"'.
Hot Deformation Behavior of
Magnesium AZ31
by
Geremi Vespa
Department of Mining, Metals and Materials Engineering
McGill University
Montreal, Quebec
December 2006
A thesis submitted to the Faculty of Graduate and Postdoctoral studies in
partial fulfillment of the requirements of the degree of Master of
Engineering
© G. Vespa, 2006
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Abstract
Automobile manufacturers are interested in lightweight materials, including magnesium,
to increase vehicle fuel economy, improve performance and reduce emissions. In this
work the deformation behavior of as-cast and rolled magnesium AZ31 alloy has been
studied. In as-cast material, it was found that reheating at 400°C and above for 60
minutes increased the homogeneity of the as-cast structure and gave rise to repeatable
deformation. At compression temperatures above 300°C dynamic recrystallization
occurred; below 200°C, there was significant twinning. Annealing completely
recrystallized the structure deformed below 200°C, but did not change the dynamically
recrystallized structure. AZ31 alloy was also rolled at temperatures of 350, 400 and
450°C and rolling speeds of 20 and 50 rpm for 15 and 30% reduction in thickness to
produce sheet. Before rolling, the alloy was preheated for 1 and 10 hours at the rolling
temperatures. The sheets were then tensile tested at 300, 400 and 450°C with strain rates
of 0.1, 0.01 and 0.001 S-l. The flow curves and microstructures indicated that the tensile
deformation mechanism changed with processing conditions. Two deformation
mechanisms were present in the magne sium sheet depending on the strain rate and grain
size. At slow strain rates and small grain size, the active deformation mechanism was
grain boundary sliding. As grain sizes increased there was also a component of
dislocation creep. At the fast strain rate, the deformation mechanism, regardless of grain
size, was dislocation creep. At a true strain rate of O.OOIS-I, it was found that rolling at
350°C with 30% reduction per pass yielded the finest microstructure and subsequently,
the best hot deformation characteristics. At a true strain rate of 0.1 S-l, rolling at 450°C
with 30% reduction per pass yielded a coarser, more recrystallized microstructure with
best hot deformation characteristics.
Résum.é
Afin d'augmenter l'efficacité énergétique, d'améliorer la performance et de réduire les
émissions de gaz, les fabricants d'automobiles s'intéressent aux matériaux légers, incluant le
magnésium. L'étude du comportement lors de la déformation d'un alliage de magnésium
AZ31 coulé et laminé est l'objet de ce projet. Dans le cas des pièces coulées, il a été
remarqué que l'homogénéité de la structure est améliorée lorsque les pièces sont réchauffées
à 400°C pendant plus de 60 minutes. Ce revenu permet aussi d'augmenter la reproductibilité.
Lors d'essais de compression à des températures supérieures à 300°C, une recristallisation
dynamique se produit et pour des températures inférieures à 200°C, des maclages significatifs
se produisent. Les pièces comprimées en deçà de 200°C recristallisent complètement lors
d'un recuit alors que celles déformées à plus de 300°C maintiennent leur structure obtenue
lors de la recristallisation dynamique. La production de tôles a été effectuée en laminant
l'alliage AZ31 à des températures de 350, 400 et 500°C, à des vitesses de 20 et 50 tpm et
avec
des réductions d'épaisseur de 15 et 30%.
Avant le laminage, les pièces ont été
préchauffées pendant 1 et 10 heures aux températures de laminage. Les tôles ont ensuite été
soumises à des essais de tractions à 300, 400 et 450°C aux taux de déformations de 0.1, 0.0 1,
0.001 s·\. Les courbes de déformations obtenues ainsi que les microstructures observées
indiquent un changement du mécanisme de déformation selon les conditions d'opération.
Deux mécanismes différents de déformation ont été observés dans les tôles de magnésium en
fonction du taux de déformation et de la taille des grains. Lorsque le taux de déformation est
bas et que la taille des grains est petite, le mode principal de déformation est le glissement
des joints de grains. Lorsque la taille des grains augmente, une composante de déformation
par mouvements de dislocations s'ajoute au glissement des joints de grains. Lors d'essais à
haut taux de déformation, le mécanisme de déformation principal est le mouvement de
dislocations et ce, peu importe la taille de grains. Avec un taux de déformation vraie de
0.001 s·\, les meilleures caractéristiques de déformation à chaud ont été obtenues sur des
pièces laminées à 350°C avec une réduction de 30% par passe. Avec un taux de déformation
vraie de O.ls-1, les meilleures caractéristiques de déformation à chaud ont été observées sur
des pièces laminées à 450°C avec une réduction de 30% par passe.
11
Acknowledgelllents
l would like to express my gratitude to my thesis supervisor, Professor Steve Yue, for
giving me the opportunity to work in such an interesting and engaging project and for his
constant guidance throughout the course of this work. l would also like to thank Ravi
Verma, Jon Carter and Paul KIajewski for their invaluable help throughout the project,
but more specifically during my three month stay at General Motors.
l am grateful to the Materials Technology Laboratory (CANMET) of Natural Resources
Canada in Ottawa and more specifically Claude Galvani, Amjad Javaid and Elhachmi
Essadiqi for their casting and rolling work. l would also like to thank Mihriban
Pekguleryuz, Faramarz Zarandi, Abdel Elwazri, Pierre Vermette, Edwin Femandez,
Barbara Hanley and Lorraine Mello. Additionally, many thanks to Luke Mackenzie for
the countless reviews of my thesis and to my fellow graduate students, especially Geoff,
Vikram, Ana, Umu, Sean, Ehab, Phuong, Graeme, Linda, Emilie, Lan, Lihong and
Etienne with whom l spent a great deal of time with.
Last but not least, l would like to thank my family: Mom & Dad, Alisa & Pat for their
continued love and support.
111
Table of Contents
ABSTRACT ........................................................................................................................ 1
RESUME .......................................................................................................................... II
ACKNOWLEDGEMENTS ........................................................................................... 111
1
INTRODUCTION..................................................................................................... 1
2
LITERATURE REVIEW ........................................................................................ 3
2.1
ELEMENTAL CHARACTERISTICS AND STRUCTURE ...............................................
3
2.1.1 Crystal Structure ............................................................................................. 3
2.1.2 Magnesium Characteristics ............................................................................ 4
2.2
ALLOYING AND ALLOY DESIGNATIONS ............................................................... 4
2.2.1 Physical Metallurgy of Magnesium Alloys ..................................................... 4
2.2.2 Alloying Elements ........................................................................................... 5
2.2.3 Alloy Designations .......................................................................................... 6
2.3
DEFORMATION ..................................................................................................... 8
2.3.1 Slip .................................................................................................................. 8
2.3.2 Twinning ....................................................................................................... 10
2.3.3 Additional Deformation Mechanisms ........................................................... 12
2.4
RECRYSTALLIZATION ......................................................................................... 14
2.4.1 Static Recrystallization ................................................................................. 15
2.4.2 Dynamic Recrystallization ............................................................................ 15
2.4.2.1 Twin Dynamic Recrystallization ....................................................... 18
2.4.2.2 Additional Mechanisms of Dynamic Recrystallization ................... 19
2.4.2.3 Effect of Grain Size on Dynamic Recrystallization ......................... 19
2.5
DEFORMATION CREEP ........................................................................................ 20
2.5.1 Deformation Creep at High Strain Rates and Low Temperature ... .............. 22
2.6
SHEET PRODUCTION ........................................................................................... 23
2.6.1 Rolling Parameters ....................................................................................... 24
3
EXPERIMENTAL PROCEDURE ........................................................................ 26
3.1
3.2
3.3
3.4
3.5
MATERIALS ........................................................................................................
HOT ROLLING ....................................................................................................
TENSILE TES TING ...............................................................................................
COMPRESSION TESTING .....................................................................................
CHARACTERIZATION ..........................................................................................
26
26
28
30
34
3.5.1 Optical Metallography .................................................................................. 34
3.5.1.1 Sam pie Preparation for Compression Samples ............................... 34
3.5.1.2 Sample Preparation for Tension Samples ........................................ 34
IV
~-,
3.5.2
3.5.3
4
Scanning Electron Microscopy ..................................................................... 35
Electron Backscattered Diffraction .............................................................. 36
RESULTS ................................................................................................................ 37
DEFORMATION OF AS-CAST AZ31 OF COMPOSITION A. ..................................... 37
4.1
4.1.1 As-cast Grain Size ......................................................................................... 37
4.1.2 Compression Behavior .................................................................................. 38
4.1.3 Microstructures after Deformation ............................................................... 38
4.2
EFFECT OF REHEAT TREATMENTS ON MICROSTRUCTURES AND DEFORMATION
BEHA VIOR OF AS-CAST AZ31 OF COMPOSITION B ........................................................ 39
4.2.1
4.2.2
4.3
Microstructures after Reheat Treatments ..................................................... 40
Deformation Response to Reheating Cycles ................................................. 42
TENSILE BEHA VIOR OF ROLLED AZ31 SHEET OF COMPOSITION C .................... 44
4.3.1 Tensile Curves ............................................................................................... 44
4.3.2 Microstructural Analysis .............................................................................. 47
4.3.2.1 Low Strain Rate (O.OOls-I) .................................................................. 47
4.3.2.2 High Strain Rate (O.ls-I) ..................................................................... 52
4.3.3 Effect ofProcessing Conditions on Yield Stress ........................................... 55
5
DISCUSSION .......................................................................................................... 60
5 .1
DEFORMATION OF AS-CAST AZ31 OF COMPOSITION A ...................................... 60
5.2
EFFECT OF HEAT TREATMENTS ON MICROSTRUCTURES AND DEFORMATION
BEHA VIOR OF AS-CAST AZ31 ........................................................................................ 61
5.3
DEFORMATION BEHAVIOR OF ROLLED AZ31 SHEET ......................................... 63
5.3.1
5.3.2
Tensile Curves ............................................................................................... 63
Microstructural Analysis .............................................................................. 65
5.3.2.1
Low Strain Rate (0.00Is- 1) .................................................................... 65
5.3.2.2
5.3.3
Processing Parameters ................................................................................. 68
Low Strain Rate (0.00Is- 1) .................................................................... 69
5.3.3.1
5.3.3.2
6
1
High Strain Rate (0.ls- ) ....................................................................... 67
1
High Strain Rate (0.ls- ) ....................................................................... 71
CONCLUSIONS ..................................................................................................... 73
REFERENCES ................................................................................................................ 75
v
List of Figures
FIGURE 2.1: HEXAGONAL CLOSE PACKED CRYSTAL STRUCTURE OF MAGNESIUM ................
3
FIGURE 2.2: BASAL PLANE IN A HEXAGONAL CLOSE PACKED LATTICE .................................
8
FIGURE 2.3: PRISMATIC PLANE IN A HEXAGONAL CLOSE PACKED LATTICE ..........................
9
2.4: SECOND ORDER PYRAMIDAL PLANE IN A HEXAGONAL CLOSE PACKED LATTICE .
..................................................................................................................................... 9
FIGURE
2.5: CRSS FOR VARIOUS DEFORMATION SYSTEMS AS A FUNCTION OF TEST
TEMPERATURE IN PURE MAGNESIUM [13] ................................................................... 10
FIGURE
FIGURE 2.6: TWINNING DIRECTION AND COMPONENTS
FIGURE
[13] ............................................... Il
2.7: TENSION, COMPRESSION AND DOUBLE TWINS [13]. ........................................ 12
2.8: SCHEMATIC DIAGRAM OF DEFORMATION BANDS, TRANSITION BANDS AND KINK
BANDS. THE REGION LABELED "T" IS A TRANSITION BAND, REGION A TO C TO A IS A
KINK BAND [22] ......................................................................................................... 13
FIGURE
FIGURE 2.9: THE DEVELOPMENT OF A "NECKLACE" MICROSTRUCTURE DURING DYNAMIC
RECRYSTALLIZATION [22] ..........................................................................................
16
FIGURE 2.10: THE PROCESS BY WHICH DYNAMIC RECRYSTALLIZA TI ON "NECKLACING"
FORMS SHEAR ZONES
[27] .......................................................................................... 17
FIGURE 2.11 : INFLUENCE OF INITIAL GRAIN SIZE ON THE DYNAMICALL Y RECRYST ALLIZED
GRAIN SIZE AND VOLUME PERCENT DYNAMIC RECRYSTALLIZATION. THE CONTINUOUS
LINE REFERS TO THE DYNAMICALL Y RECRYSTALLIZED GRAIN SIZE, WHILE THE
DISCONTINUOUS LINE IS A FIT OF DA TA ACCORDING TO THE VOLUME PERCENT
DYNAMIC RECRYSTALLIZATION
[13] .......................................................................... 20
FIGURE 2.12: MANTLE REGION WITHIN REGIONS ADJACENT TO GRAIN BOUNDARIES IN
SUPERPLASTIC MATERIALS [41] .................................................................................
21
FIGURE 2.13: EFFECTS OF ROLLING TEMPERA TURE AND REDUCTION ON DEFECTS OF AZ31 B
SHEETS ROLLED AT 2000 MlMIN [57] ......................................................................... 25
FIGURE 3.1: (A) STANAT MILL USED IN ROLLING EXPERIMENTS AND (B) SAND BATH USED
FOR HEAT TREATING THE SAMPLES BEFORE ROLLING .................................................
vi
27
FIGURE 3.2: (A) TENSILE MACHINE AND (B) A VIEW INTO THE FURNACE WHERE SAMPLES
ARE TESTED ................................................................................................................ 29
FIGURE 3.3: TENSILE SAMPLE DIMENSIONS USED FOR HOT TENSILE TESTING ..................... 29
FIGURE 3.4: MTS 100 USED FOR HOT COMPRESSION TESTING ............................................ 31
FIGURE 3.5: AZ31 ALLOY OF COMPOSITION B CAST IN A WATER-COOLED COPPER MOLD
USED FOR HEAT TREATMENT STUDY ........................................................................... 32
FIGURE 3.6: REHEATING FOLLOWED BY COMPRESSION SCHEDULE ..................................... 33
FIGURE 3.7: SCANNING ELECTRON MICROSCOPE USED IN THE HOT-TENSILE STUDY ........... 35
FIGURE 4.1: EBSD IMAGE OF AZ31 OF COMPOSITION A AND ASSOCIATED GRAIN SIZE
DISTRIBUTION ............................................................................................................ 37
FIGURE 4.2: COMPRESSION SAMPLES DEFORMED TO A TRUE STRAIN OF 0.4 AT A TRUE
STRAIN RATE OF 0.1 S-l ............................................................................................... 38
FIGURE 4.3: MICROSTRUCTURES OF COMPRESSION SAMPLES DEFORMED TO A TRUE STRAIN
1
OF 0.4 AT A TRUE STRAIN RATE OF 0.1s- AT TEMPERATURES OF (A) 350°C AND (B)
150°C AND AFTER AN ANNEALING STEP AT (C, D) 450°C FOR 5 MINUTES ................... 39
FIGURE 4.4: W ATER-COOLED COPPER PLATE SAMPLES (A) AS-RECEIVED AND REHEATED TO
(B) 400°C FOR 30 MINUTES AND (C) 450°C FOR 60 MINUTES ..................................... 40
FIGURE 4.5: WATER-COOLED COPPER PLATE SAMPLES REHEATED TO 400°C FOR 15,30,60
AND 120 MINUTES (A, D, G, J), 450°C FOR 15,30,60 AND 120 MINUTES (B, E, H, K) AND
480°C FOR 15, 30, 60 AND 120 MINUTES (C, F, l, L) 480°C FOR 120 MINUTES ............ 41
FIGURE 4.6: FLOW STRESS AFTER REHEA TING TO 400°C FOR 15 MINUTES AND DEFORMING
AT 350°C AT A TRUE STRAIN RATE OF 0.1 S-l TO A TRUE STRAIN OF 0.6 ...................... 42
FIGURE 4.7: FLOW STRESS AFTER REHEA TING TO 400°C FOR 60 MINUTES AND DEFORMING
AT 350°C AT A TRUE STRAIN RATE OF O.lS-1 TO A TRUE STRAIN OF 0.6 ...................... 43
FIGURE 4.8: SPREAD IN MCS vs. REHEA TING TIME FOR AZ31 REHEATED AT 400,450 AND
480°C ........................................................................................................................ 43
FIGURE 4.9: STRESS-STRAIN CURVES AFTER PROCESSING AT 400 AND 450°C WITH STRAIN
1
RATES OF 0.001, 0.01 AND 0.ls- . IN THE LEGEND, THE FIRSTNUMBER IS THE TESTING
TEMPERATURE, THE SECOND NUMBER IS THE STRAIN RATE ........................................ 44
FIGURE 4.10: YIELD STRESS VS. STRAIN RATE FOR SAMPLES TESTED AT 450, 400 AND
300°C ........................................................................................................................ 45
vu
FIGURE 4.11: STRESS VS. STRAIN FOR SPECIMENS WITH THE HIGHEST AND LOWEST YIELD
1
STRESSES TESTED AT 450°C AT STRAIN RATES OF 0.1 AND 0.001s· ••••••••••••••••••••••••••• 46
FIGURE 4.12: STRESS VS. STRAIN FOR SPECIMENS WITH THE HIGHEST AND LOWEST YIELD
1
STRESSES TESTED AT 400°C AT STRAIN RATES OF 0.1 AND 0.001s· ........................... 46
FIGURE 4.13: MICROSTRUCTURES OF LOWEST (LEFT SIDE) AND HIGHEST (RIGHT SIDE) YIELD
1
STRESS SAMPLES TESTED AT 450°C AT A STRAIN RATE OF 0.001s· •
MICROSTRUCTURES INCLUDE THE AS-ROLLED CONDITION (A, B), PRE-PULLING
CONDITION (C, D), AND THE FAILURE TIP (E, F, G, AND H). THE PROCESSING CONDITION
FOR THE LEFT SIDE IS 350°C + 1 HR REHEAT, 30% RED., 50 RPM. THE PROCESSING
CONDITION FOR THE RIGHT SIDE IS 450°C + 1 HR REHEAT, 15% RED., 20 RPM ............ 49
FIGURE 4.14: MICROSTRUCTURES OF LOWEST (LEFT SIDE) AND HIGHEST (RIGHT SIDE) YIELD
STRESS SAMPLES TESTED AT 400°C AT A STRAIN RA TE OF 0.001 S·I.
MICROSTRUCTURES INCLUDE THE AS-ROLLED CONDITION (A, B), PRE-PULLING
CONDITION (C, D), AND THE FAILURE TIP (E, F, G, AND H). THE PROCESSING CONDITION
FOR THE LEFT SIDE IS 350°C + 10 HR REHEAT, 30% RED., 50 RPM. THE PROCESSING
CONDITION FOR THE RIGHT SIDE IS 450°C +1 HR REHEAT, 15% RED., 20 RPM ............ 51
FIGURE 4.15: MICROSTRUCTURES OF LOWEST (LEFT SIDE) AND HIGHEST (RIGHT SIDE) YIELD
1
STRESS SAMPLES TESTED AT 450°C AT A STRAIN RATE OF 0.ls· • MICROSTRUCTURES
INCLUDE THE AS-ROLLED CONDITION (A, B), PRE-PULLING CONDITION (C, D), THE
FAlLURE TIP (E, F,) AND A GOOD DISTANCE A WAy FROM THE FAlLURE TIP IN THE
DEFORMATION ZONE (G, H). THE PROCESSING CONDITION FOR THE LEFT SIDE IS 450°C
+ 10 HR REHEAT, 30% RED., 20 RPM. THE PROCESSING CONDITION FOR THE RIGHT SIDE
IS 350°C +10 HR REHEAT, 15% RED., 20 RPM ............................................................. 53
FIGURE 4.16: MICROSTRUCTURES OF LOWEST (LEFT SIDE) AND HIGHEST (RIGHT SIDE) YIELD
STRESS SAMPLES TESTED AT 400°C AT A STRAIN RATE OF 0.1 S·I. MICROSTRUCTURES
INCLUDE THE AS-ROLLED CONDITION (A, B), PRE-PULLING CONDITION (C, D), THE
FAlLURE TIP (E, F,) AND A GOOD DISTANCE A WA y FROM THE FAILURE TIP IN THE
DEFORMA TI ON ZONE (G, H). THE PROCESSING CONDITION FOR THE LEFT SIDE IS 450°C
+10 HR REHEAT, 30% RED., 20 RPM. THE PROCESSING CONDITION FOR THE RIGHT SIDE
IS 350°C +1 HR REHEAT, 30% RED., 20 RPM ............................................................... 54
FIGURE 4.17: YIELD STRESS VS. STRAIN RATE AT400 AND 450°C FOR SAMPLES ROLLED AT
350°C, 10 HOUR REHEAT TIME, 50 RPM ROLL SPEED WITH REDUCTIONS PER PASS OF 15
(BLACK) AND 30% (RED) ............................................................................................ 56
FIGURE 4.18: YIELD STRESS VS. STRAIN RATE AT 300, 400 AND 450°C FOR SAMPLES
ROLLED AT 450°C, 10 HOUR REHEA T TIME, 20 RPM ROLL SPEED WITH REDUCTIONS PER
PASS OF 15 (BLACK) AND 30% (RED) .......................................................................... 56
V1l1
FIGURE 4.19: YIELD STRESS VS. STRAIN RATE AT 300,
400 AND 450°C FOR SAMPLES
1 HOUR REHEAT TIME, 30% REDUCTION PER PASS, A ROLL SPEED OF 20
RPM AT TEMPERATURES OF 350 (BLACK), 400 (BROWN) AND 450°C (RED) ................ 57
ROLLED WITH A
FIGURE 4.20: YIELD STRESS VS. STRAIN RA TE AT 300,
400 AND 450°C FOR SAMPLES
ROLLED AT 350°C, 1 HOUR REHEAT TIME, 30% REDUCTION PER PASS, WITH A ROLL
SPEED OF 20 (RED) AND 50 RPM (BLACK) ...................................................................
58
FIGURE 4.21: YIELD STRESS VS. STRAIN RA TE AT 400 AND
450°C FOR SAMPLES ROLLED AT
350°C, 10 HOUR REHEAT TIME, 15% REDUCTION PER PASS, WITH A ROLL SPEED OF 20
(RED) AND 50 RPM (BLACK) ....................................................................................... 59
FIGURE 4.22: YIELD STRESS VS. STRAIN RATE AT 300,
ROLLED AT 450°C,
(BLACK) AND
400 AND 450°C FOR SAMPLES
15% REDUCTION PER PASS, 20 RPM WITH A REHEA T TIME OF 1
10 HOURS (RED) ................................................................................... 59
5.1: SUBMICRON FIBERS VISIBLE PARALLEL TO THE TENSION AXIS IN SAMPLE WITH
VERY FINE GRAIN SIZE ................................................................................................ 66
FIGURE
IX
List of Tables
TABLE 2.1: TYPICAL PHYSICAL AND MECHANICAL PROPERTIES OF PURE MAGNESIUM [1]. .. 4
TABLE 2.2: CODE LETTERS FOR THE DESIGNATION SYSTEM OF MAGNESIUM ALLOYS [1]. .... 7
TABLE 2}: TEMPER DESIGNATIONS FOR MAGNESIUM ALLOYS
[1] ....................................... 7
TABLE 3.1: CHEMICAL COMPOSITION OF AZ3l MAGNESIUM ALLOYS ................................ 26
TABLE 3.2: HOT ROLLING CONDITIONS USED FOR PRODUCING AZ3l SHEET ..................... 27
TABLE 3.3: HOT ROLLING SCHEDULES FOR AZ3l ALLOY REHEA TED AND ROLLED AT 350,
400 AND 450°C. THE ROLLING TEMPERA TURE FOR EACH PASS IS THE SAME FOR A
GIVEN ROLLING SCHEDULE ........................................................................................ 28
TABLE 3.4: PROCESS PARAMETERS TESTED IN THIS STUDY ................................................ 30
TABLE 3.5: REHEA TING SCHEDULE EMPLOYED .................................................................. 33
TABLE 4.1: SUMMARY OF YIELD STRESSES AND ELONGATIONS AT DIFFERENT TESTING
TEMPERA TURES AND STRAIN RA TES USED DURING TENSION TESTING ......................... 47
TABLE 5.1: PROCESSING CONDITIONS YIELDING THE STUDIED MICROSTRUCTURES ........... 69
-x
Chapter 1
Introduction
Magnesium has the lowest density of the common structural metals, and is thus an
attractive proposition where weight is a primary consideration [1]. It is 36% lighter per
unit volume than aluminum and 78% lighter than steel [2]. Its high strength-to-weight
ratio, good machinability, weldability and damping characteristics make magnesium and
its alloys useful in many engineering applications. In contrast to steel and aluminum,
however, a very limited range of magne sium wrought products (especially flat rolled
material) is available, and this represents only 1% of the total annual magnesium usage
[3].
Recently, renewed interest in magnesium has arisen because of environmental concerns
due to emissions from the transportation industry. Magnesium's low density can help
increase vehic1e fuel economy, improve performance and reduce emissions [1,2,4].
Current automotive applications are limited to castings: gearboxes, valve covers, wheels,
c1utch housings, and brake pedal brackets. The usage of magnesium could be greatly
increased if it could be employed on a large scale in the wrought form for application in
the primary structures of automobiles.
At present, magnesium sheet is substantially more expensive than that of aluminum and
steel. There is also limited literature available on the fabrication of magne sium sheet. The
processing conditions required for the production of magnesium sheet with acceptable
properties (strength and elongation) is of great interest. At present, magnesium rolling
schedules inc1ude an extended homogenization treatment at elevated temperatures prior
to rolling. The slabs are then rolled with 12 - 18 passes at high temperatures [5]. This
1
production route is costly and increased knowledge of the response to rolling processing
conditions will help improve efficiency of production and therefore reduce costs. In this
work, the deformation response of as-cast magnesium alloy AZ31 is analyzed.
Subsequently, the effect of the as-rolled microstructure on the elevated temperature
tensile properties of AZ31 sheets is investigated.
The goal is to design hot-rolling
schedules that produce the microstructures suitable for forming techniques such as
superplastic forming.
The pertinent theories goveming the phenomena considered in this study are reviewed in
Chapter 2. Pure magnesium characteristics and structure will be reviewed. Alloying with
specific reference to the most common alloying elements will be discussed. Additionally,
there will be a strong focus on the deformation of magne sium and magnesium alloys with
sections on deformation modes, recrystallization and deformation creep. Lastly, a review
of the present research on magnesium sheet is presented.
Chapter 3 describes the various experimental techniques and materials used in this study.
The experimental procedure used for rolling, tension and compression is thoroughly
discussed. Characterization methods are also reviewed.
Results from the experiments are presented in Chapter 4 and are discussed in Chapter 5.
Lastly, conclusions obtained from this study are given in Chapter 6.
2
Chapter 2
Literature Review
2.1 Elemental Characteristics and Structure
2.1.1 Crystal Structure
The crystal structure of pure magne sium is hexagonal c10sed packed (HCP) (Figure 2.1).
1
c
J
HCP
Figure 2.1: Hexagonal close packed crystal structure of magnesium.
The c and a in this figure represent the lattice parameters (dimensions) of the unit cell.
The lattice parameters of pure magnesium at 25°C are: a
=
0.32092 nm and c
=
0.52105
nm ±0.01% [1]. In an HCP structure, the theoretical cfa ratio, according to geometry, is
1.633, but the actual value for magnesium at room temperature is 1.6236 [1].
3
CHAPTER 2 - LITERATURE REVIEW
2.1.2 Magnesium Characteristics
Table 2.1 highlights sorne key physical and mechanical properties of pure magnesium. As
will be seen in forthcoming sections, alloying can have a profound influence on
properties such as strength and ductility.
. 1 prope rfles 0 f plure magnesium [1].
Table 2.1: Typical ~Jlyslca
h . 1 an d mec h amca
Typical Physical/Mechanical Properties
1738 kg/m;1
Densit}!
Melting Point
Boiling Point
Elastic Modulus
Poisson's Ratio
Tensile Strength
Yield Strength
Percent Elongation
650°C
1090°C
44 GPa
0.35
90-220 MPa
21-140 MPa
2-15%
2.2 Alloying and Alloy Designations
Primary magnesium has a minimum purity of99.8% [1]. This purity is sufficient for most
chemical and metallurgical uses. Magnesium, however, is rarely used for engineering
applications without being alloyed with other elements to obtain the particular properties
required for industrial applications.
2.2.1 Physical Metallurgy of Magnesium Alloys
The key features that dominate the physical metallurgy of magne sium alloys are the
hexagonal lattice structure of solid magnesium and the fact that its atomic diameter
(0.320 nm) is such that it enjoys a favorable size factor for solid solubility with a diverse
range of solute elements [1]. Appreciable quantities of a solute may be accommodated in
solid solution only when the difference in atomic radii between the two atom types is less
than approximately 15%. Apart from size factor, several features of the solute and solvent
atoms determine the degree of solubility inc1uding: crystal structure, electronegativity,
and valences. When the crystal structure of both solute and solvent atoms is the same,
4
CHAPTER 2 - LITERATURE REVIEW
solid solubility occurs. Consequently, continuous solid solutions are only possible with
hexagonal metals, e.g. zinc, cadmium, beryllium, titanium and zirconium [6].
However, with increasing atomic size difference, the solute atoms create substantial
lattice distortions and a new phase eventually forms [7]. With increasing difference in
electronegativity between two elements, cornes a greater likelihood that these will form
an intermetallic compound instead of a substitutional solid solution [7]. Lastly, a metal
will have a higher tendency to dissolve in a metal of higher valency than in one of a
lower valency [6, 7].
2.2.2 Alloying Elements
Aluminum is the most common alloying element. It reduces as-cast grain size, improves
strength, hardness, castability, and widens the solidification range [1,6,8,9]. When
present in excess of 6wt% the alloy becomes commercially heat treatable, but
commercial alloys rarely exceed 10wt% aluminum. An aluminum content of 6wt% yields
the optimum combination of strength and ductility (as-cast) [1].
Zinc is often used in combination with aluminum to improve room temperature strength
[1]. Grain size in the as-cast condition is reduced markedly by the addition of zinc [6, 8].
Zinc also reduces the harmful corrosion due to iron and nickel impurities, but may
increase susceptibility to hot shortness [1].
Manganese increases the yield strength slightly, but has limited effect on the tensile
strength. Its most important function is to improve the salt water resistance of Mg-Al and
Mg-AI-Zn alloys by removing iron and other heavy elements as solutes into relatively
harmless intermetallic compounds [1].
Zirconium is a grain refiner used in castings; it produces an "ultra fine grain" size
comparable to other grain refining treatments [6, 8]. Tensile properties of rolled sheet are
----
increased with its addition [6]. Zirconium, however, cannot be added in alloys that
5
CHAPTER 2 - LITERA TURE REVIEW
contain aluminum or manganese because stable compounds with these elements are
formed and thus remove zirconium from solid solution [1].
Lithium has relatively high solubility in magnesium and, because of its low density it has
attracted interest as an alloying element. Eleven weight percent lithium is needed to form
the
~
phase, which has a body-centered cubic structure and therefore improves
formability. The addition of lithium decreases strength but increases ductility [1]. The
addition oflithium is, however, detrimental to corrosion properties [10].
Rare earth elements increase the strength of magnesium alloys at high temperatures.
Additionally, they narrow the freezing range of the alloys, thus minimizing crack welding
and porosity in castings [1]. Precipitation hardening is also possible with the addition of
rare earth elements [11]. The addition of Yttrium with other rare earth elements increases
creep resistance above 300°C [1]. Yttrium can also significantly increase the specific
strength [11].
Many elements adversely affect the corrosion resistance of magnesium including: copper,
iron, nickel, and silicon. Iron is one of the more harmful impurities in magne sium alloys
in that it greatly reduces the corrosion resistance if present even in small amounts
(>0.005wt%) [1]. Nickel also reduces the corrosion resistance. Silicon has been found to
increase the fluidity of magnesium in the molten state, but it decreases corrosion
resistance in the presence of iron [1].
2.2.3 Alloy Designations
The system used by the American Society for Testing and Materials (ASTM) is most
commonly used for designating magnesium alloys. This method is a three part letternumber-Ietter system.
The first letters indicate the two principal alloying elements (listed in order of decreasing
alloying content) [1]. The second part consists of weight percentages of these two
6
CHAPTER 2 - LITERATURE REVIEW
elements (rounded off to the nearest whole number and listed in the same order as the
code letters). The third part consists of an assigned letter (beginning with A) to
distinguish between alloys having the same nominal designation, or an "X" to indicate
that the alloy is still experimental. For example, AZ91X is an experimental magnesium
alloy containing approximately 9wt% aluminum and 1wt% zinc.
The ASTM system aiso includes a code system for the temper of magnesium alloys. This
consists of a letter plus one or more digits. The temper designation follows the alloy
designation and is separated by a hyphen. Tables 2.2 and 2.3 show the code Ietters for
both the alloying elements and temper designations, respectively.
Table 2.2: Code letters ~or t h e d eSlgnatlon
.
sys t em 0 f magnesium alloys [1].
Letter
Alloying Element
A
C
E
H
Aluminum
Copper
Rare earth metals
Thorium
Zirconium
Lithium
Manganese
Silver
Silicon
Yttrium
Zinc
K
L
M
Q
S
W
Z
Table 2.3:
Tem~er
designations for magnesium alloys [1].
Divisions
Description
F
As fabricated
0
Annealed, recrystallized (wrought products only)
H1
H2
H3
T1
T2
T3
T4
T5
T6
T7
T8
Tg
T10
Strain hardened only
Strain hardened and then partially annealed
Strain hardened and then stabilized
Cooled and naturally aged
Annealed (cast products only)
Solution heat treated and then cold worked
Solution heat treated
Cooled and artificially aged
Solution heat treated and artificially aged
Solution heat treated and stabilized
Solution heat treated, cold worked and artificially aged
Solution heat treated, artificially aged, and cold worked
Cooled, artificially aged, and co Id worked
,--
7
CHAPTER 2 - LITERATURE REVIEW
2.3 Deformation
2.3.1 Slip
The von Misses criteria reqmre that five independent slip systems are active for
homogeneous plastic deformation [12]. Magnesium has few slip systems at low
temperatures and the deformation of magne sium is, thus, very limited. The main active
slip system is basal slip. Figure 2.2 shows a basal plane on a hexagonal closed packed
lattice. Basal slip occurs along the (0001) plane in the [1 flO] direction.
c
1)
Figure 2.2: Basal plane in a hexagonal close packed lattice.
As temperature is increased, other slip systems become active; prismatic and pyramidal
<c+a> slip. The glide of dislocations on non-basal planes leads to a dramatic increase in
ductility. Figures 2.3 and 2.4 show the different planes for prismatic and pyramidal
<c+a> slip, respectively. Prismatic slip occurs along the (1010) plane in the [1120]
-
-
direction. Pyramidal <c+a> slip occurs along the (1122) plane in the [1123] direction.
8
CHAPTER 2 - LITERATURE REVIEW
c
Figure 2.3: Prismatic plane in a hexagonal close packed lattice.
c
-
[1123]
Figure 2.4: Second order pyramidal plane in a hexagonal close packed lattice.
Figure 2.5 shows the critical resolved shear stress (CRSS) for various slip and twin
systems in a single crystal of magnesium. The CRSS is the stress required to activate a
deformation mode.
9
CHAPTER 2 - LITERA TURE REVIEW
100
r------------.
o
c+a> slip
{1012}
Twinnin~
_10
ca
a.
~
----
( ,)
1
•
•
Prismatic slip
Basal slip
0.1
o
200
400
600
Temperature (OC)
Figure 2.5: CRSS for various deformation systems as a function of test temperature
in pure magnesium [13J.
As is illustrated by Figure 2.5, basal slip has the lowest CRSS at all temperatures and is
therefore the primary deformation mode. A study performed by Kaibyshev and Sitdikov
in pure magne sium revealed slip traces on the basal plane at low and intermediate
temperatures (T = 20-300°C) [14]. At temperatures above 300°C, slip lines due to basal
slip, prismatic and pyramidal <c+a> non-basal slip were identified [14, 15]. These
additional slip systems increase ductility by coming closer to fulfilling the five
independent slip systems required by von Mises [16].
2.3.2 Twinning
Twinning is when a part of the crystal undergoes a re-orientation to bec orne structurally
the mirror image of the remainder as reflected in sorne crystallographic plane, i.e. the
twin plane [7]. Figure 2.6 shows twinning planes and directions in a magnesium lattice.
10
CHAPTER 2 - LITERATURE REVIEW
Figure 2.6: Twinning direction and components [13].
Figure 2.5 indicated that the CRSS for twinning remains constant at aIl temperatures.
Concurrently, twinning plays a significant role in magnesium almost independently of
temperature. The combination of basal, prismatic and pyramidal <c+a> slip does not
provide the necessary five independent slip systems required for homogeneous
deformation and thus, twinning is often necessary to satisfy the von Mises requirement
[12]. Twins nucleate at low strains in grains poorly oriented for basal slip and are
frequently observed in deformed magne sium [17, 18]. While twinning does occur at aIl
temperatures, at low temperatures twinning is more prominent because of the lack of
independent slip systems [19, 20]. As the CRSS for the non-basal systems faIls below
that for twinning at higher temperatures, twinning gives way to slip and very little
twinning is observed [13].
The types oftwins that occur depend on the deformation direction. Figure 2.7 shows that
two types of twins occur depending on whether the deformation is under tension ([lOT 2]
twin) or compression ([1011] twin). Under compression double twinning is more likely
([1012] + [1011] double twin) [13]. Multiple twinning in magne sium at room
-
-
temperature develops in two systems: a primary [1012] and a secondary [1011] one
[17].
11
CHAPTER 2 - LITERATURE REVIEW
{lOI2} twin
{l0 Il} twin
t
l
{lOI 1}+ {l0 12} double twin
l
.;1
t
56°<1210>
Basal planes
Figure 2.7: Tension, compression and double twins [13].
During defonnation, the size and volume fraction of twins will increase [21]. Twins may
also disappear as a result of the twinned material undergoing slip and at higher strains,
the 10ss of twins can thus occur by the above process or by "detwinning" in which twins
in close proximity to one another will grow in size and annihilate each other [4].
2.3.3 Additional Deformation Mechanisms
Defonnation in magne sium often involves: defonnation bands, transition bands, kink
bands and shear bands. A defonnation band is a volume of constant orientation that is
significantly different to the orientation(s) present elsewhere in that grain [22]. The
different types of bands, as described by Humphreys and Hatherly, are shown
schematically in Figure 2.8.
12
CHAPTER 2 - LITERATURE REVIEW
Figure 2.8: Schematic diagram of deformation bands, transition bands and kink
bands. The region labeled "T" is a transition band, region A to C to A is a kink
band [22].
In Figure 2.8, region B has a different orientation to that in grain A. The region labeled
"T" at the edge of the deformation band where the orientation changes from B to A has a
finite width and is termed a transition band. In many cases, deformation bands occur with
approximately parallel sides and involve double orientation change in the form A to C to
A. A deformation band of this type is called a kink band.
Macroscopic inhomogeneities called shear bands usually develop at high strains and have
a morphology that is related explicitly to the deformation geometry. These bands
correspond to narrow regions of intense shear that occur independently of the grain
structure and independently of normal crystallographic considerations. In rolled material,
they occur at ~35° to the rolling plane and parallel to the transverse direction. At still
higher levels of strain, large shear bands develop which cross a rolled sheet from one
surface to the other, and eventually, when a uniform population of these bands exists,
failure may occur [22]. The volume of shear bands in heavily deformed metals is such
that they make a significant contribution to the rolling texture; shear bands are also a
major nucleation site for recrystallized grains. The tendency for shear banding increases
with increasing grain size and decreasing deformation temperature [22].
13
CHAPTER 2 - LITERATURE REVIEW
2.4 Recrystallization
Recrystallization is the formation of new set of dislocation-free grains within a deformed
or recovered structure [22, 7]. Recovery is the restoration of original material properties
by the annihilation and rearrangement of dislocations [22]. The growth of recrystallized
strain free grains can eventually consume the deformed microstructure [22].
Recrystallization is an important process by which grain refinement can be achieved in
many alloys including magnesium.
Recrystallization can be termed either continuous or discontinuous. Discontinuous
recrystallization can be divided into two regimes; nucleation and growth. A
recrystallization nucleus can be defined as a crystallite of low internaI energy growing
into deformed material from which it is separated by a high angle grain boundary.
Nucleation corresponds to the first appearance of new grains and growth is when the new
grains replace the deformed material. The nucleation and growth phenomena are
controlled by thermally activated processes whose driving force is provided by the stored
energy of deformation. In order for recrystallization to occur, a minimum amount of
deformation is required. Continuous recrystallization is the process by which new grains
evolve without any distinct nucleation and growth processes [22]. In general, during
continuous recrystallization, dislocations will remain in the recrystallized grains whereas
discontinuous recrystallization removes dislocations through the sweeping action of high
angle boundaries.
In regards to deformation there are two types of recrystallization: static and dynamic.
Static recrystallization refers to recrystallization after deformation has ended. Dynamic
recrystallization refers to recrystallization that occurs during deformation. Static and
dynamic recrystallization have many features in common, and sorne important
differences [22].
14
CHAPTER 2 - LITERATURE REVIEW
2.4.1 Static Recrystallization
After deformation, a dislocation substructure with associated stored energy is capable of
driving static recrystallization. Static recrystallization is the elimination of dislocations
due to the motion of high angle boundaries. The new grains grow until aU the deformed
grains are consumed. A critical strain is required for static recrystallization to take place,
and this critical strain rises as the temperature rises and the strain rate falls. At higher
strains or strain rates, the rate of recrystallization increases due to the reduction of the
subgrain size and increase in sub-boundary density giving rise to a higher density of
nuc1eation sites [23,24,25].
2.4.2 Dynamic Recrystallization
In metals with low or medium stacking fault energy such as magnesium, dynamic
recrystallization may take place when a critical deformation condition is reached. The
stacking fault energy for aluminum is high (200 ml m-2) therefore aluminum alloys often
require complex thermomechanical processing treatment in order to undergo dynamic
recrystallization (due to dynamic recovery). In contrast, the stacking fault energy for
magne sium is reported to be 78 ml m-2, and thus it undergoes dynamic recrystallization
more readily [19].
A simplified description of the phenomenon of dynamic recrystallization, according to
Humphreys and Hatherly, is as follows. Recrystallized grains form at the pre-existing
grain boundaries, but, as the material continues to deform, the dislocation density of the
new grains increases, thus reducing the driving force for their further growth. The
recrystallized grains eventually cease growing. Dynamic recrystallization originates at
high angle grain boundaries. These may be original grain boundaries, boundaries of
dynamically recrystallized grains or boundaries created during straining (e.g. those
associated with deformation bands) [22].
15
CHAPTER 2 - LITERA TURE REVIEW
A "necklace" type structure is therefore often obtained in magne sium samples deformed
at high ternperatures. A "necklace" type structure refers to srnall recrystallized grains that
surround the original grains. Figure 2.9 shows the evolution of a "necklace" structure.
(a)
(b)
(d)
(c)
(e)
Figure 2.9: The development of a "necklace" microstructure during dynamic
recrystallization [22].
In addition to stacking fault energy, the recrystallization of rnagnesiurn is affected by
ternperature, applied strain and alloying. In general, large reductions favor a lower
recrystallization ternperature [6]. Very pure magne sium can recrystallize at roorn
temperature while sorne highly alloyed magne sium alloys recrystallize at very high
temperatures [6,7].
There is sorne debate as to whether the dynamic recrystallization of rnagnesiurn and its
alloys is a continuous or discontinuous process. In work by Watanabe and coworkers,
bulging is suggested to be the dynamic recrystallization rnechanisrn [19]. The undulations
of the boundaries intensify as the stain increases, and new grains are nucleated on the
boundaries of the original grains. Subsequently, the grain refinernent proceeds over the
whole region as the strain increases [19]. This is a discontinuous dynamic
recrystallization process.
16
CHAPTER 2 - LITERATURE REVIEW
Galiyev and coworkers have described the continuous dynamic recrystallization process.
The formation of new grains is connected with extensive cross-slip of dislocations in the
vicinity of pre-existing high angle boundaries. The formation of a subgrain structure
increases by cross-slip of screw dislocations [26].
Necklacing can lead to the formation of shear zones. According to Ion and coworkers,
there is a strong microstructural similarity between shear zones and dynamically
recrystallized regions [27]. The small recrystallized grains formed at the old boundaries
have been found to be favorably oriented for basal slip, but because of constraints
imposed by adjacent grains, shear along AB cannot easily occur (Figure 2.10a). The
recrystallized bands broaden with increasing strain, and regions such as AB and CD
merge to provide a path of easy slip through the specimen (Figure 2.1 Ob). At the tips of
these bands, the high local stresses promote local strain and recrystallization (Figure
2.1 Oc). Once a clear path for shear is produced through the specimen, deformation will be
concentrated in this region, thus producing a shear zone [27].
Figure 2.10: The process by which dynamic recrystallization "necklacing" forms
shear zones [27].
Despite the regular occurrence of dynamic recrystallization at high temperatures, and
even though the recrystallized grains are fine, the grain refinement at higher temperatures
17
CHAPTER 2 - LITERATURE REVIEW
(400°C and above) may be less effective due to rapid grain growth, which offsets the
positive effects of dynamic recrystallization [28]. Conversely, it can be difficult to
achieve high volume fraction of fine grains at 200°C and higher strain rates since
dynamic recrystallization rate is greatly reduced by the relatively low temperature [28].
As the temperature is lowered, a greater strain
lS
needed in order for dynamic
recrystallization to occur. Unfortunately, at low temperatures the deformability of
magnesium is very limited. This reduces the possibility of dynamic recrystallization at
lower temperatures, because high strains cannot be attained by conventional methods
[17]. Normally, the phenomenon of dynamic recrystallization occurs in the temperature
range of 0.5-0.6 to 0.9-0.98 Tm [17]. There are, however, many instances where low
temperature dynamic recrystallization is observed in magnesium. In work by McQueen,
low temperature dynamic recrystallization has been observed in magnesium at T = O.3Tm
[17].
2.4.2.1 Twin Dynamic Recrystallization
Sitdikov, Kaibyshev and Sakai reported that twins can lead to the nucleation of new
grains. Plastic deformation at room temperature leads to extensive twinning on multiple
systems and the formation of dense dislocation pile-ups within initial grains [29]. The
mutual intersection of primary twins and secondary twinning within primary twins results
in the formation of crystallites surrounded by twin boundaries. This is a process termed
twin dynamic recrystallization.
Twin dynamic recrystallization involves three stages. In the first stage, nucleation occurs
by either intersection of various systems of twins or rearrangement of lattice dislocations
within the twin lamellae. In the second stage, twin boundaries are changed into random
high-angle boundaries due to formation of orientation misfit dislocations. As a result, the
nuclei transform into recrystallized grains. In the third stage, boundary migration begins
-,
[29]. At twin intersections small diamond shaped cells are associated with the first
dynamic recrystallization nuclei [18].
18
CHAPTER 2 - LITERATURE REVIEW
2.4.2.2 Additional Mechanisms of Dynamic Recrystallization
There are varlOUS theories on the mechanisms of high temperature dynamic
recrystallization in magnesium. Kaibyshev and Sitdikov, for example, reported that the
mechanism of grain formation in low alloy magnesium was found to involve dynamic
polygonization [30]. The formation of recrystallized grains along the original grain
boundaries occurred following the graduaI conversion of sub-grain boundaries into highangle boundaries [21]. Dynamic recrystallization can also include the formation of kink
bands. These bands are evolved in sorne grain interiors, and fine grains are developed
along corrugated grain boundaries and at triple junctions [28]. The new boundaries
correspond exactly to those of the kink bands. The misorientation and the number of the
boundaries of kink bands rapidly increase with deformation, resulting in the evolution insitu of new grains with high angle boundaries after high strains.
2.4.2.3 Effect of Grain Size on Dynamic Recrystallization
The strongest effect on the kinetics of dynamic recrystallization is exerted by the original
microstructure. A fine starting microstructure will yield a more uniformly recrystallized
microstructure [31]. Figure 2.11 shows the effect that the initial grain size has on the
volume fraction of recrystallized grains and their size. In this work by Bamett,
compression was performed at 300°C, to a strain of 0.5 at a strain rate of O.Ols- 1 and the
microstructure was analyzed. It can be se en that as the initial grain size is reduced, the
volume fraction of material that undergoes dynamic recrystallization increases. The size
of the recrystallized grains is independent of initial grain size.
19
CHAPTER 2 - LITERA TURE REVIEW
100
,
00
\
0
80
7
0
"
""
6
\
60
Vol. 0/0
• • •
•
DRX 40
•
0' ,
"-
...
•
5
d DRX (f.lm)
4
20
3
o
~~~~~~~~~~~~~2
o
5
10
15
20
25
Figure 2.11: Influence of initial grain size on the dynamically recrystallized grain
size and volume percent dynamic recrystallization. The continuous line refers to the
dynamically recrystallized grain size, while the discontinuous line is a fit of data
according to the volume percent dynamic recrystallization [13].
2.5 Deformation Creep
Superplastic magnesium alloys have been developed to improve the formability of
magne sium alloys [32]. Superplastic deformation is the capacity to undergo extensive,
neck-free, tensile deformation prior to fracture. Superplastic tensile elongations are
usually above 200% and can be as high as 1000% before final failure [33,34]. The
microstructural prerequisites for superplasticity are weIl established. They include fine
grain size and grain boundaries capable of sliding while resisting tensile separation
[35,36]. The need for a fine grain size reflects the independent contributions of grain
boundary sliding and dislocation creep during elevated temperature deformation, as
explained below [35, 37-39].
The two main types of deformation creep at high temperatures are grain boundary sliding
and dislocation creep. These mechanisms operate independently and in an additive
20
CHAPTER 2 - LITERA TURE REVIEW
manner. The faster of these two mechanisms will control the deformation response at a
given temperature and strain rate and depending on the forming conditions one or both of
those deformation creep mechanism will be active [40]. When grain boundary sliding
dominates deformation, superplasticity can occur, and when dislocation creep dominates
deformation, normal ductility is expected [41]
Grain boundary sliding is accommodated within the grain boundaries and in adjacent,
mantle-like regions of the grains (Figure 2.12) [42,43]. Accordingly, the deformation
rate associated with grain boundary sliding increases as the grain size decreases.
Mantle
region
Figure 2.12: Mantle region within regions adjacent to grain boundaries in
superplastic materials [41].
Dislocation creep is presumed to occur independently by dislocation motion within the
core regions of the grains [41]. The slip process involves both glide on slip planes and
climb over physical obstacles [41]. Because dislocation creep occurs within the core
region of grains, the deformation rate for dislocation creep is unaffected by changes in
grain size. This suggests a grain size dependent transition in the deformation creep
mechanism, and superplasticity at higher strain rates in materials of tiner grain size. Such
transitions in rate-controlling deformation creep should occur over a narrow range
corresponding to about an order of magnitude in strain rate [44]. Solute-drag creep occurs
21
CHAPTER 2 - LITERA TURE REVIEW
when the rate-controlling step for dislocation motion arises from the drag of solute atoms
on dislocations gliding on their slip planes [45]. Solute drag is found in substitutional
alloys, such as Mg-Al when solute atoms have a significant volumetrie size difference
with the matrix atoms, causing strain fields that interact with dislocations in motion [45].
In aluminum alloys, large elongations can still be achieved when solute drag dominates
the deformation as elongations up to 250% have been seen [45]. Little research has been
undertaken on solute drag in magnesium alloys.
There have been several studies showing superplasticity in magnesium alloys including
AZ31. Elongations of up to 265% have been achieved in samples tested in tension at
450°C at a strain rate of2xl0-4 sol [46,47]. Elongations ofup to 600% have been se en in
AZ61 with initial grains of 6 to 8Jlm deformed at 290°C with a strain rate of 3.3x1O-4s-1
[41].
During superplastic deformation, grain growth has also been observed. In work by Tan
and Tan, specimens initially possessed a relatively fine grain microstructure
(~12um),
but
eventually the grains coarsened due to grain growth during deformation at elevated
temperature [47]. This grain growth during superplastic deformation causes apparent
strain hardening behavior [48].
2.5.1 Deformation Creep at High Strain Rates and Low Temperature
Superplasticity is normally observed in samples tested at low strain rates ranging from
10-5 to 10-3 sol. These strain rates are too low for the commercial forming of structural
materials, and the commercial viability of superplastic materials is therefore limited.
Research into superplasticity indicates the potential for utilizing superplastic forming
capabilities at much higher strain rates [49].
High-stain-rate superplasticity is defined as superplasticity occurring at strain rates at or
above 1O-2s-1 [41,49]. High-strain-rate superplasticity is of great interest because it is
expected to result in economically viable, near net-shaped forming techniques. Grain size
22
CHAPTER 2 - LITERATURE REVIEW
is an important microstructural parameter contributing to high-strain-rate superplasticity.
In general, high-strain-rate superplasticity is observed at relatively high temperatures
(0.8Tm) and very fine grain sizes.
Magnesium alloys have a higher potential for superplasticity at lower temperatures
compared to aluminum alloys [49]. The grain boundary diffusivity of magnesium is
higher because the pre-exponential factor for grain boundary diffusion is two orders of
magnitude larger than that for aluminum (though the activation energies are close to each
other) [49]. This potential has yet to be fulfilled.
2.6 Sheet Production
The high price of magnesium sheet does not result from the raw material costs (which
compete with aluminum), but from a complex direct chill casting and hot rolling
schedule. Thick slabs must be rolled at high temperatures, and low deformations per pass
[5].
Current rolling schedules consist of a homogenization treatment at elevated
temperatures (300 - 500°C, 0.5 - 1.5h) and continuous hot rolling on a reversing mill
until the desired thickness is achieved (approx. 12 - 18 hot rolling passes and 2-3 heating
cycles during rolling to maintain the rolling temperature) [5,50].
A promising technology has been developed by Commonwealth Scientific and Industrial
Research Organization (CSIRO) where the twin roll casting of magne sium sheet
demonstrates potential for large savings on the cost of production (depending on the
tonnage). Twin roll casting enables magne sium alloy strip to be produced directly from
the melt with a thickness at near-net-shape, eliminating the need for extensive rolling of
the cast slab, thereby reducing capital investment and operational costs [50].
Additionally, the rapid solidification achieved through twin-roll casting can potentially
improve alloy properties by improving the homogeneity of microstructures, refining alloy
grain size, reducing segregation, increasing solid solubility, enhancing precipitate
nucleation within the matrix and generating a distribution of fine precipitates [50].
23
CHAPTER 2 - LITERATURE REVIEW
The final sheet microstructure requires investigation. In most research to date a nonuniform grain size is obtained after rolling, although twin-roll casting and rolling does
appear to minimize this effect [50,51]. Other rolling methods that are currently being
developed for magne sium sheet production include cross-rolling and asymmetric rolling
[52,53].
2.6.1 Rolling Parameters
The rolling temperature is an important parameter. As work by Yarita and coworkers
demonstrated, rolling forces decrease with increasing rolling temperature. The authors
found that the limit in reduction to achieve sheet without cracks was around 40% for
temperatures above 200°C, while less than 20% for temperatures below 155°C [54]. It
was found that shear bands developed at rolling temperatures of 175°C while at 255°C
none were found. It was proposed that the shear bands hinder the deformability.
Additionally, below 175°C, twins were observed, while above 200°C dynamic
recrystallization improved deformability. The benefits of low temperature rolling have
also been outlined by other researchers including Ataka and Shinohara [55].
Rolling reduction per pass is important. Bowles and Horton generated severe plastic
deformation in AZ31 sheet material through accumulated roll bonding. This caused
significant grain refinement by dynamic recrystallization, which leads to increased high
temperature ductility [56]. It was found by Sakai and coworkers, that reductions greater
than 60% can be applied by single pass rolling at high speeds at room temperature, with
only minor cracks developed. The speeds involved in this process are above 1000 mlmin
and the rolled sheets exhibited dynamically recrystallized grains below
5~m
2.13 shows the mapping of defects on an AZ31 alloy rolled at 2000 mlmin.
24
[57]. Figure
CHAPTER 2 - LITERA TURE REVIEW
Minor edge
cracks
Periodic edge
cracks
.--~ T -
60
-~
55
-
-oc: 50
-
•
- -- - - / -
.
..:.
j- -
- - _ -:.--
- _. + -
:;; 45
- -
u
:::J
'C 40
- -
~
-
~:!
-.~
1
c-
L
~
- - '--
- ----: - -
- 1- - -
~
•
•••• -
~-- - 1 - ~- - .._,ie·_.. -
r- - -
-
-
... -
J
-
1-
....
• ,.....
: ••
~- c -- • ---4
-~
- - l ..;. - ---'
- -
1
'
-
-
!
: -
:
35
-
- -
•.•
+ --
~
1
.
. -.~.- :
:
•••
1 -.~ ....
---
~
·
1
-
-
-...;
- ...... -
1
-~
+
-
-
•
•
•
-
-
-
-
;1
!
!
~-
. . . !~
-.'
-
-t 1
+-- - -
•• ,
-
-
-
1
- ----i
i
-
-1
..
- ... ~.... ----'
Scissors 40-"'+---,~:- -1
1.
1.
cracks 30
.~ ': - - - - ---;e - - - - - - •
•
•
••
1
-
••••
1
- _ J- - ~
iNo cracks
•
- -
- -
- -
i
25~----·------~----~·~~--------------------~
o
100
200
300
Rolling Temperature (oC)
400
Figure 2.13: Effects of rolling temperature and reduction on defects of AZ31B
sheets rolled at 2000 rn/min [57].
It was determined that warm rolling with heavy reduction per pass is very effective for
grain refinement, and grain sizes in the order of 2.2~m are attainable [58]. The studies
show that large deformations produce greatly refined grain sizes, and thus these
processmg
routes
may
lead
to
sheet
25
amenable
to
superplastic
forming.
Chapter 3
Experim.ental Procedure
3.1 Materials
Three different AZ31 aUoys were investigated in this work. The chemical compositions
of the three alloys are shown below in Table 3.1.
Table 3.1: Ch emlca
. 1 composi Ion 0 fAZ31 magnesi·um alloys.
Alloy
A
B
C
AI (wt %) Zn (wt %) Mn (wt %)
3.5
0.98
0.47
2.7
0.9
0.26
3.3
1.0
0.25
Composition A was used for the study on the recrystallization behavior of AZ31 during
compressive deformation. Composition B was used for a study of the effect of reheating
on as-cast microstructures and hot compression behavior. Composition C was hot roUed
into sheet, and the hot tensile behavior was evaluated.
3.2 Hot Rolling
The hot roUing was performed using a Stanat pilot-scale reversing mill at Materials
Technology Laboratory-CANMET in Ottawa. Figure 3.1a is a picture of the Stanat mill;
Figure 3.1 b is the sand bath used to heat the samples before rolling and in between
passes.
26
CHAPTER 3 - EXPERIMENTAL PROCEDURE
Figure 3.1: (a) Stanat mill used in rolling experiments and (b) sand bath used for
heat treating the samples before rolling.
The 5 mm thick plates sliced from the as-cast ingots were hot-rolled to 1.65 mm thick
sheets in the Stanat pilot reversing mill (two high configuration, 150 mm diameter x 190
mm wide rolls). To generate different as-hot rolled microstructures, various rolling
schedules were devised by varying the combinations of the processing conditions
indicated in Table 3.2. The temperature was monitored by a thermocouple inserted at
mid-thickness into the si de of each plate.
. AZ31 sheet.
Tabl e 32
. . H ot ro Irmg con thons usedf,or prod ucmg
Hot-rolling
Reheat and
Reheat
Reduction
Roll
Soak Time
per pass
speed
Rolling
(Hr)
(%)
(rpm)
Temperatures
eC)
350
400
450
15
30
1
10
20
50
The reduction on the reversing mill was carried out in either three passes with 30%
reduction per pass, or seven passes with 15% reduction per pass (according to the two
rolling schedules given in Table 3.3). To keep the hot rolling entry temperature of the
strip the same for each schedule, the strip, after each pass, was transferred back to the
sand bed for intermediate heating at the initial homogenizing temperature. After the last
pass, the specimen was air-cooled. The total reduction in thickness was 67%.
27
CHAPTER 3 - EXPERIMENTAL PROCEDURE
Table 3.3: Hot rolling schedules for AZ31 alloy reheated and rolled at 350, 400 and
450°C. The rolling temperature for each pass is the same for a given rolling
schedule.
Rolling Schedule A - Roll Speed 50 rpm
Pass #
Roll Gap (mm)
Total
Reduction per
Strain Rate
(S-I)
pass (%)
Reduction (%)
1
4.24
14.8
14.8
8.2
2
3.61
27.6
15.0
8.9
3.07
38.3
14.8
9.6
3
10.4
4
2.62
47.5
14.9
11.6
5
2.22
55.6
15.5
1.88
62.2
14.9
12.3
6
7
1.60
67.9
14.9
13.3
Rolling Schedule B - Roll Speed 50 rpm
Pass #
Roll Gap (mm)
1
2
3
3.48
2.44
1.70
Total
Reduction (%)
30.1
51.0
65.8
Reduction per
pass (%)
30.1
29.9
30.2
Strain Rate
(s-1
11.9
14.5
16.9
3.3 Tensile Testing
The elevated temperature tension testing, which was used to assess formability, was
performed at General Motors Research and Development center located in Warren,
Michigan. The tensile machine used was a model 5568 screw-driven Instron, with an
Instron 3119-007 fumace and a Merlin data acquisition system (Figure 3.2). Temperature
was monitored with two thermocouples. One thermocouple was suspended in the fumace
chamber and the other was attached to the lower anvil.
28
CHAPTER 3 - EXPERIMENTAL PROCEDURE
Figure 3.2: (a) Tensile machine and (b) a view into the furnace where samples are
tested.
Tensile specimens with gauge section 6.4mm wide and 25.4mm long were machined
from the as-rolled sheets using wire electro discharge machining. The rolling direction is
parallel to the gauge length. A schematic diagram of the sample used in tension is shown
in Figure 3.3.
+--25----+
t
20
65
i
!
~
~6
L
25
0.5 radius -
;
~ four locations
rn
!
Figure 3.3: Tensile sample dimensions used for hot tensile testing.
Prior to starting the tensile tests, the fumace was heated to the test temperature and the
sample was inserted into the cross-head. The furnace was then allowed to stabilize at the
desired temperature and was then held at temperature for 2 minutes. The sample was then
29
CHAPTER 3 - EXPERIMENTAL PROCEDURE
pulled in tension until failure, after which the specimen was removed from the fumace
and immediately quenched in water.
Table 3.4 shows all the corresponding process parameters of the sheet samples that were
rolled. Every "x" indicates sheet that was tested in this study.
.
Ta bl e 34 : P rocess parame ters t es ted·ID th"IS St U dIy.
Process Parameters
Reheating Temp. C
1
Soaking Time, hr
% Reduction per pass
Roll Speed, rpm
400
350
15
1
10
30
15
450
30
15
1
10
30
15
30
10
15
30
2 5 2 5 2 5 2 5 2 5 2 5 2 5 2 5 2 5 2
o 0 o 0
0 o 0 o 0 o 0
0 00 0 0 0
o
o
15
30
5 2 5 2 5
o
0 0
o
x
x
0
Alloy
AZ31 as hot-rolled
x x x x
x
x
x x
x
Hot tensile experiments on the rolled sheets were performed at 300, 400 and 450°C at
strain rates of 0.1, 0.01 and 0.001 S-1. When insufficient material was available the lowest
tempe rature (300°C) and the highest strain rate (0.1s- 1) were eliminated. After tensile
testing the effect of as-hot rolled microstructures, and therefore process parameters, were
compared on the basis of flow stress, elongation and microstructure.
3.4 Compression Testing
The unixial compression tests were performed using an MTS 100 coupled with a radiant
fumace (Figure 3.4). Temperature was measured by a chromel-alumel (K-type)
thermocouple in contact with the compression sample. Stainless steel compression anvils
were attached to the frame of the MTS by water cooled steel supports, and enveloped
with an argon atmosphere. Samples were quenched in water within 1 second of the end of
the test.
30
CHAPTER 3 - EXPERIMENTAL PROCEDURE
Compression Anvils
Cooling Hoses
Radiant Furnace
Figure 3.4: MTS 100 used for hot compression testing.
The MTS computer records data in terms of force and linear displacement as measured
by a load cell and a linear variable differential transformer (LVDT). Temperature and
time during deformation are also recorded. Testar software was utilized to control the
mechanical test. AIl experiments were performed at constant true strain rate.
Compression samples of 1104 mm in height and 7.6 mm in diameter (heightldiameter
ratio of 1.5) were machined from the cast plates of compositions A and B. Thin sheets of
mica coated with a boron nitride powder were placed between the sample and the anvils
to minimize friction during compression.
Isothermal compression was performed from room temperature to 400°C on AZ31 of
composition A. Specimens were heated to the deformation temperature, held for 10 min
to achieve temperature homogeneity, and then deformed at a true strain rate of 0.1 S-1 to a
true strain of 004. After deformation the samples were quenched in water. The samples
31
CHAPTER 3 - EXPERIMENTAL PROCEDURE
were then heat treated for 5 minutes at 450°C, quenched and the microstructures were
studied. This anneal 'simulates' the heating cycle applied after hot rolling but before
elevated temperature forming.
A reheating investigation was undertaken using two plates of composition B (Table 3.1).
The plates were sectioned into 6 pieces (Figure 3.5); sections that contain the marker
lines were discarded as they contained much po rosity (confirmed by x-ray radiography)
and possible chemical composition differences due to casting turbulence.
Figure 3.5: AZ31 alloy of composition B cast in a water-cooled copper mold used for
heat treatment study.
The procedure used in this study is outlined in Figure 3.6. The samples were first
reheated to the test temperature for the appropriate length of time. The heat treatment
conditions are shown in Table 3.5. After reheating the samples were quenched in water.
The specimens were then heated to 350°C and compressed at a true strain rate ofO.1s- 1 to
a true strain of 0.6.
32
CHAPTER 3 - EXPERIMENTAL PROCEDURE
400, 450, 480°C
350°C
Temp
quench
Time
Figure 3.6: Reheating followed by compression schedule.
Table 3.5: Reheating schedule employed.
Temperature("C) Time (min)
400
15
30
60
120
15
30
60
120
15
30
60
120
450
480
33
CHAPTER 3 - EXPERIMENTAL PROCEDURE
3.5 Characterization
3.5.1 Optical Metallography
At McGill, the examination was performed on an Epiphot 200 Nikon Optical Microscope
and the micrographs were taken with a Clemex Image Analyzer. At General Motors,
quantitive analysis was performed using Image Pro software.
3.5.1.1 Sample Preparation for Compression Samples
Compression samples were sectioned along the longitudinal (compression) axis and
mounted in Technovit 4004 cold curing resin. Samples were then ground using 240, 400
and 600 grit silicon carbide papers. Polishing was performed with 9/lm and 3/lm diamond
suspension and finally a 0.05/lm (OPS-0.05) alumina powder suspension. Grains and
twin boundaries were revealed using an acetic-picral solution consisting of 4.2 g of picric
acid, 10 ml of acetic acid, 10 ml of distilled water and 70 ml of ethanol. The dendritic
structure of the as-cast material was revealed using a glycol etch consisting of 1 ml nitric
acid, 75 ml ethylene glycol, and 25 ml distilled water.
3.5.1.2 Sam pie Preparation for Tension Samples
Optical samples were taken from the grip and gauge sections of the tensile bars. In order
to perform optical microscopy, samples used for microstructure were cut using a diamond
blade and cold mounted in Technovit 4004 cold curing resin with the transverse direction
face down along the thickness. The samples were then ground with Buehlers 600, 800
and 1200 grinding papers, polished with 3/lm and l/lm diamond paste then 0.05/lm
(OPS-0.05) alumina powder suspension. The grain boundaries were then revealed using
acetic-picral solution.
34
CHAPTER 3 - EXPERIMENTAL PROCEDURE
After mechanical polishing, the samples were etched in the acetic-picral solution for
approximately 2 seconds. The grain size was measured using the linear intercept method
with a minimum of 80 grain boundaries per image and a minimum of 3 images.
3.5.2 Scanning Electron Microscopy
Scanning electron microscopy was performed using a Zeiss EVO 50 on the surface of
selected samples (Figure 3.7). Both secondary and back scattered detection modes were
used. Surfaces were prepared by cutting the pulled specimen down to a length adequate
for SEM analysis followed by a cleaning step in an ethanol solution to remove any debris
or lubricant present.
Figure 3.7: Scanning electron microscope used in the hot-tensile study.
35
CHAPTER 3 - EXPERIMENTAL PROCEDURE
3.5.3 Electron Backscattered Diffraction
Electron backscattered diffraction (EBSD) was carried out on a Philips XL30 field
emission gun scanning electron microscope (FEGSEM). EBSD data were acquired and
analyzed using HKL Channel 5 EBSD software (HKL Technology, Denmark). The
FEGSEM was operated at 20kV and the specimen tilted at 70°. For orientation mapping
the scan step size was set at 0.4 or 0.5J.lm. High angle grain boundaries >15° and low
angle grain boundaries > 1.5° are represented by thick and thin black lines, respectively
[59].
36
Chapter 4
Results
4.1 Deformation of As-cast AZ31 of Composition A
Preliminary studies on the deformation of as-cast AZ31 were undertaken to help develop
a greater understanding of the deformation behavior of magnesium alloys. As such, a
study on the compression behavior of as-cast AZ31 was performed. AlI samples are of
composition A in Table 3.1.
4.1.1 As-cast Grain Size
Figure 4.1 is an electron backscattered diffraction (EBSD) image of the as-cast
microstructure. The grains are generally equiaxed with an average grain diameter of
297~m.
Figure 4.1: EBSD image of AZ31 of composition A and associated grain size
distribution.
37
CHAPTER 4 - RESULTS
4.1.2 Compression Behavior
After deformation to a strain of 0.4 at different temperatures, the macroscopic appearance
of the cylindrical compression samples is shown in Figure 4.2. At higher temperatures the
deformed sample is more uniform and shearing has not taken place, although, at 150°C
and below fracture due to shearing is visible. Shear bands are visible below 300°C.
Figure 4.2: Compression samples deformed to a true strain of 0.4 at a true strain
rate of 0.1 S-l.
4.1.3 Microstructures after Deformation
The microstructures deformed at low and high temperature before and after annealing are
different. These are shown in Figure 4.3. Figures 4.3a and 4.3b are the as-deformed
microstructures after compression at 350 and 150°C respectively. Figures 4.3c and 4.3d
are microstructures obtained after annealing the deformed specimens at 450°C for 5
minutes. Recrystallized grains are evident when the deformation is performed at 350°C
(Figure 4.3a). A necklaced microstructure is visible with small recrystallized grains
surrounding coarse undeformed grains. At 150°C, no recrystallization is visible as
extensive twinning is observed (Figure 4.3b). Note that the microstructures at 150°C were
taken in the vicinity of the sheared region.
38
CHAPTER 4 - RESULTS
Recrystallized grains are visible in Figure 4.3c, but grains appear slightly larger than in
Figure 4.3a with similar amounts of recrystallized volume. Figure 4.3d shows a
completely recrystallized structure consisting of small grains on the order of 3-20Jlm.
Figure 4.3: Microstructures of compression samples deformed to a true strain of 0.4
at a true strain rate of 0.1s-1 at temperatures of (a) 350°C and (b) 150°C and after an
annealing step at (c, d) 450°C for 5 minutes.
4.2 Effect of Reheat Treatments on Microstructures and
Deformation Behavior of As-cast AZ31 of Composition B
This study was performed to determine if reheat treatments on the as-cast material affect
the microstructure and subsequent deformation behavior. AH samples are of composition
B in Table 3.1.
39
CHAPTER 4 - RESUL TS
4.2.1 Microstructures after Reheat Treatments
Figure 4.4 shows optical microstructures of samples as-received, reheated to 400°C for 30
minutes and 450°C for 60 minutes, respectively. The microstructures are similar; the
grains appear similar in size and shape after both heat treatments and are on the order of
100-200llm.
Figure 4.4: Water-cooled copper plate samples (a) as-received and rebeated to (b)
400°C for 30 minutes and (c) 450°C for 60 minutes.
Figure 4.5 shows the evolution of microstructures as the reheat temperature increases
from 400 to 480°C and reheat time increases from 15 to 120 minutes. Etching has been
performed to reveal dendrites (Figure 4.5). With increasing time and temperature the
dendrites disappear. After 30 minutes at 480°C the dendrites appear to have completely
disappeared (Figure 4.5f). For the case of 15 and 30 minutes of reheating, there appears
to be a large effect of temperature on the disappearance of dendrites (Figures 4.5a, b, c
and 4.5d, e, f). However, in the case of 1 hour of reheating, the effect of temperature
diminishes (Figure 4.5g, h, i). As shown in Figure 4.5j, k, l, after two hours of reheating
"...-~,
at 400, 450 and 480°C the dendrites appear to be completely dissolved.
40
CHAPTER 4 - RESULTS
Figure 4.5: Water-cooled copper plate samples reheated to 400°C for 15,30,60 and
120 minutes (a, d, g, j), 450°C for 15,30,60 and 120 minutes (b, e, h, k) and 480°C
for 15, 30, 60 and 120 minutes (c, f, i, 1) 480°C for 120 minutes.
41
CHAPTER 4 - RESULTS
4.2.2 Deformation Response to Reheating Cycles
Figures 4.6 and 4.7 are true stress vs. true strain plots for speCImens tested in
compression at 350°C to a true strain of 0.6 at a true strain rate of 0.1s- 1 after being
reheated to 400°C for 15 and 60 minutes, respectively. Each curve represents a specimen
that has been subjected to the same reheat condition, and therefore illustrates the
variability of the structure after reheating. After 15 minutes of reheating (Figure 4.6),
there is a large spread in the flow curves. After 60 minutes of reheating, the flow stress
values are very similar (Figure 4.7). At 120 minutes, the flow stress curves are very
comparable to those obtained after 60 minutes.
Figure 4.8 is a plot of spread (i.e. the difference between the highest and lowest values) in
maximum compressive strength (MCS) vs. time for aH three testing temperatures. The
spread in MeS decreases with increasing reheating time. After 60 minutes, the spread in
MCS is minimized for aH three reheating temperatures, with 450°C giving the least
spread.
-120 I-------.::::;::::;::;;;;:;;;;;;;;;;;;;;:::::=----------------~
-100
Ci
c..
::::il:
-; -60 f--I-W--- - - - - - -
UI
I!!
f i)
-40
-20
o
--~
... _ - - - - - - _ . _ - - - - - - -
~---------------~-----------~----~
o
-0.1
-0.2
-0.3
-0.4
-0.5
-0.6
Strain
Figure 4.6: Flow stress after reheating to 400°C for 15 minutes and deforming at
350°C at a true strain rate of O.1s- I to a true strain of 0.6.
42
CHAPTER 4 - RESULTS
-120
- - - - - -...- - . - - - - - . - - - - - - - -....- - - . - - - - . - . - . - - - - - - - - - .
-40
------.
-20
o
~----------------------------~-----------~
o
-0.1
-0.3
-0.2
-0.4
-0.5
-0.6
Strain
Figure 4.7: Flow stress after reheating to 400°C for 60 minutes and deforming at
350°C at a true strain rate of 0.1s-1 to a true strain of 0.6.
30,--------------------------------------------~
25
li
c...
~
en
20
::E
.5
15
-------_.
-----\-------
------_.-------------------- - - - -
.- - - - - _ .
1
----
,-+-400
___ 450
'-,------1
-'-480:
o
1
"'C
III
~
~
-'1..'----_._-------_._------
10
.--------------
!
1
~====:::::::::::=========i--l
5 ---------------- .
1
!
o~---------------------------------------------~
o
20
40
60
80
100
120
140
Time (min)
Figure 4.8: Spread in MCS vs. reheating time for AZ31 reheated at 400, 450 and
480°C.
43
CHAPTER 4 - RESULTS
4.3 Tensile 8ehavior of Rolled AZ31 Sheet of Composition C
4.3.1 Tensile Curves
Typical stress-strain curves of sheet samples pulled in tension are shown in Figure 4.9.
The alloy exhibits a dependence of flow behavior on temperature and strain rate and there
appears to be three distinctive types of behavior. At the lower deformation temperatures
and highest strain rates there is considerable work hardening initially, although the UTS
is reached after a low strain. At the higher deformation temperatures and lowest strain
rates there is graduaI work hardening and a very long pre-UTS region. Under all other
conditions, there is little work hardening and a short pre-UTS region, followed by a
relatively long post-UTS region.
60
i
~---
i
-.-------
-4500.1
1
!
1
Ci
--j
40
1
·--l
1
-4500.01
-4500.001
1
---j
1
D.
t
!.
-4000.1
UI
...=30
-
-4000.01
en
CI)
;
1
[
1
.
1
U
-4000.001
...
:::l
1- 20
1
1
-----1
10
1
!
o~------------------------------~--~------------~----~
o
0.2
0.4
0.6
0.8
1.2
1.4
True Strain
Figure 4.9: Stress-strain curves after processing at 400 and 450°C with strain rates
of 0.001,0.01 and 0.1s-1• In the legend, the first number is the testing temperature,
the second number is the strain rate.
44
CHAPTER 4 - RESULTS
Figure 4.10 shows the yield stress against strain rate for deformation temperatures of
450°C, 400°C and 300°C for aU hot-rolling processing conditions. The yield stress was
determined by using a 5% offset method. This was chosen because, in high tempe rature
tests, the standard 0.2% offset typically gives values in the pseudo-elastic region, not in
the plastic region, of the curve. With decreasing temperature and increasing strain rate,
the yield stress increases. There is a variation in the yield stress of each specimen tested
at a specifie strain rate and temperature. This may be due to the different as-hot rolled
microstructures. Hence, the specimens that exhibited the lowest and highest yield stress
values at 450 and 400°C at strain rates of 0.001 and 0.1s- 1 were taken for microstructural
examination.
120 ,---------------------------------------------------~
100
~
1
60
CIJ
~
r--
1
80
::!:
l
----~~-
~----~~1
40
~----~~---------'=---
20
~~~~~::
I~------~~~~~-1
450°C
o 1 ____________________________________________________
~I
0.0001
0.001
0.01
~
0.1
Strain Rate (S·1)
Figure 4.10: Yield stress vs. strain rate for samples tested at 450, 400 and 300°C.
The flow curves of the specimens exhibiting maximum and minimum yield stress tested
at 450°C and 400°C and at 0.ls- 1 and 0.00Is- 1, are shown in Figures 4.11 and 4.12. The
samples with lower yield stress displayed better ductility than the higher yield stress
curves for the same strain rate. At 400°C and 0.00Is- 1, the difference in elongation is not
-~
as large as at 450°C.
45
CHAPTER 4 - RESULTS
40,---------------------------------------------------------,
35
30
c;-
25
Il.
~
1/)
-
e 20
- - - - - + - - - - - - - - - - - - - t - - - - - - - - - - - - - . - - -...
1/)
en
G)
2
1-
~
15 +1---
_ _ _\Ols_l
10 -+11----"..-='------
1
5-+ __- - - - - -
1
O~------~----------------------~----------------~~--~
o
0.2
0.6
0.4
0.8
1.2
1.4
True Strain
Figure 4.11: Stress vs. strain for specimens with the highest and lowest yield stresses
tested at 450°C at strain rates of 0.1 and O.OOls-l.
60,--------------------------------------------------------~
50+--/----
c;-
40
t-ir'~- ....-
- - - .
_._--
...
----
- - - - - _ .. _ . - - - _ .
---
- - _.. _ - - - - - - - - j
Il.
~
1/)
~
...
30 +-....-
.. - - - - -..-.-.. --~\
_ .._ - - - - -
en
G)
2
1- 20
I,·~··--:::·:::=:~~~~:::::--
. .---O.OOls-1
- -..._ - - - . _ . - - - - _....
10
- - -
-----\---+ -------
O~----------------------------------------------~--------~
o
0.2
0.4
0.6
0.8
1.2
True Strain
Figure 4.12: Stress vs. strain for specimens with the highest and lowest yield stresses
tested at 400°C at strain rates of 0.1 and O.OOls-l.
46
CHAPTER 4 - RESULTS
Table 4.1 summarizes the yield stresses and elongations obtained in the specimens that
were taken for micro structural analysis. At a given strain rate and temperature the
samples with the lowest yield stresses had the longe st elongations. At 400°C and 0.1s- l ,
however, the specimen with the higher yield stress had the longer elongation.
Table 4.1: Summary ofyield stresses and elongations at different testing
' rates used d
' tensIOn
. t estmg.
.
temperatures an d
urm2
stram
Temperature
450
450
450
450
400
400
400
400
re)
Strain rate (s- ) Yield Stress (MPa) % Elongation
8.7
248
0.001
13.1
116
0.001
31.5
0.1
92
35.3
32
0.1
13.9
176
0.001
20.5
164
0.001
44.2
0.1
60
50.6
104
0.1
4.3.2 Microstructural Analysis
4.3.2.1 Low Strain Rate (0.00Is- 1)
The microstructures at key stages of deformation of specimens tested at 450°C are
presented in Figure 4.13. The microstructures on the left hand side of the page (Figures
4.13a, c, e, g) have the lowest yield stress, while those on the right hand side (Figures
4.13b, d, f, h) have the highest yield stress. The as-hot rolled microstructure with the
greatest amount of visible deformation, shear bands, and unrecrystallized grains is the
microstructure that led to the lowest tensile yield stress (Figure 4. 13a). On the other hand,
the microstructure with relatively uniform and recrystallized grains had the highest yield
stress (Figure 4. 13b).
Figures 4.13c and 4.13d are microstructures that are present before pulling the tensile
specimens (i.e. after exposure to the test temperature for several minutes). The
microstructures have average grain sizes of 9.8Jlm and 16.2Jlm respectively. There is,
however, bimodality in the grain sizes. Bimodality exists when there is a large difference
in grain sizes present throughout a specimen. In the case of the 9.8Jlm grain size the range
47
CHAPTER 4 - RESUL TS
is from 6.9 to 10.4llm while for that with a grain size of 16.21lm the range is from 13.2 to
21. 1Ilm.
Figures 4.13e and 4.13f are microstructures at the failure tip. In Figure 4.13f, there is
sorne necking visible at the tip. In contrast, Figure 4.13e shows relatively uniform
elongation failure without necking. There is more cavitation at the failure tip in Figure
4.13e than 4.13f. Figures 4.13g and 4.13h are higher magnification images of near the
failure tip of the samples. Again, more cavitation is se en in Figure 4.13g than Figure
4.13h. The final grain sizes after deformation are 17.8 and 17.3llm respective1y.
48
CHAPTER 4 - RESULTS
Figure 4.13: Microstructures oflowest (Ieft side) and highest (right side) yield stress samples tested at
450°C at a strain rate ofO.00ls· 1• Microstructures include the as-rolled condition (a, b), pre-pulling
condition (c, d), and the failure tip (e, f, g, and h). The processing condition for the left side is 350°C
+ 1 hr reheat, 30% red., 50 rpm. The processing condition for the right side is 450°C +1 hr reheat,
15% red., 20 rpm.
49
CHAPTER 4 - RESULTS
The microstructures at key stages of deformation of specimens tested at 400°C are
presented in Figure 4.14. The observations are similar to those made for 450°C, except
for the grain sizes.
The microstructures just before deformation (after a few minutes exposure to 400°C)
have average grain sizes of 9.21lm (the range is from 8.2 to 15.4llm) and 16.2llm
respectively (the range is from 10.8 to 18.9Ilm). The final grain sizes after deformation
are 10.8 and 111lm respectively.
50
CHAPTER 4 - RESULTS
Figure 4.14: Microstructures of lowest (Jeft side) and highest (right side) yield stress samples tested at
400°C at a strain rate ofO.OOls· J• Microstructures include the as-rolled condition (a, b), pre-pulling
condition (c, dl, and the failure tip (e, f, g, and hl. The processing condition for the left side is 350°C
+10 hr reheat, 30% red., 50 rpm. The processing condition for the right side is 450°C +1 hr reheat,
15% red., 20 rpm.
51
CHAPTER 4 - RESULTS
4.3.2.2 High Strain Rate (O.ls- 1)
Using the same approach as above, the microstructures at key stages of deformation of
specimens tested at 450°C and 400°C are presented in Figures 4.15 and 4.16,
respectively. The microstructures on the left hand side (Figure 4.15a, c, e, g and 4.16a, c,
e, g) have the lowest yield stresses, while those on the right hand side (Figure 4.15b, d, f,
h, and 4.16b, d, f, h) have the highest yield stresses. In contrast to the lower strain rate
results, the as-hot rolled microstructure with relatively uniform and recrystallized grains
had the lowest yield stress (Figure 4.15a, 4.16a). On the other hand, the microstructures
with the greatest amount of visible deformation, shear bands, and unrecrystallized grains
are the microstructures that led to the highesttensile yield stresses (Figures 4.15b, 4.16b).
Figures 4.15c, 4.16c, 4.l5d, and 4.16d are microstructures that are present immediately
prior to tensile deformation. Figures 4.15e, 4.16e, 4.15f, 4.16f are microstructures after
deformation both close to the fracture point and away from the tip. For the specimen
---
exhibiting the lower flow stress at 450°C, at the fracture point (Figure 4.15e) the grains
are approximately the same size as the original one (Figure 4.15c). However, moving
away from the tip, there is much grain coarsening (Figure 4.15g). In contrast, for the
higher flow stress specimen, the grains at the fracture tip (Figure 4.15f) are coarser than
the original one (Figure 4.15d); moving away from the tip, the grains are refined (Figure
4.15h).
At 400°C, the specimen with the lower yield stress exhibits little change in the grain size
after reaching the test temperature and after deformation and fracture. The higher yield
stress specimen also exhibits little change after test temperature is reached, but the
bimodality after fracture appears to be much more pronounced (Figure 4.16h).
At both 450°C and 400°C the same hot rolling schedule produced the lowest flow
stresses. This was the condition with 450°C; 10 hr reheat time, 30% reduction per pass
and a roll speed of20 rpm.
52
CHAPTER 4 - RESULTS
Figure 4.15: Microstructures of lowest (Ieft side) and highest (right side) yield stress samples tested at
450·C at a strain rate ofO.1s- I • Microstructures include the as-rolled condition (a, b), pre-pulling
condition (c, dl, the failure tip (e, f,) and a good distance away from the failure tip in the deformation
zone (g, hl. The processing condition for the left side is 450·C +10 hr reheat, 30% red., 20 rpm. The
processing condition for the right side is 350·C +10 hr reheat, 15% red., 20 rpm.
53
CHAPTER 4 - RESUL TS
Figure 4.16: Microstructures of lowest (left side) and highest (right side) yield stress samples tested at
400°C at a strain rate ofO.1s- I • Microstructures include the as-rolled condition (a, b), pre-pulling
condition (c, d), the failure tip (e, f,) and a good distance away from the failure tip in the deformation
zone (g, h). The processing condition for the left side is 450°C +10 hr reheat, 30% red., 20 rpm. The
processing condition for the right side is 350°C +1 hr reheat, 30% red., 20 rpm.
54
CHAPTER 4 - RESULTS
4.3.3 Effect ofProcessing Conditions on Yield Stress
In many cases, sheets were tested with aU processing conditions identical except for one
process condition. Inspection of Table 3.4 reveals that there are sorne cases where aU
processing conditions except one were kept identical.
Figure 4.17 is a plot of the yield stress vs. strain rate for sheet samples roUed at 350°C, 10
hour reheat, 50 rpm roU speed and reductions per pass of 15 and 30%. Figure 4.18 is a
plot of the yield stress vs. strain rate for sheet samples that have been roUed at 450°C, 10
hour reheat time, 20 rpm roU speed and reductions per pass of 15 and 30%. Both figures
indicate that the samples roUed with 30% deformation per pass have the lowest yield
stress at aIl temperatures and strain rates involved. In Figure 4.18, as the testing
temperature is increased the difference in yield stress between the samples rolled with 15
and 30% reduction diminishes.
55
CHAPTER 4 - RESUL TS
40 .
30
/~-+--------c-/~-r--~~~-----------------i
- - - .--------7''--
10
o
..
_-
---~_._--
~---------------------------------------------~
0.0001
0.001
0.01
0.1
Strain Rate (5.1)
Figure 4.17: Yield stress vs. strain rate at 400 and 450°C for samples rolled at 350°C,
10 hour reheat time, 50 rpm roll speed with reductions per pass of 15 (black) and
30% (red).
80
70
60
~
--_.~._~~~
--~-----
-.----.--:/'---------1
50
~
fil
~
;;;
40
"C
1
>
30
20
----.-.--=~
10
o
~.------------------------------------------------~
0.0001
0.001
0.01
0.1
Strain Rate (5.1)
Figure 4.18: Yield stress vs. strain rate at 300, 400 and 450°C for samples rolled at
450°C, 10 hour reheat time, 20 rpm roll speed with reductions per pass of 15 (black)
and 30% (red).
56
CHAPTER 4 - RESUL TS
Figure 4.19 is a plot of the yield stress vs. strain rate for sheet samples rolled at 350,400
and 450°C, 1 hour reheat, 20 rpm roll speed and a 30% reduction per pass. There is a
difference in the yield stress behavior. It can be seen that for both samples tested at 400
and 450°C the specimens rolled at 350°C (black) had different yield stress behaviors
depending on the strain rate. At the fastest strain rate, the samples rolled at 350°C had the
highest yield stresses while at the slowest strain rate they had the lowest yield stresses.
The samples rolled at 450°C (red) also had very different yield stresses at the different
strain rates (especially at 400°C testing temperature). At a strain rate ofO.1s· J the sample
rolled at 450°C had the lowest yield stress while it also has the highest yield stress at a
strain rate ofO.001s· J •
120,----------------·------------------------------------,
100
"ii'
---_._---
-----------------~-------
80
0..
~
.....
II)
II)
CIl
60
1/)
"t:I
Ci
>=
40
20
0
0.0001
0.001
0.01
0.1
Strain Rate (S·l)
Figure 4.19: Yield stress vs. strain rate at 300, 400 and 450°C for samples rolled with
a 1 hour reheat time, 30% reduction per pass, a roll speed of 20 rpm at
temperatures of 350 (black), 400 (brown) and 450°C (red).
Figure 4.20 is a plot of the yield stress vs. strain rate for sheet samples rolled at 350°C, 1
hour reheat, roll speeds of20 and 50 rpm with a 30% reduction per pass. Figure 4.21 is a
plot of the yield stress vs. strain rate for sheet samples rolled at 350°C, 10 hour reheat,
roll speeds of20 and 50 rpm with a 15% reduction per pass. In Figure 4.20 a roll speed of
50 rpm had the lowest yield stress while in Figure 4.21, a roll speed of 20 rpm had the
57
CHAPTER 4 - RESUL TS
lowest yield stress. Figure 4.22 is a plot ofthe yield stress vs. strain rate for sheet samples
rolled at 450°C, 1 and 10 hour reheats, 20 rpm roll speed and 15% reduction per pass.
There is no difference in the yield stresses with different reheat times.
100------·----·---------------
1
BO
L--------.---,
~
::
~
60
~-----i
40 [-
>
1
20
~----1
o ~I____________________________________________~
0.0001
0.001
0.01
Strain Rate
0.1
1
(5- )
Figure 4.20: Yield stress vs. strain rate at 300, 400 and 450°C for samples rolled at
350°C, 1 hour reheat time, 30% reduction per pass, with a roll speed of 20 (red) and
50 rpm (black).
58
CHAPTER 4 - RESULTS
40 i
l
~- ..~-~
30
'ii
II..
~
i
VI
..
i ---
~ 20
II)
"C
Gi
>=
1
10~:--~~~-~~--~~~
~~----~-----
1
!
o
'~----------------------------------~----------~
0.0001
0.001
0.01
0.1
Strain Rate (S-l)
Figure 4.21: Yield stress vs. strain rate at 400 and 450°C for samples rolled at 350°C,
10 ho ur reheat time, 15% reduction per pas s, with a roll speed of 20 (red) and 50
rpm (black).
100 , - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - - ,
l
!
80
!~----- . .--------.. ---------
1
lso ;-1--j
i
!
40"""1~--
>=
1
201-----
- - -
-~~-
-~-
._-----
!
o
~!------------------------------------------------~
0.0001
0.001
0.01
0.1
Strain Rate (S-l)
Figure 4.22: Yield stress vs. strain rate at 300,400 and 450°C for samples rolled at
450°C, 15% reduction per pas s, 20 rpm with a reheat time of 1 (black) and 10 hours
(red).
59
Chapter 5
Discussion
5.1 Deformation of As-cast AZ31 of Composition A
Figure 4.2 showed that magnesium sheared at low temperatures (150°C and below), while
at high temperatures, the ductility is increased. This transformation is directly related to
the activity of slip systems in magnesium's HCP structure. Due to CRSS considerations,
the main deformation mechanisms in magnesium at low temperatures are basal slip and
twinning [13,17,18]. At high temperatures, CRSS for prismatic and <c+a> slip decreased
and they contributed to deformation [14, 16].
Dynamic and static recrystallization are prevalent in magnesium alloys. In Fig. 4.3a
(deformed at 350°C), a necklace type structure was seen where recrystallized grains were
distributed along
pre-existing
boundaries.
These
recrystallized grains
evolved
dynamically because quenching of the specimens after deformation prevented static
recrystallization from occurring. Dynamic recrystallization was attributed to the low
stacking fault energy (~78 ml m-2) and limited slip systems in magne sium and its alloys
[19].
Figure 4.3b showed a twinned structure when deformed at 150°C. Basal slip cannot be
seen by optical microscopy; therefore the only sign of deformation was twinning [21]. In
Figure 4.3d, after annealing, the microstructure was completely recrystallized. Static
recrystallization has occurred during the annealing step. The stored energy in the form of
dislocations was the driving force for static recrystallization. Aside from dislocations,
60
CHAPTER 5 - DISCUSSION
most of the literature has shown that twins can be the location of dynamic
recrystallization [17, 18, 21, 30]. However, the influence of twins on static
recrystallization is not well researched.
In this study, there was no obvious strong
influence of twins on the annealed microstructure.
Note that the microstructures in Figures 4.3a (deformed at 350°C) and 4.3c (annealed for
5 min at 450°C) are similar, although grain growth may have occurred in the latter. This
indicates that dynamic recrystallization has relieved much of the deformation introduced
during compression. This may suggest that as soon as dynamic recrystallization takes
place, further plastic strain was concentrated in the dynamically recrystallized regions,
and there is little deformation in the grain interiors. This will promote shearing as well as
generating very inhomogeneous structures. Therefore, it may be that dynamic
recrystallization is not beneficial for attaining a homogenous as-hot rolled structure.
However, if the as-cast grain size was initially small, dynamic recrystallization may be
effective for grain refinement by hot deformation. The level of as-cast grain size required
for full dynamic recrystallization would appear to be of the order of 10 microns, based on
the dynamically recrystallized grain size in Figure 4.3a. Clearly, low temperature
deformation followed by a static anneal would appear to be much more feasible in terms
of grain refinement strategies.
5.2 Effect of Heat Treatments on Microstructures and
Deformation 8ehavior of As-cast AZ31
The as-cast grain size did not increase significantly with higher reheat temperature or
longer reheat times (Figure 4.4). The Burke and Turnbull analysis of grain growth
determined that the rate of grain growth can be represented by the following equation:
(1)
61
CHAPTER 5 - DISCUSSION
where R is the mean grain size at time t, Rois the initial mean grain size and
C2
is a
constant. As the initial grain size is increased, more time is needed for grain coarsening to
occur [22]. The as-cast grain size of the water-cooled copper plate material was over
lOOllm and therefore grain growth was relatively slow. Sorne specimens contained
anomalous large grains but this may been due to grain size inhomogeneity in the cast
plate, and not due to actual grain growth.
The dendritic structure of the cast plate was affected by reheat time and temperature
(Figure 4.5). At high temperatures and long reheat times, dendrites dissolved. With
increasing reheat temperature, the dendrites were further dissolved. This is expected as
diffusion is related exponentially to temperature [60]. At higher temperatures, the
increased thermal energy increases the probability of atomic jumps. Additionnally, the
Mg-Al eutectic is metastable and exposure to high temperatures resulted in a complete
dissolution of the Mg 17A1 12 intermetallic phase [61].
Homogenization of the as-cast structure improved the repeatability of deformation
curves. As the reheat time was increased (Figures 4.6 and 4.7) repeatability of
deformation curves improved. Following a 60 minute reheat at 400, 450 and 480°C the
repeatability increased. There was an exponential relation between the spread in the data
(a measure ofrepeatability) and the reheat time (Figure 4.8). This was not expected as the
diffusion equation is exponentially related to temperature and not time. However, the
temperature range (400 to 480°C) was quite small compared to the time range (15 to 120
min), which may account for this observation.
After 1 hour of reheat at all test tempe ratures (400, 450 and 480°C) the as-cast structure
was sufficiently homogenized to allow for consistent results. Although the flow stress of
400°C after an hour is repeatable, the dendrites did not appear to be [ully dissolved. This
suggested that repeatability can be achieved before complete dissolution of dendrites.
Additionally, there was potential incipient melting at 480°C. The best reheating schedule
is, thus, 450°C for 60 minutes, and, in fact, is seen to have minimized the spread of MCS
62
CHAPTER 5 - DISCUSSION
(Figure 4.8). However, 400°C for 60 min was adequate with regard to repeatability of hot
deformation processing.
5.3 Deformation Behavior of Rolled AZ31 Sheet
5.3.1 Tensile Curves
The temperature and strain rate dependence of the flow stress of magnesium alloys was
shown in Figure 4.9. As the temperature was decreased and the strain rate was increased,
work hardening occurred. This occurs in many magnesium alloys and has been observed
by many researchers [12,13]. The relationship between true stress and true strain is given
by:
(2)
where cr is the true stress, K is the material constant, ë is the strain rate, and m is the
strain-rate sensitivity factor. Because the m factor was been found to be >0.3 for fine
grained magnesium, the material exhibited a significant degree of strain-rate sensitivity
[12,41,62].
Figure 4.10 showed all the yield stresses obtained at the 3 different test temperatures and
strain rates. Again, the general dependence on temperature and strain rate is clear as the
yield stress increased with decreasing temperature and increasing strain rate. For a given
test temperature and strain rate, there was a large variation in yield stress. The difference
was on the order of 10 MPa. This variation directly corresponded to the different
microstructures produced by the various processing conditions used in hot rolling
(discussed in chapter 5.3.2). Thus the processing conditions had an effect on the high
temperature tensile testing flow curves.
Due to insufficient samples, less data was available on the samples tested at 300°C
(Figure 4.10). All further analysis in this study was performed on specimens tested at 400
63
CHAPTER 5 - DISCUSSION
and 450°C. These temperatures were chosen because they are within the temperature
range for present forming techniques such as superplastic forming.
The highest and lowest yield stress samples produced at strain rates of 0.001 and 0.1 S·I
were sectioned for micro structural analysis. The flow curves (true stress vs. true strain) of
those samples were plotted against each other in Figures 4.11 and 4.12. A low flow stress
is usually an indication of good ductility. Flow stress is a function of the material's grain
size, test temperature and strain rate. If the test conditions are the same, then the flow
stress can be directly related to the grain size. A fine grain material, because of more
grain boundary area per unit volume, and thereby more grain boundary diffusion, will
show a lower flow stress and higher tensile elongation under superplastic conditions,
than a coarse grain material. At 450°C, the samples with the lowest yield stresses had the
highest elongations (Figure 4.11). At O.OOls· l , the 248% elongation suggests that
superplastic deformation has occurred [33,34]. The 116% elongation achieved by the
sample with the highest yield stress suggests that normal ductility is occurring in this
sample [41]. At the higher strain rate (0.1 S·I), the elongations are quite low for both
samples. It should be noted that defects can affect the ductility, and perhaps the influence
of defects was more marked at higher strain rates.
A few of the samples received in this study contained defects. Through energy dispersive
spectrometry and back-scattered imaging, it was determined that these defects were due
to rolled-in oxide and rolled-in scale. This caused them to fail prematurely. In Figure
4.12, both samples with the lowest yield stresses at the 400°C testing temperature had
lower elongations then the samples with the highest yield stresses. Defects may have
been present in these two samples. Unfortunately, because ofthe sampling methodology,
it was not possible to examine the fracture surfaces for defects.
To obtain statistically good elongation values, 3-4 samples for each condition would be
needed. Since these numbers of specimens were not available, accordingly, the
comparison of hot-rolling processing conditions was done on the basis of yield stress.
64
CHAPTER 5 - DISCUSSION
5.3.2 Microstructural Analysis
The bulk of the study on the hot-rolling processmg conditions was based on the
micro structural analysis. The analysis will be divided into two sections, low strain rate
(0.00Is- 1) and high strain rate (0.ls- 1).
5.3.2.1 Low Strain Rate (O.001s- 1)
As described in the results section 4.3.2.1, there was a difference in the grain sizes before
tension between samples with the highest and lowest yield stresses (Figures 4.13, 4.14).
This difference in grain size directly translated into different deformation creep modes
[41]. Creep is a deformation process that occurs in solids at high temperatures, typically
when T/Tm> 0.5 where Tm is the melting point of the solid [41]. Three mechanisms can
occur at the atomic level: slip by dislocation movement, sliding of adjacent grains along
grain boundaries, and diffusional flow. Normally, the mechanisms are considered to
occur independently of each other.
For both 450 and 400°C the same deformation creep mechanisms were present. The
deformation creep mode of the lowest yield stress samples (left side of the page of
Figures 4.13 and 4.14) had the characteristics of grain boundary sliding: uniform
elongation, cavitation and large elongations. Uniform elongation refers to the lack of
necking as was seen in Figures 4.13e and 4.14e. The lack of flow localization indicated
that failure was primarily controlled by cavitation [63]. Cavities that are distributed along
grain boundaries or around the interfaces of the magne sium matrix have been observed
by many investigators [64-67]. It was reported that cavities are thought to be nucleated by
the continuous condensation of vacancies on grain boundaries that experience a normal
tensile stress or by vacancy clustering due to stress concentration on grain boundary
inclusions produced by strain incompatibility and grain boundary sliding [66]. On the
left hand side of the figures the final failure appeared to be due to the interlinkage of
internaI voids (Figures 4.13 and 4.14). This brittle fracture was reported to be as a result
of the nucleation, growth and interlinkage of internaI voids [68]. The elongations
obtained (248 and 176% ) are in the superplastic range [33,34].
65
CHAPTER 5 - DISCUSSION
Grain boundary sliding was mostly likely the deformation creep mode due to the
extremely fine grains present in the microstructure prior to deformation (Figures 4.13c,
4.l4c). It has been seen that a grain size less than 1O~m is necessary for grain boundary
sliding in magne sium [35-41]. Fine grain sizes are needed because they increase the area
of the mantle-like region adjacent to the grain boundaries. In this process (the core and
mantle theory proposed by Gifkins), grain boundary sliding is accommodated by slip
within the mantle region [41]. The grain size of the microstructures prior to deformation
was 9.8 and 9.2~m at 450 and 400°C respectively. These extremely fine grains were
produced from the microstructures that had the greatest amount of visible deformation
after being hot-rolled (Figures 4.l3a, 4.14a). These deformed microstructures
recrystallized after exposure to high temperature for several minutes. The possibility of
static recrystallization after deforming at low temperatures was also noted in the results
obtained after compression testing in chapter 4.1.
Another proof of grain boundary sliding was the presence of submicron fibers. These are
shown below in Figure 5.1. Submicron fibers are thought to be evidence of local grain
boundary plasticity. These fibers bridge surface cavities along the direction of tensile
deformation. Fiber formation is distinctly associated with grain boundary sliding creep
and may be attributed to local grain boundary plasticity [63].
Figure 5.1: Submicron fibers visible parallel to the tension axis in sam pie with very
fine grain size.
66
CHAPTER 5 - DISCUSSION
rr-'
On the right hand side of the page of Figures 4.13 and 4.14, there appeared to be more
than one deformation creep mechanism active. This was due to its larger initial grain size,
the presence of necking and low levels of cavitation. It has been seen that both grain
boundary sliding and dislocation creep occur at high temperatures, and that the
mechanisms will occur independently and in an additive manner [40]. Thus, due to the
larger initial grain sizes (16.211m at both 450 and 400°C) dislocation creep and grain
boundary sliding creep were the active deformation creep mechanisms. This grain size
dependence on the deformation mechanisms has been observed by many researchers
[35,41].
The grain sizes after deformation in the highest and lowest yield stress specimens were
very similar. At 450°C, the final grain size in the lowest and highest yield stress samples
was 17.8 and 17.311m, respectively. At 400°C, the final grain sizes were 10.8 and llllm,
respectively. These grain sizes suggest that there is a steady-state grain size at a particular
temperature. While the initial grain size is dictating the deformation mechanism, after a
certain amount of deformationlexposure to high temperatures, the grain sizes bec orne
similar in size. At 400°C the final grain sizes are smaller due to the lower thermal energy
caused by the 50°C temperature difference. In the case of grain boundary sliding, in this
work, grain growth occurs from the start of deformation to the end, which is probably due
to grain coarsening.
By contrast coarsening or refinement was seen in the case of
dislocation creep, and this could be due to dynamic recrystallization since the
dynamically recrystallized grain size is generally thought to be only a function of the
Zener-Hollomon parameter.
5.3.2.2 High Strain Rate (O.ls- 1)
At the high strain rate (O.ls- I ), the favorable microstructure was, in fact, the exact
opposite of the lower strain rate case. This was due to the fact that at a strain rate ofO.ls- 1
dislocation creep was the main active deformation creep mode. Dislocation creep is
controlled by dislocation slip within grains. In the majority of alloys such as magnesium,
the creep mechanism is controlled by the glide step in the glide/climb mechanism
67
CHAPTER 5 - DISCUSSION
~--.
because solute atoms impede dislocation motion [41, 69]. These types of alloys,
including magnesium alloys, are termed Class 1 solid solutions. The glide-control
mechanism is independent of grain size, therefore, the relatively larger grains present
prior to testing (Figures 4.l5c, 4.16c) are of no significance.
Another factor that indicated that the main deformation creep mechanism is dislocation
creep is that, at this strain rate, little cavitation was seen in any of the specimens. It is of
value to view the microstructures in the vicinity of the fracture tip. Figures 4.15e, 4.15f,
4.16e and 4.16f revealed the microstructures very close to the fracture surface. This is
within the heavily necked region and the grains were somewhat coarser than the asreheated grains (Figures 4.l5c and 4.15d). Moving away from the fracture tip, there was a
much coarser grain size in the lower yield stress specimen (Figure 4.15g), but little
change in the higher yield stress specimen (Figure 4.15h); although, there could be grain
refinement.
These changes in grain size may be due to dynamic recrystallization throughout the
specimen. In the neck area the grains should be small because the strain rate is greatest
near the fracture point. Similarly, the region away from the neck should have a larger
dynamically recrystallized grain size (lower strain) [70]. However, the only specimen that
follows this hypothesis is the lower yield stress specimen tested at 450°C (Figure 4.15).
5.3.3 Processing Parameters
The micro structural analysis has revealed that two deformation mechanisms were present
in the magne sium sheet depending on the strain rate and grain size. At slow strain rate
and small grain size, the active deformation mechanism was grain boundary sliding. At
the same strain rate with larger grain size, there was a component of dislocation creep. At
the fast strain rate, the deformation mechanism regardless of grain size was dislocation
creep.
If the processing conditions associated with the lowest flow stress and highest elongation
are compared to those with the highest flow stress and lowest elongation sorne guidelines
68
CHAPTER 5 - DISCUSSION
for hot-rolling schedule design can be made. Of the strain rates used in this work, the
most important one, with respect to CUITent forming practice, is the lowest strain rate
(O.OOls- I ). This is in the regime of superplastic forming. On the other hand, higher strain
rates improve productivity, and will therefore be of importance in the future. Thus, the
effect of microstructure and rolling schedule on the lowest and highest strain rate tensile
behavior will be examined below.
5.3.3.1 Low Strain Rate (O.OOls- 1)
Table 5.1 shows the list of the processing conditions that are associated with low and
high yield stress. These processing conditions are the ones that were used for the
micro structural analysis.
..
'Id'mg th esu
t d'le d micros t ruc t ures.
T a bl e 51 P rocessmg con dt'
1 Ions yle
TemJ>erature
450°C
400°C
Processing Condition
Lowest Yield Stresses
350°C, 1 hr reheat, 30% red., 50 rpm
350°C, 10 hr reheat, 30% red., 50 rpm
Highest Yield Stresses
450°C,1 hr reheat, 15% red., 20 rpm
450°C, 1 hr reheat, 15% red., 20 rpm
From Figure 4.10, other processing conditions that had low yield stresses at 400 and
450°C are:
>>-
400°C, 10 hr reheat, 30% reduction., 50 rpm
400°C, 1 hr reheat, 30% reduction., 20 rpm
Additionally, another processing condition with high yield stress from Figure 4.10 was:
>-
450°C, 10 hr reheat, 15% reduction., 20 rpm
From these conditions, the rolling parameters that are most important (in decreasing
order) for obtaining a good microstructure for elevated temperature tensile properties are:
deformation per pass, rolling temperature, roll speed and reheat time. The parameters will
be considered separately in order of importance.
The deformation per pass appears to be the single greatest factor in the processing of
AZ31 alloy. In all cases where a low flow stress was observed, a larger reduction per pass
69
CHAPTER 5 - DISCUSSION
of 30% was used. The large reduction per pass created finer grains upon recrystallization
during reheating before testing (Figures 4.13c and 4.l4c). This is despite the fact that the
as-rolled material is very inhomogeneous to begin with. This is an indication that static
recrystallization is occurring in the samples with the very deformed microstructure during
the preheating. The. static recrystallization is creating very fine grains that lead to
improved high temperature formability. Large deformations per pass have been seen to
be beneficial in recent research. Bowles and Horton showed dynamic recrystallization
occurred after a reduction per pass of 60-80% reduction at 400°C rolling temperature.
The fine grains obtained through dynamic recrystallization lead to improvements in
ductility and strength [56]. Sakai and coworkers also found that reductions per pass
(>60%) created fine recrystallized grains along shear bands [57]. Figures 4.17 and 4.18
reinforce the fact that the amount of deformation per pass is of critical importance. At aIl
strain rates and temperatures, the samples rolled with a greater reduction per pass had the
lowest yield stress.
The rolling temperature is second in importance. The lower rolling temperature produced
a partially recrystallized, as-hot-rolled microstructure that recrystallized to very fine
grains during the preheating before tension testing. At 350°C, the low deformation
tempe rature induced the shear bands and large deformation that then recrystallized to
very fine grains suitable for grain boundary sliding (as seen in the microstructure analysis
section). By contrast, rolling at higher temperature produced a more equiaxed,
homogeneous, recrystallized microstructure, but upon reheating in preparation for tensile
testing the grains increased in size and were substantially larger compared to the samples
rolled at lower temperatures. Thus, rolling at even lower temperatures should be
examined to see if any further grain refinement just prior to elevated temperature tensile
testing is possible. Several authors have rolled at low temperatures in order to generate a
very fine microstructure. Yarita and coworkers found that warm rolling at temperatures
around 200°C can be beneficial for the deformability of magnesium sheet [54].
Preliminary research by Ataka has found that warm rolling around lOO-200°C has
..... / --
adequate rollability [55]. Figure 4.19 showed that the samples rolled at 350°C at the
slowest strain rate displayed the lowest yield stress.
70
CHAPTER 5 - DISCUSSION
The roll speed and the reheat time have less effect on the as-rolled microstructure
(accordingly, on the tensile behavior). Nevertheless, three of the four low yield stress
samples were obtained at a roll speed of 50 rpm. Both samples with higher yield stress
had a roll speed of 20 rpm. The greater roll speed increases the strain rate imposed on the
sample and perhaps this is making dynamic recrystallization during hot rolling more
difficult, increasing the level of work hardening in the as-hot rolled structure. It has been
found by several authors that high strain rate rolling normally coincides with large
deformations per pass [55,56,57]. Figures 4.20 and 4.21 do conflict as in one case the
sample with the slowest roll speed had the lowest yield stress, while the opposite was true
in the other. There does not appear to be any difference in flow behavior with longer
reheat soak times. Therefore, with regard to ductility, it is clear that a reheat time of 10
hours is not necessary. In Figure 4.22 there does not appear to be a large difference (if
any) on the yield stress with different reheat times. This conclusion is consistent with
research performed in the reheat study that demonstrated that one hour at elevated
temperatures was sufficient to homogenize the as-cast structure.
5.3.3.2 High Strain Rate (O.ls- 1)
At this higher strain rate (0.1s- 1), the favorable microstructure was, in fact, the exact
opposite of the lower strain rate case. The main deformation mode was dislocation creep.
Thus, fine grains were not advantageous because grain boundaries are barriers to
dislocation motion.
For both 450°C and 400°C the same processing condition produced the lowest flow
stress; 450°C, a 10 hr reheat, 30% reduction per pass and a roll speed of 20 rpm. If the
desired as-hot rolled microstructure is coarse and fully annealed then most of this rolling
schedule is appropriate. The only major surprise is the high reduction per pass, which
would be expected to either increase the level of work hardened structure or decrease the
dynamically recrystallized grain size.
71
CHAPTER 5 - DISCUSSION
More research is needed to study the deformation mechanisms at this fast strain rate. The
deformation mechanism is most likely dislocation creep, but whether or not solute drag is
the mechanism has yet to be seen.
72
Chapter 6
Conclusions
1. The as-cast dendritic structure was broken down by exposure to elevated
temperatures. This is a diffusion based process and a reheat of 450°C for 60
minutes gives the best combination of dissolved dendrites and repeatable flow
stresses.
2. During compression on as-cast AZ31, dynamic recrystallization occurred at high
temperatures (above 300°C) and resulted in a "necklace" type structure consisting
of recrystallized grains surrounding larger unrecrystallized grains. The extent of
dynamic recrystallization was low, possibly because strain is concentrated in the
necklaced regions. Therefore, it may be that dynamic recrystallization was not
beneficial for attaining a homogenous as-hot rolled structure.
3. In as-cast material, low temperature deformation followed by a static anneal
generated a much more homogenously recrystallized microstructure compared to
dynamic recrystallization, in compression testing.
4. Magnesium AZ31 sheets have a strong dependence on the temperature of tensile
deformation and imposed strain rate. At low temperatures and high strain rates
there was considerable work hardening initially as the UTS is reached after a low
strain. At higher tempe ratures and low strain rates, there was graduaI work
hardening and a very long pre-UTS region.
73
CHAPTER 6 - CONCLUSIONS
5. The micro structural analysis on sheet after tensile testing has revealed that two
deformation mechanisms were present in the magne sium sheet depending on the
strain rate and grain size. At slow strain rate and small grain size, the active
deformation mechanism was grain boundary sliding. At the same strain rate with
larger grain size, there was a component of dislocation creep. At the fast strain
rate, the deformation mechanism regardless of grain size was dislocation creep.
6. The ideal processing condition for rolling depends on the strain rate and
temperature to be used in the forming process, since these will dictate the
deformation mechanisms.
•
At the relatively slow strain rates of O.OOls- l , the best as-hot rolled
microstructure consisted of a heavily deformed structure that when
reheating during testing produced very fine grains that then deformed by
grain boundary sliding creep at a strain rate ofO.OOls- l .
•
At higher strain rates (O.1s- l ) the best microstructure was fully
recrystallized, leading to a relatively coarse structure after heating to the
tensile test temperature
7. In terms of rolling schedule, lower temperatures, and larger strains per pass were
favorable in generating the heavily worked as-hot rolled structures, which were
beneficial to low strain rate elevated temperature tensile properties.
8. Rolling parameters that did not strongly affect the elevated tempe rature tensile
properties were the reheat time and the roll speed.
74
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