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
ii
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.
iii
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
iv
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
v
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
vi
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
vii
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
(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. . . . . . . . . . . . . . . . . . . . . . . . 44
Figure 4.4
Micrographs of carbide-rich regions in 16MnCr5 CR 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. . . . . . . . . . . . .45
Figure 4.5
Micrographs of 16MnCr5 Norm 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. . . . . . . . . . . . . . . . . . . . . . . . 52
Figure 4.6
Micrographs of carbide-rich regions in 16MnCr5 Norm 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. . . . . . . . . . . . .53
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
viii
Figure 4.8
Histograms of the particle area for various heat treatment times for the 15MnCr5
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
Figure 4.9
Histograms of the particle area for various heat treatment times for the 15MnCr5
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
(1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm
microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .64
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
ix
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
(1277 °F) heat treatment. (a) HR microstructure, (b) CR microstructure, (c) Norm
microstructure, and (d) all microstructures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . .76
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
16MnCr5 steel for various prior microstructures. (a) HR microstructure, (b) CR
microstructure, (c) Norm microstructure, (d) all microstructures. . . . . . . . . . . . 85
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
16MnCr5 steel for various prior microstructures. (a) HR microstructure, (b) CR
microstructure, (c) Norm microstructure, (d) all microstructures. . . . . . . . . . . . 88
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
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.
xiii
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
Figure 2.1:
(a)
(b)
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]
Figure 2.2:
(a)
(b)
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 heightto-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 capillaryinduced 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
ln
a c 2γv m
=
ae
RTr
9
(2.1)
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 capillaryinduced 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
ln
2γv m
Cm
=
Ce
RTrcrit
(2.2)
Solving Equation (2.2) for the critical radius yields
rcrit
⎛ 2γv m
=⎜
⎝ RT
⎞⎛ C m
⎟⎜⎜ ln
⎠⎝ C e
⎞
⎟⎟
⎠
−1
(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 cementiteferrite interface, the Lifshitz-Wagner-Greenwood theory predicts a radius squared dependency as
given by
r 2 − ro2 =
8γΩDkC e
*t
9 RT
(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-WagnerGreenwood theory predicts a radius cubed dependency represented by [2]
r 3 − ro3 =
8γΩDC e
*t
9 RT
(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
Figure 2.9
Schematic of the fault migration theory of spheroidization breakup. [7]
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 ironcarbon 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]
xc Dceff =
xc Dc x Fe DFe
⎛ VFe3C − 3VFe
⎜⎜
VFe
⎝
18
2
⎞
⎟⎟ x c Dc + x Fe DFe
⎠
(2.6)
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]
Activation Energy kJ/mole
System
(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 ( S s ), and the mean thickness of unspheroidized particles
( x ). The volume fraction of spheroidized particles is represented by [14]
Vs = N s S s
(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]
dN s dVs 1
≅
dt
dt a x 2
(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]
dN s
= CVu
dt
(2.9)
Atasoy combined Equation (2.8) with Equation (2.9) to form [11]
2
dVs
= CVu a x
dt
20
(2.10)
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]
2
Vu = AVc exp(−Ca x t )
(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]
(
2
Vc 1
= exp Ca x t
Vu A
Figure 2.12
)
(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]
2
Atasoy defined the spheroidization rate k as equal to Ca x t . 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
εf
∫ σ * dε = C
(2.13)
0
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]
ε θ = ln[(R + t ) / (R + t / 2 )]
(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
Figure 3.1
(a)
(b)
(c)
(d)
(e)
(f)
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)
Figure 3.3
(b)
(c)
(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
Figure 4.1
(a)
(b)
(c)
(d)
(e)
(f)
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.
42
Figure 4.2
(a)
(b)
(c)
(d)
(e)
(f)
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
Figure 4.3
(a)
(b)
(c)
(d)
(e)
(f)
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
Figure 4.4
(a)
(b)
(c)
(d)
(e)
(f)
Micrographs of carbide-rich regions in 16MnCr5 CR 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.
45
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.
% Spheroidized =
AT − AS
* 100
AT
(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
Figure 4.5
(a)
(b)
(c)
(d)
(e)
(f)
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
Figure 4.6
(a)
(b)
(c)
(d)
(e)
(f)
Micrographs of carbide-rich regions in 16MnCr5 Norm 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.
53
30
30
1200
10
20
800
10
400
400
0
0
Frequency
800
Percentage of Total (%)
20
Frequency
Percentage of Total (%)
1200
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
0
(a)
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
(b)
30
30
3000
10
0
1000
0
20
800
10
400
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
Spheroidized Particle Area (µm2)
0
0
(c)
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
Spheroidized Particle Area (µm2)
(d)
30
30
2000
10
400
20
1500
Frequency
800
Frequency
20
Percentage of Total (%)
1200
Percentage of Total (%)
Frequency
2000
Percentage of Total (%)
20
Frequency
Percentage of Total (%)
1200
1000
10
500
0
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
Figure 4.7
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
(e)
(f)
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
30
2000
10
0
1000
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
3000
20
2000
Frequency
20
Percentage of Total (%)
3000
Frequency
Percentage of Total (%)
30
10
0
1000
0
(a)
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
(b)
30
30
3000
20
2000
10
1000
10
0
1000
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
Spheroidized Particle Area (µm2)
Frequency
2000
Percentage of Total (%)
20
Frequency
Percentage of Total (%)
3000
0
0
(c)
30
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
Spheroidized Particle Area (µm2)
(d)
30
3000
10
1000
20
1200
800
Frequency
2000
Percentage of Total (%)
20
Frequency
Percentage of Total (%)
1600
10
400
0
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
Figure 4.8
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
(e)
(f)
Histograms of the particle area for various heat treatment times for the 15MnCr5
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
0
30
30
800
10
20
1200
800
10
400
0
400
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
0
0
(a)
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
(b)
30
30
1200
10
400
0
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
Spheroidized Particle Area (µm2)
20
800
10
400
0
0
(c)
(d)
1200
800
10
400
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
Figure 4.9
Percentage of Total (%)
20
Frequency
Percentage of Total (%)
1200
20
800
10
400
0
0
0
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
2
Spheroidized Particle Area (µm )
(e)
(f)
Histograms of the particle area for various heat treatment times for the 15MnCr5
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
0.04 0.08 0.12 0.16 0.2 0.24 0.28 0.32 0.36 0.4
Spheroidized Particle Area (µm2)
30
30
0
Frequency
800
Percentage of Total (%)
20
Frequency
Percentage of Total (%)
1200
Frequency
0
Frequency
1200
Percentage of Total (%)
20
1600
Frequency
Percentage of Total (%)
1600
0.03
Spheroidized Particle Area (µm2)
Spheroidized Particle Area (µm2)
0.03
0.02
0.01
0
0
4
8
12
Holding Time (Hours)
16
0.02
0.01
0
20
0
4
(a)
Spheroidized Particle Area (µm2)
Spheroidized Particle Area (µm2)
20
0.03
0.02
0.01
0
Figure 4. 10
4
8
12
Holding Time (Hours)
16
20
0.02
Microstructures
HR
CR
Norm
0.01
0
0
4
8
12
Holding Time (Hours)
16
(c)
(d)
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
16
(b)
0.03
0
8
12
Holding Time (Hours)
20
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
Holding
Time
(Hours)
0.003
1
2
4
6
10
20
CR
Norm
Avg.
Area
-(µm2)
+(µm2)
Avg.
Area
-(µm2)
+(µm2)
0.015
0.014
0.023
0.018
0.020
0.033
0.034
0.009
0.005
0.015
0.013
0.013
0.022
0.022
0.024
0.040
0.047
0.040
0.043
0.061
0.066
0.008
0.012
0.010
0.019
0.017
0.018
0.027
0.005
0.008
0.006
0.013
0.011
0.012
0.019
0.012
0.022
0.015
0.044
0.033
0.034
0.059
Avg.
Area
-(µm2)
+(µm2)
0.015
0.011
0.012
0.016
0.014
0.020
0.009
0.006
0.007
0.010
0.009
0.013
0.020
0.012
0.017
0.029
0.022
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
100
100
80
80
Percent Spheroidized
Percent Spheroidized
microstructures.
60
40
40
20
20
0
60
0
4
8
12
Holding Time (Hours)
16
0
20
0
4
100
80
80
60
40
20
0
Figure 4. 11
4
8
12
Holding Time (Hours)
16
20
40
Microstructures
HR
CR
Norm
0
0
4
8
12
Holding Time (Hours)
16
(c)
(d)
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
20
60
20
0
16
(b)
100
Percent Spheroidized
Percent Spheroidized
(a)
8
12
Holding Time (Hours)
20
3
ln(Vc/Vu)
2
1
0
Figure 4.12
Microstructure
HR
CR
Norm
0
400
800
Holding Time (Seconds)
ln(Vc/Vu) with respect to time for the 15MnCr5 steels subcritically spheroidized
at 692 °C (1277 °F)
60
1200
260
260
85
200
80
180
160
4
8
12
Holding Time (Hours)
16
90
220
85
200
80
180
75
0
240
Hardness (HRB)
90
220
Carbide Microhardness (VHN)
95
Hardness (HRB)
Carbide Microhardness (VHN)
95
240
160
20
75
0
4
8
12
Holding Time (Hours)
(a)
20
(b)
260
260
Microstructures
HR
CR
Norm
Hardness (HRB)
90
220
85
200
80
180
160
4
Figure 4. 13
8
12
Holding Time (Hours)
16
95
90
220
85
200
80
180
75
0
240
Hardness (HRB)
240
Carbide Microhardness (VHN)
95
Carbide Microhardness (VHN)
16
160
20
75
0
4
8
12
Holding Time (Hours)
16
20
(c)
(d)
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
Holding
Time
(Hours)
0.003
1
2
4
6
10
20
CR
Norm
Hardness
(HV)
Std Dev
(HV)
Hardness
(HV)
Std Dev
(HV)
Hardness
(HV)
Std Dev
(HV)
231
206
191
190
189
182
175
4
7
5
5
5
5
4
232
218
223
206
192
175
173
10
12
7
10
8
8
6
254
213
211
206
175
180
177
10
7
13
7
6
3
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
170
160
150
80
140
Ferrite Microhardness (HV)
84
Hardness (HRB)
130
84
150
80
140
130
76
120
0
4
8
12
Holding Time (Hours)
16
76
120
20
0
4
8
12
Holding Time (Hours)
(a)
170
160
150
80
140
Ferrite Microhardness (HV)
84
Hardness (HRB)
Ferrite Microhardness (HV)
160
130
84
150
80
140
Microstructures
HR
CR
Norm
130
76
0
4
Figure 4. 14
Table 4.3
8
12
Holding Time (Hours)
16
120
20
0
4
8
12
Holding Time (Hours)
76
16
20
(c)
(d)
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.
Hardness in the Ferrite Regions for the 15MnCr5 Steel After Various Heat
Treatment Times at 692 °C (1277 °F).
HR
Holding
Time
(Hours)
0.003
1
2
4
6
10
20
CR
Norm
Hardness
(HV)
Std Dev
(HV)
Hardness
(HV)
Std Dev
(HV)
Hardness
(HV)
Std Dev
(HV)
165
146
151
155
140
161
147
8
3
4
9
4
2
3
155
147
148
143
138
137
142
7
6
4
3
2
8
3
127
126
123
139
129
139
131
4
3
5
4
3
3
3
64
20
(b)
170
120
16
Hardness (HRB)
Ferrite Microhardness (HV)
160
Hardness (HRB)
170
90
85
85
Hardness (HRB)
Hardness (HRB)
90
80
75
70
65
80
75
70
0
4
8
12
Holding Time (Hours)
16
65
20
0
4
(a)
Microstructures
HR
CR
Norm
Hardness (HRB)
85
80
75
70
80
75
70
0
4
Figure 4. 15
8
12
Holding Time (Hours)
16
20
65
0
4
8
12
Holding Time (Hours)
16
(c)
(d)
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
20
90
85
Hardness (HRB)
16
(b)
90
65
8
12
Holding Time (Hours)
20
Table 4.4
Macrohardness for the 15MnCr5 Steel After Various Heat Treatment Times at
692 °C (1277 °F).
HR
Holding
Time
(Hours)
0.003
1
2
4
6
10
20
CR
Norm
Hardness
(HRB)
Std Dev
(HRB)
Hardness
(HRB)
Std Dev
(HRB)
Hardness
(HRB)
Std Dev
(HRB)
87
83
79
79
76
80
79
1
2
1
2
1
1
1
82
79
79
77
79
76
76
1
1
1
1
1
1
1
78
75
75
70
69
75
73
1
1
1
2
1
1
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
ΔT =
ησdε
ρC p
(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).
CR
HR
Norm
As-Received
Delta T Std Dev
(°C)
(°C)
41.3
1.0
39.8
1.6
45.2
1.4
20 Hours
Delta T Std Dev
(°C)
(°C)
44.7
2.7
43.6
1.2
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 asreceived 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.
As-Recieved
20 hrs
6 hrs
40
200
0
0
0.1
0.2
0.3
Engineering Strain (in/in)
0
0.4
80
As-Recieved
400
6 hrs
40
20 hrs
200
0
0
0.1
0.2
0.3
Engineering Strain (in/in)
(a)
0
0.4
(b)
80
400
As-Recieved
6 hrs
20 hrs
200
0
Figure 4.17
0
0.1
0.2
0.3
Engineering Strain (in/in)
0
0.4
(c)
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
40
Engineering Stress (ksi)
Engineering Stress (MPa)
600
Engineering Stress (ksi)
400
Engineering Stress (MPa)
600
80
Engineering Stress (ksi)
Engineering Stress (MPa)
600
74
72
72
Reduction in Area (%)
Reduction in Area (%)
74
70
68
66
64
70
68
66
0
4
8
12
Holding Time (Hours)
16
64
20
0
4
74
72
72
70
68
0
4
Figure 4.18
8
12
Holding Time (Hours)
20
70
68
Microstructures
HR
CR
Norm
66
66
64
16
(b)
74
Reduction in Area (%)
Reduction in Area (%)
(a)
8
12
Holding Time (Hours)
16
20
64
0
4
8
12
Holding Time (Hours)
16
20
(c)
(d)
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
0.18
0.16
0.16
Uniform Elongation (in/in)
Uniform Elongation (in/in)
0.18
0.14
0.12
0.1
0
4
8
12
Holding Time (Hours)
16
0.14
0.12
0.1
20
0
4
0.18
0.16
0.16
0.14
0.12
0.1
0
Figure 4.19
4
8
12
Holding Time (Hours)
16
20
(b)
0.18
Uniform Elongation (in/in)
Uniform Elongation (in/in)
(a)
8
12
Holding Time (Hours)
16
20
0.14
Microstructures
HR
CR
Norm
0.12
0.1
0
4
8
12
Holding Time (Hours)
16
20
(c)
(d)
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
36
34
34
Percent Elongation (%)
Percent Elongation (%)
36
32
30
28
26
32
30
28
0
4
8
12
Holding Time (Hours)
16
26
20
0
4
36
34
34
32
30
0
Figure 4.20
4
8
12
Holding Time (Hours)
16
20
Microstructures
HR
CR
Norm
30
26
0
4
8
12
Holding Time (Hours)
16
20
(c)
(d)
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
20
32
28
28
26
16
(b)
36
Percent Elongation (%)
Percent Elongation (%)
(a)
8
12
Holding Time (Hours)
Table 4.6
Uniform and Non-Uniform Elongation Values for the 15MnCr5 Steel at Various
Heat Treatment Times at 692 °C (1277 °F).
Norm
HR
CR
Norm
CR NonHR NonNonUniform
Uniform
Uniform
Uniform
Uniform
Uniform
Elongation Elongation Elongation Elongation Elongation
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.
560
80
76
4
8
12
Holding Time (Hours)
16
88
600
84
560
80
520
20
76
0
4
(a)
84
560
80
Ultimate Tensile Strength (MPa)
88
600
Ultimate Tensile Strength (ksi)
Ultimate Tensile Strength (MPa)
640
92
76
0
Figure 4.21
4
8
12
Holding Time (Hours)
16
20
92
Microstructures
HR
CR
Norm
600
88
84
560
80
520
20
76
0
4
8
12
Holding Time (Hours)
16
(c)
(d)
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
16
(b)
640
520
8
12
Holding Time (Hours)
20
Ultimate Tensile Strength (ksi)
0
92
Ultimate Tensile Strength (ksi)
84
Ultimate Tensile Strength (MPa)
88
600
520
640
92
Ultimate Tensile Strength (ksi)
Ultimate Tensile Strength (MPa)
640
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
500
500
72
460
64
440
420
68
460
64
440
420
60
0
4
8
12
Holding Time (Hours)
16
60
400
20
0
4
(a)
16
500
72
72
Microstructures
HR
CR
Norm
68
460
64
440
Upper Yield Strength (MPa)
480
Upper Yield Strength (ksi)
480
420
64
440
420
0
Figure 4.22
4
8
12
Holding Time (Hours)
16
60
400
20
0
4
8
12
Holding Time (Hours)
16
20
(c)
(d)
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
68
460
60
400
20
(b)
500
Upper Yield Strength (MPa)
8
12
Holding Time (Hours)
Upper Yield Strength (ksi)
400
Upper Yield Strength (ksi)
68
Upper Yield Strength (MPa)
480
Upper Yield Strength (ksi)
Upper Yield Strength (MPa)
480
72
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
74
74
72
72
Reduction in Area (%)
Reduction in Area (%)
minima in the total elongation curves.
70
68
66
20
70
68
40
60
80
Area Percent Spheroidized (%)
66
20
100
40
60
80
Area Percent Spheroidized (%)
(b)
74
74
72
72
Reduction in Area (%)
Reduction in Area (%)
(a)
70
68
66
64
20
Figure 5.1
100
68
64
20
Microstructures
HR
CR
Norm
40
60
80
Area Percent Spheroidized (%)
100
(c)
(d)
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
70
66
40
60
80
Area Percent Spheroidized (%)
100
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.
Ferrite Hardness (HRB)
80
84
76
36
Total Elongation (%)
36
Total Elongation (%)
Ferrite Hardness (HRB)
80
84
76
32
28
24
120
130
140
150
Ferrite Hardness (HV)
160
32
28
24
120
170
130
140
150
Ferrite Hardness (HV)
(a)
Ferrite Hardness (HRB)
80
84
76
36
Total Elongation (%)
36
Total Elongation (%)
170
(b)
Ferrite Hardness (HRB)
80
84
76
160
32
28
24
120
Figure 5.2
130
140
150
Ferrite Hardness (HV)
160
170
32
28
24
120
Microstructures
HR
CR
Norm
130
140
150
Ferrite Hardness (HV)
160
170
(c)
(d)
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.
640
84
80
20
40
60
80
Area Percent Spheroidized (%)
76
100
88
600
84
560
80
20
40
60
80
Area Percent Spheroidized (%)
(a)
(b)
640
88
600
84
560
80
20
Figure 5.3
40
60
80
Area Percent Spheroidized (%)
76
100
Ultimate Tensile Strength (MPa)
92
Ultimate Tensile Strength (ksi)
Ultimate Tensile Strength (MPa)
640
92
88
600
Microstructures
HR
CR
Norm
84
560
80
20
40
60
80
Area Percent Spheroidized (%)
76
100
(c)
(d)
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
76
100
Ultimate Tensile Strength (ksi)
560
92
Ultimate Tensile Strength (ksi)
88
600
Ultimate Tensile Strength (MPa)
92
Ultimate Tensile Strength (ksi)
Ultimate Tensile Strength (MPa)
640
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.
560
Microstructures
HR
CR
Norm
26
28
Figure 5.4
30
32
Total Elongation (%)
80
34
76
36
Ultimate Tensile Strength (MPa)
84
Ultimate Tensile Strength (ksi)
Ultimate Tensile Strength (MPa)
88
600
88
600
84
560
Microstructures
HR
CR
Norm
64
66
68
70
Reduction in Area (%)
80
72
Ultimate Tensile Strength (ksi)
92
92
76
74
(a)
(b)
(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.
72
72
60
Upper Yield Strength (MPa)
64
440
Upper Yield Strength (ksi)
64
440
60
400
400
20
68
40
60
80
Area Percent Spheroidized (%)
56
100
20
40
60
80
Area Percent Spheroidized (%)
(a)
(b)
72
72
480
64
60
Figure 5.5
40
60
80
Area Percent Spheroidized (%)
56
100
68
64
440
60
20
40
60
80
Area Percent Spheroidized (%)
(c)
(d)
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
Microstructures
HR
CR
Norm
400
400
20
Upper Yield Strength (MPa)
68
Upper Yield Strength (ksi)
Upper Yield Strength (MPa)
480
440
56
100
56
100
Upper Yield Strength (ksi)
Upper Yield Strength (MPa)
68
Upper Yield Strength (ksi)
480
480
72
72
64
440
Microstructures
HR
CR
Norm
400
26
Figure 5.6
28
30
32
Total Elongation (%)
60
68
400
34
56
36
64
440
64
Microstructures
HR
CR
Norm
66
68
70
Reduction in Area (%)
60
72
Upper Yield Strength (ksi)
68
Upper Yield Strength (MPa)
480
Upper Yield Strength (ksi)
Upper Yield Strength (MPa)
480
56
74
(a)
(b)
(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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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)
Figure A1
(b)
(c)
(d)
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
Figure C1
Table C2
Avg.
Spheroidized
Particle Area
(µm²)
1.073
(a)
(b)
Micrographs of the commercially spheroidized 16MnCr5 Steel, picral etch (a)
light optical micrograph (b) SEM micrograph
Image Analysis Results and Tensile Test Data for the Commercially
Spheroidized 16MnCr5 Steel.
%
Spheroidized
Reduction
in Area
(%)
Total
Elongation
(%)
Uniform
Elongation
(in/in)
UTS
(MPa)
YS
(MPa)
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
σ
dσ
=
de 1 + e
( 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
HR
CR
Norm
Uniform
Elongation
0.102
0.128
0.152
Std Dev
0.002
0.006
0.001
Considère's Construction
Uniform
Elongation
0.103
0.128
0.152
103
Std Dev
0.001
0.006
0.001
Δ
Elongation
-0.001
0.000
0.000
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)
Figure E1
(b)
(c)
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
3.6
160
3.2
140
4
180
3.6
160
3.2
140
4
8
12
Holding Time (Hours)
16
20
0
4
180
3.6
160
3.2
Plastic Strain Energy Density (MJ/m )
4
200
20
Microstructure
HR
CR
Norm
3
200
Plastic Strain Energy Density (MJ/m3)
16
(b)
Plastic Strain Energy Density (106 ft.lb/ft3)
(a)
8
12
Holding Time (Hours)
140
4
180
3.6
160
3.2
140
0
Figure G1
4
8
12
Holding Time (Hours)
16
20
0
4
8
12
Holding Time (Hours)
16
20
(c)
(d)
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
Plastic Strain Energy Density (106 ft.lb/ft3)
0
Plastic Strain Energy Density (106 ft.lb/ft3)
180
200
Plastic Strain Energy Density (MJ/m3)
4
Plastic Strain Energy Density (106 ft.lb/ft3)
Plastic Strain Energy Density (MJ/m3)
200