1. 1 Copyright © 20xx by ASME
April 15, 2016, Kingston, Ontario, Canada
THE EFFECT OF ZN ON COLD ROLLING TEXTURES IN CAST ALUMINUM
Brad Merkley
Queen’s University
Kingston, Ontario, Canada
ABSTRACT
Aluminum alloys with a target Zn-wt% of 0, 1, 10,
30 and 60 have been cast and prepared for rolling.
Each cast will have samples rolled to 75% and 90%
elongation. The samples have been scanned with an x-
ray diffractometer in order to obtain a measurement of
the bulk alloy texture. An orientation distribution
function has been calculated using the net results of the
measurements and pole figure plots have been
generated for qualitative analysis. The ODF is
represented in terms of the volume fractions of 5
characteristic components; Cube, Goss, Brass, S and
Copper. The reduction in the Copper component
suggests that there may be a decrease in the stacking
fault energy when zinc is added to aluminum.
However, the relatively unaffected Cube and Goss
components infer that a change in the stacking fault
energy does not influence texture development as
much as in Cu-Zn alloys [1]. The fluctuation in the
Brass and S component may be due to characteristic
aging phenomenon that can occur at the 10 Zn-wt%
and 30 Zn-wt% alloys, but it is predominantly
assumed that large reductions during the initial rolling
of the 30 Zn-wt% lead to this behavior.
OBJECTIVE AND ANALYSIS METHODS
In order to represent the bulk texture in a
comparative fashion, the orientation distribution
function (ODF) needs to be calculated from x-ray
diffractometer measurements. The ODF represents the
average volume fraction of a specific crystal
orientation in a sample from a common reference
frame. It is common to segment the data from the ODF
into the volume fraction of specific crystal orientation
components that attribute to characteristic metallic
behaviors [2]. If desired, a range of orientations
between these components can be represented in a
fiber plot. This method can be preferred as movement
of characteristic distribution fibers for a specific
crystalline structure can be tracked in Euler ODF plots
and it can give more insight into how certain texture
components form into others [1]. This paper compares
the key texture components measured in Al-Zn alloys
to the values determined in a characteristic fcc fiber for
Cu-Zn alloys. The analyzed orientation components
have been listed in table 1 below.
Table 1: List of key fcc texture components. S is depicted in
a multitude of orientations, but the (123)<63-4> is the
dominant one [3]. In this report, S will refer to the
summation of the following orientations: S1(1,2,-3)
<6,3,4>, S2(1,2,3)<6,3,4>, S3(1,-2,-3)<6,-3,4>,
S4(1,2,3)<6,-3,4>
Component Orientation
Cube (100)[010]
Goss (011)[100]
Brass (011)[211]
S* (123)<63-4>
Copper (112)[111]
Figure 1: Pole figures from the 111 and 001 orientation.
Taken from 90% reduced Cu. The upwards triangle
represents Copper, o as various S components, the square as
Brass and the diamond as Goss. The inverted triangle
represents the Dillamore texture which is not analyzed in
this experiment [3].
Texture development in fcc metals can be
compared to two extreme cases: “pure metal” and
“alloy” type texture distributions [3]. These two
characteristic texture patterns are usually depicted by
the volume fraction of each peak component they
contain. The pure metal pattern has a high amount of
the Copper component while the alloy type contains a
peak at the Brass component. Figure 2 depicts an
example 111 stereographic projection of these extreme
texture patterns.

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Figure 2: Characteristic "pure metal" and "alloy" 111 pole
figures at 95% elongation: pure Cu in a) and 70:30 brass in
b) [1] [4]. The extremes can be corrosponded to the
components in Figure 1
A transition between the two extreme texture
patterns in Figure 2 may infer a change in the stacking
fault energy with the addition of more alloying
elements. The Brass orientation is known to exhibit
copious twinning compared to the other orientations,
which is a common phenomenon to occur in low
stacking fault energy metals [3] [5]. The key objective
of this experiment is to indirectly observe a change in
the stacking fault energy in Al-Zn alloys through a
change in the rolling texture development.
It is important to note that these characteristic
textures can change on a variety of other phenomenon
unrelated to the stacking fault energy. Experiments
where plutonium-gallium alloys were rolled at varying
temperatures showed that a decrease in temperature
transitioned the rolling textures from the typical pure
metal distribution to the alloy distribution [4]. This
acts as a reminder that the change in stacking fault
energy can only be inferred as there are multiple
mechanisms that can influence the texture
development.
The results of the experiment will be
compared to the results in the Hirsch & Lücke paper,
analysing the effects of Zn in Cu rolling textures [1].
The data obtained from the experiment will be
compared to the α and β fibers produced by Hirsch &
Lücke. The α fiber represents a distribution of
orientations along the 𝜑1 axis (as the orientation of
textures are represented by Euler angles) which
intersects the Brass and Goss components. The β fiber
goes through a multitude of axes in a non-linear
fashion, dubbing itself the name of the “skeleton” line
[1]. This β fiber shows the transition from the Copper
component, through the major S components and
ending at the Brass.
In contrast to Copper, Aluminum has a much
higher stacking fault energy (𝛾𝑆𝐹𝐸
𝐴𝑙
≅ 170 J/m2
and
𝛾𝑆𝐹𝐸
𝐶𝑢
≅ 80 J/m2
) [4]. It is predicted that the higher
stacking fault energy aluminum has may postpone the
transformation from the pure metal type texture to the
alloy type texture with the addition of zinc content.
EXPERIMENTAL METHODS AND
OBSERVATIONS
Aluminum ingots with a Zn-wt% of 0, 1, 10,
30 and 60 were the target compositions to cast for the
experiment. A Ti-wt% of 0.15 was targeted for each
alloy to act as a grain refiner [6]. The grain refiner
consisted of 5 Ti-wt% and 1 B-wt% as TiB2 in an
aluminum matrix. 99.99% Al was used for the samples
containing a Zn-wt% of 0 and 1 while 99.7% Al was
used for the samples containing a Zn-wt% of 10, 30
and 60. The Al, Zn and grain refiner was cut with a
bandsaw to the respective ratios required to fill 150%
of the casting volume (126mm x 62.8mm x 9.56mm)
to mitigate the effects of equiaxed grain growth and
shrinkage during casting. Table 2 below shows the
respective amounts of each metal used in each casting.
They were measured using a triple beam balance.
Table 2: Mass of each metal used to create the castings
The aluminum of each alloy was placed in a
sand crucible and placed in a gas furnace to melt. Each
sample was to be poured into the casting mold when
above the liquidus region by 70K. The melts were
measured with a thermocouple and poured into a steel
mold covered in graphite spray. The pouring
temperatures determined from the phase diagram in
Figure 3 are displayed in Table 3.
Table 3: Pouring temperatures used in the casting process
Figure 3: Phase Diagram of Al-Zn alloys
The aluminum in the first casting (pure Al)
was heated up to 745o
C, then the grain refiner was

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added in one solid chunk. It seemed that the solid
chunk of grain refiner sank to the bottom and was not
melting well. After some mechanical stirring with a
steel rod covered in a graphite spray, the melt was
picked out of the furnace at 750o
C then poured into the
mold. This was above the 730o
C melting temperature
to account for possible heat lost while bringing the
melt to the mold. The sample was air cooled, then
removed from the cast. Because the chunk of grain
refiner had issues melting, the remaining grain refiner
for the four other samplers were cut into small pellets
with shears to increase the surface area to promote
faster melting
The aluminum in the 1 Zn-wt% alloy was
melted up to 990o
C, then the zinc and grain refiner was
added into the melt. A porous dross started to form at
the top of the melt with the addition of the grain
refiner, assumed to be caused by dirt and impurities on
the surface. The melt was removed at 750o
C again and
poured into the mold.
The aluminum in the 10 Zn-wt% was heated
up to 900K. The zinc was added to the melt at 850o
C
and the grain refiner at 800o
C. However, the grain
refiner had issues melting again and the pellets had
sunk to the bottom. After some mechanical stirring
with the graphite coated rod, the melt was removed at
740o
C and poured into the mold. The pour was done
discontinuously, stopping halfway through due to
misalignment. At the bottom of the crucible, un-
melted pellets of the grain refiner were observed. Due
to dissatisfaction of the grain refiner melt, an Al-Ti
phase diagram was observed and it was assumed that
temperatures above 1100o
C would be required to
adequately melt the grain refiner [7]. The gas furnace
used could only reach temperatures up to 1000o
C
safely, so temperatures above 900o
C were used when
adding the grain refiner in for the sequential alloys.
The aluminum in the 30 Zn-wt% alloy was
heated to 950o
C. The grain refiner was added at this
temperature, followed by the zinc at 940o
C. The dross
similar to the second casting appeared after inserting
the grain refiner. The mold was then removed from the
furnace at 710o
C and poured into the mold.
The aluminum in the final casting, the 60 Zn-
wt%, was initially heated to 1000o
C. The grain refiner
and zinc were added at this temperature. Once at
660o
C, the crucible was removed for pouring.
Misalignment with the mold caused the casting pour to
be interrupted halfway through the process.
When the castings were cooled, the top
overflow region of the castings were cut off with a
bandsaw. 30 thou was milled off of each side of the
casting to ensure a flush surface during rolling.
Each of the samples were annealed in nitric
molten salt at an average temperature of 428o
C for 96
hours to ensure adequate homogenization. The
samples were each wrapped in aluminum foil before
insertion.
Table 4: Calculated annealing diffusion distances for
420o
C. The actual salt bath averaged at 428 o
C
Once the samples were removed, they were
placed in a bucket of warm water until cooled to room
temperature.
A Struers Accutom-5 precision cutter was
planned to be used to cut the alloys into strips that
would be used for the rolling tests and the chemical
analysis. However, before it could be used, all of the
alloys had to be cut 8 centimeters from the bottom of
the casting with a bandsaw in order to fit into the
precision cutter. 1cm of the left side of the casting was
also cut off to be sent for chemical analysis. The cuts
that were made on the alloys can be found in Figure 4
and the results of the chemical analysis are shown in
Table 5.
Figure 4: Cutting profile used to fit the castings into the
Struers Accutom-5 precision cutter. The left side (relative to
bottom depicted in the figure) of the casting is what was sent
for chemical analysis
Table 5: Casting compositions from chemical analysis.
Notice the Ti values were lower than desired
The bottom half of figure 4 was then cut into
strips; dimensioned to fit into the channel die of
12.7mm. The Struers Atccutom-5 was used with
Metlab 6NF20 blades (0.020” thick) with a feed rate
of 0.1mm/s, on a low force setting and a RPM of 2000.
The filleted section at the bottom (seen in Figure 4)
was cut off first through visual inspection to ensure

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that all of the samples used would be completely
rectangular to fit in the channel die. Cuts spaced
12.7mm from the fillet cut were then made to produce
the samples.
Figure 5: Picture example of the samples used in the
following experiments. Between sample 5 and T1 is where
the 8cm cut was made to fit into the precision cutter. The
missing left side is what was sent for chemical analysis and
the top budge is the casting overflow region
The 3 bottom strips in the alloy were cut in
half to ensure that the material would not shoot out of
the channel die during compression.
The samples were etched using “aluminum
bright etch”; 50% NaOH, 45% distilled water and 5%
NaNO2. They were rinsed using a 10% nitic acid and
ethanol mix to clean the surface, rinsed with distilled
water and dried off with a hair drier. They were stored
in a freezer at -28o
C to mitigate recrystallization.
Before any of the samples were compressed
or rolled, they were solutionized in the nitric salt bath
again for 10 minutes at 424o
C and quenched in ice
water (temperatures of 4-0o
C) to reintroduce the alloy
as a full solid solution. They were each tested
approximately 5 minutes after the process.
The left half of the 3rd
strip (can be depicted
in Figure 5) in the pure aluminum sample was prepared
and put into the steel channel die. It was wrapped in a
Teflon tape to promote a frictionless surface during
deformation at and reapplied after every compression.
The die was then loaded into the 8502 Instron test
frame. A die feed rate of 0.1mm/s was used for each
compression. Each test was set to hit a target of 50%,
75% and 90% compression with an additional 1mm of
compression added to account for elasticity in the
sample and the compliance of the Instron frame. The
results from the compression tests can be seen in Table
5.
Table 6: Results from the S1-3L sample compression tests in
the Instron frame
The back end of the samples after
compressions 2, 4 and 6 were cut with a handsaw for
future texture analysis. After each test, the top of the
die was used with the Instron machine to push the
sample out of the channel.
The second sample used was the right side of
the 3rd
strip in the 3rd
sample (10 wt%-Zn), seen in
Figure 5. The tests were successfully completed up to
75% compression until the die started to budge and
yield. Due to the failure of the channel die, the rest of
the tests were conducted using a Stanat 6” rolling mill.
The fourth strip in the casting (seen in Figure
5) was used for each alloy in the rolling tests. The
samples were always inserted into the die in the same
direction, alternating which surface was facing
upwards after every reduction. After the targeted 50%,
75% and 90% reductions were approximately reached,
the back end of the sample was sheared off using tin
snips for future analysis. The 1 wt%-Zn sample was
reduced first, with the results displayed in Table 7.
Table 7: Results from the second sample rolling
Once finished, the same testing method was
applied to the 30 wt%-Zn sample. However, the strain
applied during the first few reductions was too great
for the sample to handle and significant cracking
started to occur. In order to mitigate this, more
reductions applying less strain to the sample were used
(a maximum of 3.175mm of compression set per pass
was not to be exceeded). This can be seen after the
third compression in Table 8.

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Table 8: Results from the fourth sample rolling
The same small reduction technique was used
for the 60 wt%-Zn sample next to ensure that no
significant cracking would occur in the material.
Throughout the passes, a large amount of small edge
cracks were forming at the precision cut surfaces, but
nothing significant to affect the structural integrity in
the center face of the sample. The results of the rolling
reductions can be found in Table 9.
Table 9: Results from the fifth sample rolling
The same small rolling reductions were done
for the 10 wt%-Zn sample. Due to the reduced amount
of alloying content, no edge cracking had occurred in
the sample.
Table 10: Results from the third sample rolling
Only the 75% elongation and 90% elongation
samples for each of the alloys were measured for
textures. The resultant strains that were produced in
the rolling processes have been summarized in Table
11. Each of the samples had to be cut into a 12mm by
18mm size using tin snips to fit into the x-ray
goniometer.
Table 11: Strain values produced through compression and
rolling processes
The 75% and 90% elongation samples from
each alloy were sanded down to a target of half of the
thickness with an additional 100 micrometers on the
surface. This was done by sanding each of the samples
sequentially with 220 grit, 400 grit and 600 grit
sandpaper, lubricated with water. Once down to the
desired thickness, the samples were polished using a
lapping wheel with 6 micrometer diamond. The
samples were etched with the aluminum bright etch.
Then, all the sampled were electro polished using 70%
ethanol, 20% water, 8% phosphoric acid and 2%
butylcellosolve at 25V and -20o
C for 60 seconds.
Once electro polished, each of the samples
were set into the x-ray diffractometer and measured at
3 specific crystal planes ( (111), (220), (100), (311) ).
Only the (111) pole figures are referred to in this paper.
The last three alloys (10 Zn-wt%, 30 Zn-wt% and 60
Zn-wt%) are all measured using a chromium filter on
the diffractometer. After each alloy was scanned, pole
figure plots were generated. The data from the
diffractometer was then used to determine the ODF to

6. 6 Copyright © 20xx by ASME
show the peak components of each key texture
orientation.
RESULTS AND DISCUSSION
The resulting pole figures of each alloy has
been depicted below in Figures 6-10. The left of each
figure displays the 90% strained sample texture while
the right displays the 75% strained texture. The plots
are taken from the normal perspective of the 111 plane.
Figure 6: Pole figures of the pure aluminum sample
Figure 7: Pole figures of the 1 Zn-wt% sample
Figure 8: Pole figures of the 10 Zn-wt% sample
Figure 9: Pole figures of the 30 Zn-wt% sample
Figure 10: Pole figures of the 60 Zn-wt% sample
Figures 11 and 12 depict the specific
crystalline orientations measured in the ODF
calculations. The volume fraction of each key crystal
orientation is plotted as a function of the Zn wt%.
Figure 11: Key texture element plot for 75% strain

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Figure 12: Key texture element plot for 90% strain
From observing the pole figures in Figures 6-
10, a general transition from the pure metal type
texture to the alloy type can be seen faintly. The major
part of this texture transformation can be spotted in the
Copper regions (±30% up and down on the vertical
axis, as well as ±60o
and ±50% from the vertical axis)
and in the Brass regions (at the far horizontal edges
and ±30o
from the vertical axis at the outer perimeter).
The Copper region in the pure aluminum sample is
extremely pronounced and mimics the pure metal type
texture distribution seen in Figure 2. However, as the
additional zinc content is added to the alloys, there is
still a pronounced density of the Copper type
orientation, but its magnitude is distributed along the
tertiary textures surrounding its peak. In comparison to
the Hirsch and Lücke paper, this Copper texture
component transformation is seen with the addition of
zinc, particularly between pure copper and 95:5 brass,
which is seen in Figure 13 [1].
Figure 13: Depiction of the Copper texture component
change in pure copper with the addition of 5 Zn-wt% at 95%
elongation. Notice the spreading of the Copper texture
component along the tertiary textures [1]
Figure 10 starts shows the disappearance of
the Copper texture component at 60o
from the vertical
axis at 90% reduction. This behavior seems to
correspond to the 95% strained 95:5 brass sample in
Figure 13. The decrease in the Copper texture
components in both the 90% and 75% compressed
samples (seen in Figures 11 and 12) corresponds to the
β fiber values measured in the Cu-Zn alloys at the
same levels of compression (seen in Figures 14 and 15
[1]).
Figure 14: Beta fiber plots for Cu-Zn alloys. Note that the
Copper, S and Brass orientations can be depicted from left
to right with the orientations depicted on the top axis of the
graph. The key texture components at approximately 75%
compression and 90% compression have been labeled with
the blue and red balls respectively [1]
Figure 15: Additional beta fiber plot for 70:30 brass [1]

8. 8 Copyright © 20xx by ASME
It is worth noting that the aluminum required
much more zinc alloying content to begin to see a
significant reduction in the Copper texture component.
This strengthens the hypothesis that the higher
stacking fault energy that aluminum has relative to
copper delays the transformation from the pure metal
texture distribution to the alloy texture distribution
with the addition of zinc. This is also shown through
the Goss and Cube components in Figures 11 and 12.
The Goss and Cube texture components are also
characteristic of the pure metal type texture
distributions [4]. They see little change in volume
fraction with the addition of zinc. This contradicts the
data found in the Hirsch and Lücke paper as the Goss
component was severely decreased with the addition
of zinc [1]. Because these components are relatively
unaffected by the change of zinc content, it can be
inferred that the extremely high stacking fault energy
of aluminum is not lowered enough with high amounts
of zinc to see a complete transformation to the alloy
type texture distribution.
However, the Brass texture component
undergoes some atypical behavior, particularly
because of the 30 Zn-wt% alloy. As seen in Figures 11
and 12, the Brass and S textures seem to peak at the 30
Zn-wt%. Comparing to the β fiber plots in the Hirsch
and Lücke paper (Figures 14 and 15), the Brass texture
component decreases with the addition of zinc in
elongation ranges from 75-90% [2].
This unpredicted texture behavior may have
been caused by the excess strain that the alloy
experienced during the first few reductions in the
rolling mill, seen in Table 8. Alloys that undergo high
reductions in rolling mills tend to develop stronger
texture peaks at lower overall strains [3]. Especially
with the high zinc content and the macroscopic
cracking observed, the 30 Zn-wt% alloy must have
experienced a large amount of internal stress during
these large reductions.
If the 75% compression pole figures in
Figures 8, 9 and 10 are observed, it can be seen that
the Brass type texture is much more defined at its
respective lower reduction. The excessive initial strain
may have created the conditions necessary for the
Brass and S components to develop further than the
other respective alloys. Also, if the large reductions
have developed the texture distribution at a lower
strain, the raise in the Brass and S components
correspond to the higher reduction samples with a low
amount of zinc, seen in Figure 14. This further
strengthens the hypothesis that the change in stacking
fault energy the aluminum experiences with the
addition of zinc does not affect the rolling texture
development as much as the Cu-Zn alloys.
It is worth noting that the 1 Zn-wt% sample
was also rolled at high strains. However, due to the
lower amount of zinc particles that could pin
dislocation movement during the deformation, The 1
Zn-wt% alloy would experience a much less internal
stress during rolling.
It was theorized that the change in the texture
components could have been influenced by the
synthesis of GP zones in the 1 Zn-wt% and 10 Zn-wt%
alloys and possible spinodal decomposition in the 30
Zn-wt% alloys. This is because the texture
measurements that were taken on the samples were
stored at room temperatures for over a week and
heated up during the sanding process. This may have
contributed to aging decomposition in the respective
alloys.
The misfit strain that the GP zones induce
could promote more crystal rotation. This may have
led to a more dispersed orientation distribution, which
would ultimately explain the lower volume fractions
in Figures 11 and 12. However, since the early stage
GP zones contain a coherent interface, the influence
that the misfit stains would have on the overall alloys
would be fairly insignificant and the GP zones
themselves would be oriented in the deformed texture.
Possible spinodal decomposition may have
led to small recrystallized phases in the material that
would be biased in the typical rolling texture
orientations. This could possibly be implied to
contribute to the increased volume fraction of the
Brass and S components in Figures 11 and 12. Yet the
ambient temperatures the samples were kept at were
probably too low to influence any significant
microstructural changes.
CONCLUSIONS
Primarily, the evolution of the Copper texture
component in the Al-Zn alloy textures may signify a
slight reduction in the stacking fault energy. The
unpredictably high Brass and S component volume
fractions in the 30 Zn-wt% alloys may have been
caused by a large amount shear during the first few
reductions. However, the driving mechanisms for
these results can be isolated in further
experimentation.
In order to confirm the transition from pure
metal type textures to the brass textures, measurements
need to be made with alloy deformations greater than
90%. This is because in the Hirsch and Lücke
experiments, the Brass component decreases with
additional alloy content at 90% compression. Yet the
characteristic β fiber decomposition at lower stacking
fault energies, i.e a major peak in the Brass component
with a decrease in the S and Copper components, is
observed at deformations ≤95% [2]. Future
experiments should consist of deforming the alloys up
to 98% compression and using the β fiber as the main
set of data to infer to the texture evolution. This will
allow for more definitive observations of the pure

9. 9 Copyright © 20xx by ASME
metal to alloy type textures distributions with the
addition of zinc. The shape of the fibers can be
compared to the low zinc brasses in the Hirsch and
Lücke experiments to determine if the aluminum
alloys with high zinc content still behave like metals
with a high stacking fault energy.
The following experiments should all be
conducted in channel die compression tests to isolate
any shearing that may occur along the n-t direction or
any normal strain along the t direction. Through
computational analysis by Hirsch and Lücke, Taylor
model strain in the rolling direction (which is
predominantly active in channel die compression)
produces most of the characteristic texture properties
while shear contributes to the development of more
tertiary orientations, such as some of the S components
[2]. The magnitude of biaxial strain and additional
shear can vary between alloys of different composition
due to different hardening and deformation properties
which may be influenced by the pinning effect of zinc
particles in the matrix. This means that the magnitude
of shear and biaxial strain can vary between the
different alloy compositions under the same rolling
conditions which will ultimately lead to different
texture developments. Using a channel die will help
minimize any tertiary deformation that will occur and
promote normal Taylor type deformation in the rolling
direction. Also, if future experiments are done in a
rolling mill, the reductions should be calibrated so that
the samples experience the same amount of strain per
pass as opposed using a fixed reduction height. This
will promote a more accurate view of the texture
growth.
A stronger correlation in the transition from
the pure metal texture distribution to the alloyed type
distribution can also be observed if the alloys are rolled
at varying temperatures; in particular lower
temperatures. Future experiments with the samples
cooled with liquid nitrogen, similar experiments done
with plutonium-gallium, may show that a decrease in
temperatures can lead to a faster transition to the alloy
type texture with the addition of alloying elements [4].
The range of temperatures can lead to a better
understanding of how much influence the addition of
zinc has on the transition to alloy type texture
development through a decrease in stacking fault
energy. If the component volume fractions at the 30
Zn-wt% alloy is still higher than expected with varying
temperatures, the idea that the misfit strain that the GP
zones cause can influence significant crystalline
rotation or the effects spinodal decomposition has on
micro-recrystalization can be assumed to have a
significant effect on the texture development of the
material.
REFRENCES
[1] J. H. a. K. Lucke, "Mechanisms of Deformation
and Development of Rolling Textures in
Polycrystaline F.C.C Metals," Acta Metall Vol 36
No.11, pp. 2863-2882, Great Briton , 1988 .
[2] M. L. Weaver, Interpreting ODF's, Maryland:
University of Maryland, 2011.
[3] U. F. Kocks, Texture and Anisotropy: Preferred
Orientations in Ploycrystals and their effect on
Material Properties, New York: Cambridge
University Press, 1988.
[4] C. V. S. Lim, Length Scale Effect on the
Microstructural Evolution of Cu Layers in a Roll
Bonded CuNb Composite, Pittsburgh: Carnegie
Mellon University, 2008.
[5] J. A. Venables, The Electron Microspy of
Deformation Twinning, Cambridge: Journal of
Physicas and Chemistry of Solids, 1963.
[6] AMG Aluminum, "Titanium Boron Aluminum
Grain Refiners," Wayne, PA, 2014.
[7] MTDATA, "Calculated Al-Ti phase diagram,"
National Physics Laboratory, Teddington, 2003.
[8] Kaiser Aluminum, "Alloy 7068," Kaiser
Aluminum, Foothill Ranch, 2013.
[9] B. Diak, Interviewee, Personal inquiry.
[Interview]. Janurary - April 2016.