461 final - brad merkley

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.


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 (𝛾𝑆𝐹𝐸𝐴𝑙 ≅ 170 J/m2 and

𝛾𝑆𝐹𝐸𝐶𝑢 ≅ 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 745oC, then the grain refiner was


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 750oC then poured into the

mold. This was above the 730oC 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 990 oC, 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 750oC 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 850oC

and the grain refiner at 800oC. 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

740oC 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 1100oC would be required to

adequately melt the grain refiner [7]. The gas furnace

used could only reach temperatures up to 1000oC

safely, so temperatures above 900oC were used when

adding the grain refiner in for the sequential alloys.

The aluminum in the 30 Zn-wt% alloy was

heated to 950oC. The grain refiner was added at this

temperature, followed by the zinc at 940oC. The dross

similar to the second casting appeared after inserting

the grain refiner. The mold was then removed from the

furnace at 710oC and poured into the mold.

The aluminum in the final casting, the 60 Zn-

wt%, was initially heated to 1000oC. The grain refiner

and zinc were added at this temperature. Once at

660oC, 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 428oC for 96

hours to ensure adequate homogenization. The

samples were each wrapped in aluminum foil before

insertion. Table 4: Calculated annealing diffusion distances for

420oC. The actual salt bath averaged at 428 oC

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


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 -28oC 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 424oC and quenched in ice

water (temperatures of 4-0oC) 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.


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 -20oC 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


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




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


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]


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


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.

461 final - brad merkley

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