discrepancy is probably due to the small percentage (16%) of proeutectoid ferrite

in the HR steel.

120 130 140 150 160 170Ferrite Hardness (HV)

24

28

32

Tota

l Elo

ngat

ion

36

(%)

76 80 84Ferrite Hardness (HRB)


24

28

32

Tota

l Elo

ngat

ion

36

(%)


(a) (b)




24

28

32

36

Tota

l Elo

ngat

ion

(%)


24

28

32

36

Tota

l Elo

ngat

ion

(%)


(c) (d) Figure 5.2 Total elongation and the corresponding ferrite microhardness for the 16MnCr5

steel for various prior microstructures. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, (d) all microstructures.

The linear trend shown in Figure 5.2(d) shows a direct relationship between the hardness

of the ferrite and the total elongation. Therefore in order to obtain higher amounts of total

elongation the ferrite should be as soft as possible. The Norm steel had the highest total

83

elongation because it had the softest ferrite due to the normalizing heat treatment after hot-rolling.

However, the total elongation was reduced by the growth of carbides in the proeutectoid ferrite.

Shorter heat treatment times have a reduced number of these carbides in the proeutectoid ferrite

with an increase in total elongation observed during the early stages of the spheroidization heat

treatment.

The total elongation is a common measure of ductility in a tensile test. However, the total

elongation is not commonly used to measure workability. Reduction in area is a better

measurement of workability. Total elongation accounts for deformation under both uniform

elongation and non-uniform elongation. Cold-forming usually consists of complex stress states

more closely represented by the triaxial stress during non-uniform deformation than by the

uniaxial stress state occurring during uniform deformation. Reduction in area measurements

better capture the amount of deformation without fracture that the metal can undergo while a

triaxial stress state is imposed.

5.3 Ultimate Tensile Strength

The UTS decreases as heat treatment time and spheroidization progresses. Figure 5.3

shows the relationship between the UTS and the percent spheroidization. Figure 5.3(a) shows the

relationship between UTS and percent spheroidization for the HR steel. The UTS is 627 MPa

(91 ksi) at 60% spheroidization and decreases to 545 MPa (79 ksi) at 93% spheroidization. Figure

5.3(b) shows the relationship between UTS and percent spheroidization for the CR steel. The

UTS is 621 MPa (90 ksi) at 43% spheroidization and decreases to 537 MPa (78 ksi) at 94%

spheroidization. Figure 5.3(c) shows the relationship between UTS and percent spheroidization

for the Norm steel. The UTS is 558 MPa (81 ksi) at 46% spheroidization and decreases to

524 MPa (76 ksi) at 84% spheroidization. Figure 5.3(d) shows the relationship between UTS and

percent spheroidization for all three steels. The UTS decreases at different rates for each prior

microstructure but converge to approximately 525 MPa (76 ksi) between 95-100%

84

spheroidization. The convergence of these lines shows the decrease in UTS is controlled by prior

microstructure at low percentages of spheroidization, but at high percentages of spheroidization

(95-100%), the microstructures are similar so the UTS converges on a single value.


560

600

Ul

640

timat

e Te

nsile

Stre

ngth

(MPa

)

76

80

84

88

Ul

92

timat

e Te

nsile

Stre

ngth

(ksi

)20 40 60 80 100

Area Percent Spheroidized (%)

640

560

600

Ulti

mat

e Te

nsile

Stre

ngth

(MPa

)

76

80

84

88

Ul

92

ktim

ate

Tens

ile S

treng

th (

si)

(a) (b)


560

600

640

Ulti

mat

e Te

nsile

Stre

ngth

(MPa

)

76

80

84

88

92

Ulti

mat

e Te

nsile

Stre

ngth

(ksi

)

20 40 60 80 100Area Percen

560

600

640

Ul

76

80

84

88

92

Ul

t Spheroidized (%)

timat

e Te

nsile

Stre

ngth

(MPa

)

timat

e Te

nsile

Stre

ngth

(ksi

)


(c) (d)

Figure 5.3 Ultimate tensile strength and the corresponding percent spheroidization for the 16MnCr5 steel for various prior microstructures. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, (d) all microstructures.

As the total elongation increases, the UTS decreases. Similarly, the reduction in area

increases as the UTS decreases. Figure 5.4(a) shows the relationship between UTS and total

elongation. The relationship for all prior microstructures is similar. The general trend is an

increase in total elongation with a decrease in UTS as indicated by the trend line shown. This

85

trend agrees well with data shown by Syn et al. for spheroidized 1.8% C steel. [25] Figure 5.4(b)

shows the relationship between the UTS and the reduction in area. The general trend is an

increase in reduction in area with a decrease in UTS as indicated by the trend line shown. The

scatter in the reduction in area data is less than that of the total elongation data. This may be due

to the effect of the grain boundary carbides on the total elongation.

26 28 30 32 34 36Total Elongation (%)

560

600

Ulti

mat

e Te

nsile

Stre

ngth

(MPa

)

76

80

84

88

92

Ulti

mat

e Te

nsile

Stre

ngth

(ksi

)


64 66 68 70 72 74Reduc

76

80

84

88

92

Ul

tion in Area (%)

560

600

Ulti

mat

e Te

nsile

Stre

ngth

(MPa

)

timat

e Te

nsile

Stre

ngth

(ksi

)


(a) (b) Figure 5.4 (a) The relationship between ultimate tensile strength and total elongation for the

15MnCr5 steel. (b) The relationship between ultimate tensile strength and reduction in area for the 15MnCr5 steel.

Yield Strength

The pearlitic structures of the CR and Norm steels have lower initial yield strengths

compared to the bainitic structure of the HR steel. However as spheroidization takes place the

yield strength of all three steels drop. Figure 5.5 shows the relationship between the yield strength

and the percent spheroidization for the three steels. Figure 5.5(a) shows the relationship between

yield strength and percent spheroidization for the HR steel. The yield strength drops from

483 MPa (70 ksi) at 60% spheroidization to 407 MPa (59 ksi) at 93% spheroidization. Figure

5.5(b) shows the relationship between yield strength and percent spheroidization for the CR steel.

The yield strength drops from 441 MPa (64 ksi) at 43% spheroidization to 407 MPa (59 ksi) at

86

94% spheroidization. Figure 5.5(c) shows the relationship between yield strength and percent

spheroidization for the Norm steel. The yield strength drops from 427 MPa (62 ksi) at 46%

spheroidization to 414 MPa (60 ksi) at 84% spheroidization. Figure 5.5(d) shows the relationship

between the yield strength and the percent spheroidization for all three steels. The three distinct

lines created by the different initial microstructures show the yield strength is not merely

controlled by the percentage of spheroidization but it is also affected by the prior microstructure.

The yield strengths do, however, decrease to toward a value of approximately 410 MPa (59 ksi) at

percentages of spheroidization between 95-100%. As the microstructures become fully

spheroidized, the differences between the microstructures diminish. The yield strength

approaching a common value at high percentages of spheroidization is expected.

Figure 5.6(a) shows the relationship between yield strength and total elongation. As the

total elongation increases, the upper yield strength decreases shown by the trend line. This trend

agrees well with data shown by Syn et al. for spheroidized 1.8% C steel. [25] Figure 5.6(b) shows

the relationship between the yield strength and the reduction in area. Similarly, the reduction in

area increases as the upper yield strength decreases as shown by the trend line. The scatter in the

reduction in area data is much less than that of the total elongation data. This may be due to the

effect of the grain boundary carbides on the total elongation. Since upper yield strength is an

easily measured quantity, the relationship between the yield strength and the reduction in area

could be used to estimate the reduction in area. Knowing the reduction in area, an estimation of

workability can be obtained easily.

5.5 Effects of Initial Microstructure

The values of reduction in area, UTS, and yield strength seem to change depending on

the starting microstructure of the steel. However, at large values of spheroidization (95-100%),

the values for reduction in area, UTS, and yield strength do seem to converge to a common value.

This convergence suggests the properties of partially-spheroidized steels are dependent on the

87

microstructure prior to heat treatment. At large values of spheroidization, the microstructures

become similar, ferrite with small spherical carbides. This similarity in microstructure leads to a

commonality in the values of reduction in area, UTS, and yield strength. The total elongation,

however, may be affected by the growth of grain boundary carbides and the strength of the ferrite,

which still vary with prior microstructure.


400

440

480

Upp

er Y

ield

Stre

ngth

(MPa

)

56

60

64

68

72

Upp

er Y

ield

Stre

ngth

(ksi

)


400

440

480

Upp

er Y

ield

Stre

ngth

(MPa

)

56

60

64

68

72

Upp

er Y

ield

Stre

ngth

(ksi

)

(a) (b)

20 40 60 80 100Area Percen

56

60

64

68

72

Upp

e

t Spheroidized (%)

400

440

480

Upp

er Y

ield

Stre

ngth

(MPa

)

r Yie

ld S

treng

th (k

si)

20 40 60 80 100Area Percen

56

60

64

68

72

Upp

e

MicrostructuresHR

400

440

480

Upp

e

CR

t Spheroidized (%)

r Yie

ld S

treng

th (M

Pa)

r Yie

ld S

treng

th (k

si)

Norm

(c) (d) Figure 5.5 Upper yield strength and the corresponding percent spheroidization for the

16MnCr5 steel for various prior microstructures. (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, (d) all microstructures.

88

26 28 30 32 34 36Total Elongation (%)

400

Upp

440

480

er Y

ield

Stre

ngth

(MPa

)

56

60 U

64

68

72

pper

Yie

ld S

treng

th (k

si)


64 66 68 70 72 74Reduc

400

Upp

440

480

er Y

ield

Stre

ngth

(MPa

)

56

60 U

64

68

72

pper

Yie

ld S

treng

th (k

si)


tion in Area (%)

(a) (b) Figure 5.6 (a) The relationship between yield strength and total elongation for the

15MnCr5 steel. (b) The relationship between yield strength and reduction in area for the 15MnCr5 steel.

5.6 Industrial Relevance

Spheroidization heat treatments require vast amounts of time and energy. In order to

decrease the time associated with spheroidization the starting microstructure must be chosen

carefully. The Norm steel had higher values of reductions in area at low percentages of

spheroidization; however, the heat treatment to normalize steel only adds cost and time. The CR

steel had the greatest reduction in area and therefore the greatest cold workability after twenty

hours. The CR steel also had the second lowest values for UTS and yield strength. The twenty

hour heat treatment yielded highest workability; however depending on the forming operation, a

six to ten hour heat treatment may be all that is necessary. After six hours of heat treatment, the

values of reduction in area, UTS, and yield strength are not changing drastically. A six hour heat

treatment will also limit the growth of the grain boundary carbides and maintain high total

elongation.

89

90

CHAPTER 6

SUMMARY

15MnCr5 steel was with two different pearlite structures and a bainitic structure was

subcritically spheroidized at 692° C (1277° F) and evaluated with scanning electron microscopy

and image analysis. The resulting mechanical properties from these heat treatments were

evaluated with hardness testing and tension testing.

1. The HR steel spheroidized the most quickly reaching 90% spheroidization in just two

hours. The CR steel was the next to reach 90% spheroidization in approximately ten

hours. The cementite defects in these hot rolled steels likely accelerated the

spheroidization kinetics. The Norm steel was the slowest to spheroidize and had only

reached 84% spheroidization after twenty hours. The HR steel also had the largest

average spheroidized particle size after twenty hours, 0.033 µm2. The CR had the second

largest particle size of 0.027 µm2 after twenty hours. The Norm steel had the smallest

average particle size of 0.020 µm2 after twenty hours.

2. Even though two steels may have the same percentage of spheroidization, the properties

measured from the tensile test are dependent on the prior microstructure. The reduction in

area, UTS, and yield strength changed with heat treatment time, however these changes

were dependent on the initial microstructure. The values of reduction in area, UTS, and

yield strength approach common values at near 100% spheroidization because at high

values of spheroidization the microstructures become similar. The approximate values of

these properties at 100% spheroidization are: reduction in area 73%, UTS 525 MPa

(76 ksi), and yield strength 410 MPa (59 ksi).

3. The workability of the steel was estimated by the reduction in area during the tension test

due to the triaxial stress state in the necked region. The Norm steel had better workability

91

at lower percentages of spheroidization but after approximately six hours of heat

treatment the higher percentages of spheroidization in the CR and HR steels result in

higher workability. The reduction in area and therefore the workability approach a

common value at near 100% spheroidization, due to the similarity in the microstructures.

4. Since the upper yield strength is an easily measured quantity from the tensile test, the

correlation between the yield strength and the reduction in area may provide an easy way

to estimate the workability of the 15MnCr5 steel.

5. The CR steel has the highest reduction in area and therefore the highest workability at

times greater than six hours of heat treatment. The CR steel does not have the additional

heat treatment of the Norm steel and the CR steel has similar values for reduction in area,

UTS, and yield strength at times greater than six hours. Spheroidizing from the CR state

could lead to potential time and energy savings for spheroidization treatments.

6. The total elongation values increased, decreased, and then increased over the heat

treatment time. These variations in total elongation are related to the behavior during

non-uniform elongation. The decrease in total elongation at high heat treatment times

may be due to the growth of ferrite grain boundary carbides. The carbides may serve as a

nucleation site for microvoids and decrease the total elongation. The relationship between

the ferrite hardness and the total elongation shows that in order to obtain higher levels of

total elongation the hardness of the ferrite must be reduced. The Norm steel had the

softest ferrite and therefore had the highest total elongations. The HR steel had the

hardest ferrite and had the lowest total elongations.

92

CHAPTER 7

FUTURE WORK

1. Performing upset testing on the spheroidized steels would better characterize the

workability of these steels. A forging hammer or a high speed mechanical press with a

high strain rate would be required due to the high ductility of the material. The hammer

or press could provide a high strain rate that may cause the material to crack.

2. Since upset testing could not be performed, the reductions in area were used to estimate

the workability of these steels. A comparison of the workability generated from the upset

tests should be compared to the reduction in area data to see how well the reduction in

area data represent the cold workability of this steel.

3. The current study examined only bainitic and pearlitic microstructures. Under the proper

processing conditions, a martensitic microstructure could be formed and then

spheroidized. Evaluation of the martensitic microstructures would then be appropriate.

4. The total elongation values seemed to be affected by the growth of grain boundary

carbides. The grain boundary carbides should be examined more carefully to better

understand how and to what extent this phenomenon is occurring. Fractography could be

performed to see if these grain boundary carbides are nucleating microvoids and

decreasing total elongation.

93

94

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[23] ASTM, "E8 Standard Test Methods of Tension Testing of Metallic Materials, "Annual Book or ASTM Standards, West Conshohocken, PA: American Society for Testing and Materials, Vol. 3.01.

[24] A.S. Korhonen and H.J. Kleemola, “Effects of Strain Rate and Deformation Heating in

Tensile Testing,” Metallurgical and Materials Transactions A, vol. 9, no. 7, July 1978, pp. 979-986.

[25] C.K. Syn, D.R. Lesuer, and O.D. Sherby, “Influence of Microstructure on Tensile

Properties of Spheroidized Ultrahigh-Carbon (1.8 Pct C) Steel,” Metallurgical and Materials Transactions A, vol. 25A, 1994, pp. 1481-1493.

96

APPENDIX A

INTERCRITICAL ANNEALING

This appendix covers an intercritical annealing cycle that was applied to all three prior

microstructures and compared to the subcritical treatment. This intercritical treatment was

compared to the subcritical treatment using SEM photomicrographs and computer-aided image

analysis.

The intercritical heat treatment was initiated with a linear rise in temperature at

3.2 °C/min (5.7 °F/min) and holding for six hours at 746 °C (1374 °F) and slow cooling at

0.10 °C/min (0.32 °F/min) to 649 °C (1200 °F) for nine hours. Hence a total of 34 hours was used

for this heat treatment. The intercritical samples were then air cooled from 649 °C.

Figure A1 shows SEM micrographs of the carbide rich regions of the intercritically

annealed samples and a subcritical annealed sample. Figure A1(a)-(c) show SEM micrographs of

the carbide-rich regions of intercritical annealed HR, CR, and Norm steels. The carbides of the

intercritically annealed steels appear similar. The carbide morphology consists of coarse pearlite

and some small spheroidized carbides. The intercritically annealed samples still appear to be in

the carbide breakup stage of spheroidization. Figure A1(d) shows an SEM micrograph of the

carbide-rich regions in the CR steel after twenty hours of subcritical heat treatment. The twenty

hour sample was had the most similar heat treatment time to the intercritically annealed samples.

The percentage of spheroidization for the intercritically annealed CR steel was 27% spheroidized

and the percentage of spheroidization on the CR steel after twenty hours was 94%. Since the

intercritically annealed steels had such low percentages of spheroidization, the subcritical heat

treatment was chosen for the present study.

97

(a) (b)

(c) (d) Figure A1 SEM micrographs of the 15MnCr5 steel after spheroidization heat treatments.

(a) intercritically annealed HR steel, (b) intercritically annealed CR steel (c) intercritically annealed Norm steel, and (d) subcritically annealed CR steel after 20 hours.

98

APPENDIX B

LOGNORMAL STATISITICS

This appendix discusses the use of lognormal statistics in the analysis of the spheroidized

particles. Many particle distributions have lognormal type distributions. In order to test which

type of distribution the particles displayed, probability plots were generated. Figure B shows the

normal and lognormal probability plots for the CR steel after one hour of heat treatment at 692 °C

(1277 °F). The probability plot should form a straight line if the statistics fit well with that

distribution. In addition the Anderson-Darling (AD) coefficient shows how well the probability

plot fits the distribution. The lower the AD coefficient is the better fit for a distribution. Figure

B1(a) shows the probability plot for a normal statistical distribution. The probability plot for the

normal distribution forms a curved line and has an AD coefficient of 476. Figure B1(b) shows the

probability plot for a lognormal distribution. The probability plot forms a much straighter line and

had an AD coefficient of 89.

99

(a)

(b)

Figure B1 Probability plots for the CR 15MnCr5 steel after one hour of heat treatment at

692 °C (1277 °F) (a) normal distribution, (b) lognormal distribution.

100

APPENDIX C

COMMERCIALLY SPHEROIDIZED 16MNCR5

The commercially spheroidized 16MnCr5 steel was produced by Gerdau MACSTEEL

and had a composition shown in Table C1. The chemistry is very similar to that of the

experimental material. Figure C1 shows the microstructure of the commercially spheroidized

16MnCr5 steel. The microstructure of the spheroidized 16MnCr5 material consisted of ferrite

with large carbides between the ferrite grains and smaller carbides located inside prior pearlite

colonies. Table C2 summarizes the image analysis results and mechanical properties of the

commercially spheroidized steel. It should be noted the image analysis on this material was done

using only 2500 particles. The commercially spheroidized steel has a larger average spheroidized

particle area than any of the steels used in the present study. The reduction in area values for the

commercially spheroidized steel are similar to those of the HR and CR steels after twenty hours

of heat treatment indicating similar workability. The UTS for the commercially spheroidized steel

is approximately 50 MPa (7.3 ksi) lower than the twenty hours heat treated steels used in this

study. Similarly, the yield strength for the commercially spheroidized steel is approximately

100 MPa (14.5 ksi) lower than the steels heat treated for twenty hours.

Table C1 Chemical Composition in wt % of the Commercially Spheroidized 16MnCr5 steel.

C Mn P S Si Ni Cr Mo Cu Al

0 19 1 15 0 011 0 025 0 21 0 06 1 06 0 02 0 15 0 027

101

(a) (b) Figure C1 Micrographs of the commercially spheroidized 16MnCr5 Steel, picral etch (a)

light optical micrograph (b) SEM micrograph

Table C2 Image Analysis Results and Tensile Test Data for the Commercially

Spheroidized 16MnCr5 Steel.

Avg. Spheroidized Particle Area

(µm²)

% Spheroidized

Reduction in Area

(%)

Total Elongation

(%)

Uniform Elongation

(in/in)

UTS (MPa)

YS (MPa)

1.073 75.1 73.7 32.7 0.186 492 306

102

APPENDIX D

UNIFORM ELONGATION MEASUREMENT

Uniform elongation can be measured as the strain occurring at the maximum nominal

stress on the stress – strain curve. Considère’s construction can also be used to find the maximum

stress on a stress – strain curve. Considère’s construction can be used to find the uniform

elongation from tensile test data. Considère’s construction states the maximum load on a true

stress – engineering strain curve occurs when

eded

+=

1σσ

( D1)

where σ is the true stress and e is the engineering strain. [20] Table D1 shows the values for

uniform elongation measured by using the maximum nominal stress and by using Considère’s

construction. The difference in elongation measured is 0.001 in/in or less. The nominal method

was chosen to measure uniform elongation because of its ease of calculation.

Table D1 Uniform elongation values for the as-received 16MnCr5 steel using a nominal load method and Considère’s construction.

Nominal Considère's Construction

Uniform Elongation Std Dev Uniform

Elongation Std Dev Δ Elongation

HR 0.102 0.002 0.103 0.001 -0.001 CR 0.128 0.006 0.128 0.006 0.000

Norm 0.152 0.001 0.152 0.001 0.000

103

104

APPENDIX E

“U” SAMPLE BEND TEST

This appendix discusses the use of “U” shaped bend specimens in order to determine

workability. In order to obtain additional circumferential strain on the compression specimens, a

hole was drilled in one of the compressed samples of the commercially spheroidized material.

This compression sample was compressed to 445 kN (100 ksi) at a rate of 2 mm/min

(0.075 in/min). A 14.3 mm (0.563 in) diameter hole was drilled in the center of the sample and

then the sample was cut in half to create two “U” shaped samples. These samples were then

compressed in a vice incrementally and examined for surface cracking. The samples were

compressed until the ends of the “U” shaped sample met. The addition strain this test added was

approximately 0.31 in/in. The total circumferential strain imparted to these samples would then

be 1.17 in/in. Figure E1 shows photographs of the “U” samples after the bend tests. Figure E1(a)

shows a test sample that was bent at room temperature and Figure E1(b) shows a sample bent at

0 °C (32 °F). Figure E1(c) shows a magnified photograph of the surface of the room temperature

sample. Small microcracks can be seen in the sample but no large cracks can be seen without the

use of microscopes. Microcracks are not used to measure workability. Therefore the “U” sample

bend test was not used to characterize the cold workability of these steels in this study.

105

(a) (b) (c) Figure E1 Photographs of the region of maximum bending for the compressed “U” samples

15MnCr5 steel. (a) compressed at room temperature, (b) compressed at 0 °C (32 °F), and (c) magnified photograph of the room temperature sample showing microcracks.

106

APPENDIX F

CHARPY TESTING

In order to effectively increase the strain rate of the workability tests performed, Charpy

impact testing was evaluated. Both full-size and sub-size Charpy samples were made from the

commercially spheroidized 15MnCr5 steel and tested at room temperature. Figure F1 shows the

results of the Charpy tests. The Charpy samples did not break completely and therefore the results

were un-useable. The samples do show the very ductile nature of this spheroidized steel. Figure

F2 shows a photograph of the fracture surface of one of the full-size samples. Large shear

fractures can be seen on the surface. The fracture is 100% ductile.

Figure F1 Photograph of the broken full-size and sub-size Charpy samples for the commercially spheroidized 16MnCr5 steel.

107

Figure F2 Photograph of the fracture surface of the full-size Charpy sample of commercially spheroidized 16MnCr5 steel.

108

APPENDIX G

COCKCROFT AND LATHAM FRACTURE CRITERION

In order to further measure workability in the three steels, a modified Cockcroft Latham

criterion was evaluated. The area under the engineering stress – engineering strain curve was

measured to approximate a Cockcroft Latham coefficient. This was done to avoid using the

Bridgman correction to correct for necking behavior in the true stress – true strain curve.

Figure G1 shows the calculated Cockcroft Latham coefficient at various times for the three steels.

The values all converge toward a single value at twenty hours of spheroidization. The increasing

and decreasing trends taking place in these curves are not fully understood.

109


110

140

Plas

160

180

200tic

Stra

in E

nerg

y D

ensi

ty (M

J/m

3 )

3.2

Plas

3.6

4

tic S

train

Ene

rgy

Den

sity

(106 f

t. lb/

ft3 )


140

Plas

160

180

200

tic S

train

Ene

rgy

Den

sity

(MJ/

m3 )

3.2

Plas

3.6

4

ttic

Sra

in E

nerg

y D

ensi

ty (1

06 ft. l

b/ft3 )

(a) (b)


140

160

180

200

Plas

tic S

train

Ene

rgy

Den

sity

(MJ/

m3 )

3.2

3.6

4

Plas

tic S

train

Ene

rgy

Den

sity

(106 f

t. lb/

ft3 )


140

160

180

200Pl

astic

Stra

in E

nerg

y D

ensi

ty (M

J/m

3 ) MicrostructureHRCRNorm

)

3.2

3.6

4

Plas

tic S

train

Ene

rgy

Den

sity

(106 f

t. lb/

ft3

(c) (d) Figure G1 Cockcroft Latham coefficient for the 16MnCR5 steel heat treated at 692 °C

(1277 °F). (a) HR microstructure, (b) CR microstructure, (c) Norm microstructure, and (d) all microstructures.

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