FSW Aluminum Alloy

Report for ASM Competition Submission

Professor G. B. Olson

ClientsThe Boeing CompanyFord Motor Company

Graduate AdvisorRyan Glamm

Student DesignersWilliam Chang

Boping LiuRyan GinderThomas Yu

MSE 390 Spring 2010

June 5th, 2010

Department of Materials Science and Engineering,Northwestern University

2145 Sheridan RoadEvanston, IL 60201

Table of Contents

Section I. Background…………..………………………………………………...3i. Introduction…………………………………………………………3

ii. Precipitation Hardening……………………………………………..4

iii. Friction Stir Welding………………………………………………..5

iv. Stress Corrosion Cracking…………………………………………..8

Section II. Team Organization….;……………………………………………….10

Section III. Previous Design Approaches………………………………………...11i. Performance Objectives…………………………………………….11

ii. Performance – Property Relations………………………………….12

iii. Property – Structure Relations……………………………………...12

iv. Structure – Processing Relations…………………………………....17

Section IV. Current Results……….……......……………………………...……..18i. Summary of Objectives……………………………………………...18

ii. Determination of Phase Fractions…………………………………....18

iii. Modeling the Yield Strength of Solutionized Alloy Prototypes…….23

iv. Experimental Phase Fraction & Precipitate Size…………………….25

v. Processing: Solutionizing, Welding and Aging……………………...28

vi. Conclusions and Final recommendations…………………………....34

Section V. References……………………………………………………………..36

2

I. Background

i. Introduction

Ford Motor Company and the Boeing Company are two major corporations in

manufacturing transportation vehicles. In recent years, lightweight materials have become a

main focus for development and aluminum alloys rank high in the list of lightweight alternatives

for materials currently being used in industry. The two companies formed the Ford-Boeing

Nanoalliance to investigate the potential application of lightweight high-performance aluminum

alloys as an integral part of both companies’ products1.

Ford, along with the rest of the automobile industry, currently employs high-strength

steels as the main component for ground vehicle production. However, the push to reduce fossil

fuel consumption and CO2 emissions has increased the demand for more fuel-efficient vehicles.

Consequently, Ford is considering the possibility of replacing steel with lightweight high-

performance aluminum alloys, which will significantly reduce the weight of the vehicles and

thus improve fuel efficiency. The automotive industry already use aluminum alloys as a

secondary component to steel. However, the existing automotive aluminum alloys lack in

strength and cannot meet the impact and roof crush legislations if applied as the main component

in automotive manufacturing1.

Boeing and the aircraft industry currently utilize 2xxx (Al-Cu-Mg) and 7xxx (Al-Zn-Mg-

(Cu)) aluminum alloys, both of which are high strength heat treatable alloys2. Unfortunately the

2xxx and 7xxx aluminum alloys are susceptible to corrosion cracking, which limits the

operational life of the aircraft. In addition, the 2xxx and 7xxx aluminum alloys are difficult to

weld. As a result, components have to be machined from larger pieces of materials, which

renders the processing costly and inefficient2. The non-heat-treatable 5xxx (Al-Mg) aluminum

alloys can be efficiently welded, but the alloys do not meet the strength and toughness required

for aeronautical applications8.

This year’s design will focus on meeting the requirements of the automotive alloy and

therefore the design specifications stipulated by Ford. The new aluminum alloy will have to be

high-strength, resistant to corrosion cracking, easily welded and affordable. The application of

this new alloy will give automobile manufacturers the ability to produce more fuel-efficient

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vehicles while maintaining the necessary strength and durability. Ford can potentially utilize the

alloy across its entire production line and increase the fuel efficiency of all Ford vehicles.

ii. Precipitation Hardening

2xxx and 7xxx aluminum heat treatable alloys are high strength alloys because they

utilize the formation of precipitates that significantly harden the alloys2. To achieve

precipitation, an alloy must be first heated to a temperature below the melting temperature where

the alloy will be in solid solution. The specific temperature for solution treatment depends on

the particular alloy and the concentration of the solute9. The alloy is then quenched in water to

room temperature. Due to lack of diffusion, no transformation occurs and the solid solution

remains mostly undisturbed at room temperature9. The alloy becomes a supersaturated solid

solution (SSSS) and there exists a driving force for precipitation of the equilibrium phase2.

Precipitation occurs as the alloy is aged at a certain temperature for a period of time.

During the aging process, the supersaturated solid solution decomposes and regions of

concentrated solute or vacancies become nucleation sites for the Guinier and Preston (GP)

zones2. These zones are approximately two atomic layers thick and can considerably strengthen

the alloy. GP zones are coherent with the matrix and minimize misfit strain and interfacial

energy by forming spherical zones. Therefore, even though the driving force for precipitation of

the GP zones is less than that of the equilibrium phase, the barrier to nucleation is less for the GP

zones and GP zones nucleate first9.

After the formation of GP zones, the precipitation of transition phases follows. Similar to

GP zones, the transition phases have lower activation energy barrier for nucleation than the

equilibrium phase. The free energy of the alloy decreases rapidly through the formation of the

transition phases to the equilibrium phase. The crystal structures of the transition phases are

intermediate between those of the matrix and equilibrium phase, which provide the transition

phases a high degree of coherence and thus lower activation barriers9.

Frequently, the equilibrium phase represents over-aging with decreased alloy strength. In

most cases, peak strength of the alloy is achieved with intermediate precipitates being the

dominant phase. In most systems, there may exist multiple strengthening phases2. A list of

common strengthening phases is summarized in the table below.

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Phase Approx

Composition

Typical aging

temperature

Alloy Primary

Strengthener θ' Al2Cu 175°C AlCuMg (Cu/Mg~4) Ω Al2Cu(Ag,Mg) 200°C AlCuMgAg (Cu/Mg~4) S' Al2MgCu 150°C AlCuMg (Cu/Mg~1) X Al2MgCu(Ag) 150°C AlCuMgAg (Cu/Mg~1) T Al6CuMg4 190°C AlCuMg (Cu/Mg~1/4) Z Al3CuMg(Ag) 200°C AlCuMgAg (Cu/Mg~1/4) Q Al5Cu2Mg8Si6 200°C AlCuMgSi β'' Mg5Si6 175°C AlMgSi η' MgZn 130°C AlZnMgCu

Table I. Common strengthening precipitates in heat treatable aluminum alloys2

Precipitation hardening is an effective strengthening mechanism and will be an essential

component in designing the desired aluminum alloy. The main challenge is to utilize the right

strengthener that can remain robust to the dynamic friction stir welding thermal cycle and remain

compatible with compositional constraints.

Non-heat treatable 5xxx series alloy utilize solid solution hardening and cold work as the

main strengthening mechanism. Cold work causes lattice distortion by other atoms in the solid

solution. Sub-micron dispersions of Cr, Mn, Ti, and Zr within the alloy inhibit grain growth

during processing, which means that the standard thermo-mechanical processing methods of

heat-treatable alloys do not apply to 5xxx alloys. These alloys have better formability and

corrosion resistance than heat-treatable alloys, but the strength of 5xxx alloys does not meet the

impact and roof crush test standards2. Without using the quenching and heat-treatment of heat-

treatable alloys, the strength of the alloy could be increased by adding alloying elements such as

Sc and/or thermomechanical treatments such as severe plastic deformation1. Alternatively, a

small amount of precipitation hardening could be utilized in an alloy with similar composition to

5xxx alloys to increase strength with the possibility of retaining the formability and stress

corrosion resistance seen in 5xxx series.

iii. Introduction to Friction Stir Welding

Friction stir welding is a solid-state welding technique to join alloys in that the change of

their original metallic properties is minimal. Friction stir welding consists of a constantly rotated

pin traversing the join between two metal plates. The process uses a metal pin, consisting of high

strength tool steel, on a rotating shoulder. The necessary thermodynamic and mechanical

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changes are produced by the pin and shoulder rotating at many hundreds of revolutions per

second. The additional pressure from the shoulder also results in the material being forged. For

each material considered, many factors must be optimized to achieve a good weld, such as tool

revolution speed, the pressure of the tool on the join, the transverse speed, and the geometry of

the fixture11.

Figure 1. Schematic of Friction Stir Weld Assembly

Figure 2. Schematic of Friction Stir Weld during welding

A primary advantage of friction stir welding is the ability to weld 7xxx series alloys.

Friction stir welding does not melt the alloy at the join line, which is ideal for 7xxx series alloys,

which are precipitation hardened and when melted by conventional weld techniques, the alloys

would lose strength, but more importantly, they form brittle oxide and eutectic phases. Use of

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the friction stir welding technique will lead to more efficient use of high strength aluminum

alloys, as machining of large plates from conventional welding techniques would be rendered

unnecessary. Since friction stir welding does not melt the parent-alloy, the residual stresses

present from a liquid solidification are reduced. The absence of such a liquid solidification from

the weld melt would avoid the formation of extremely brittle phases upon solidification, and a

lower incidence of weld defects such as lack of fill or cracking. Additionally, since there is no

weld pool, friction stir welding can be done in any direction and weld geometry11. In order to

perform this type of weld, the entire assembly is translated and rotated to achieve the complex

weld needed.

There are three regions of interest post-friction stir welding: the weld nugget, the thermal-

mechanical affected zone (TMAZ), and the heat affected zone (HAZ)11.

Figure 3. Hardness in different zones after Friction Stir Welding11

The weld nugget experiences the most plastic deformation and heating from the welding

fixture. The weld nugget can reach a temperature close to ¾ that of melting point of the alloy.

Due to this temperature and deformation, recrystallization occurs at a very rapid rate, causing a

very fine grain structure to form. Within the weld nugget reprecipitation of previous precipitates

should occur upon cooling, resulting in minimal degradation in mechanical properties. In the

above figure of hardness versus microstructure region, the weld nugget shows slightly lower

hardness and resistance to plastic deformation as the base material, due to the reprecipitation and

recrystallization of the alloy in the weld nugget12.

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In the thermal mechanically affected zones (TMAZ), the mechanical deformation from

the welding process causes the grains to be distorted, while the thermal energy is insufficient to

cause full recrystallization of the material. This can result in a reduction of mechanical

properties, as is demonstrated by the Figure 3. Due to the influx of thermal energy, precipitate

dissolution will occur, followed by reprecipitation through GP zones. As the newly formed

precipitates may not be the optimal size, the hardness of the material is lower in this region12.

In the heat affected zones (HAZ), there is not enough mechanical deformation nor

thermal energy to cause recrystallization of the material. However, the heat from the welding

process is sufficient to cause coarsening of precipitates. This coarsening of precipitates essential

to maintaining the strength of the alloy causes a loss in mechanical properties. This coarsening of

precipitates results in a lowered resistance to plastic deformation, shown by the lowered hardness

values compared to the weld nugget and the base alloy in Figure 312.

The base alloy is further away from the weld than the heat affected zone. In this region,

neither the mechanical deformation nor the heat generated from the welding process affects the

microstructure significantly. No grain deformation or coarsening of precipitates is observed in

this region, leaving the desired mechanical properties unchanged. In the above graph of hardness

versus microstructure region, the measured hardness is highest in this region12.

iv. Stress Corrosion Cracking

Stress corrosion cracking (SCC) is a phenomenon found in various materials, notably in

metal alloys. The term “stress-corrosion” refers to the occurrence of crack formation caused by

the accumulative effects of both mechanical stress and corrosion. No unifying theory explains

the phenomenon. Even though the appearance of SCC is similar across materials, the

mechanism of SCC is different depending on the material. In aluminum alloys, the SCC cracks

are shown to be always intercrystalline. There are two distinct cases of SCC in aluminum alloys.

In the first case, grain boundaries are continuously dissolved by electrochemical corrosion

caused by the disturbances of corrosion products resulted from applied stress, and in the second

case, a discontinuous spontaneous crack formation occurs without any visible corrosion

interference10.

It has been shown experimentally that increased Zn content at the grain-boundary

decreased the life of the alloy under applied stress. Figure 4 demonstrates experimental results

8

where time to fracture in an Al-Zn-Mg alloy decreased with increasing Zn content at the grain

boundary and thus it can be inferred that grain boundary compositions contributes to SCC10.

Figure 4. Relationship between Zn content and time to fracture in Al-Zn-Mg10

Additionally, Zhang2 used the theoretical Rice-Wang model for grain boundary

embrittlement and the method outlined by Wu, Freeman, and Olson for using first principles to

determine boundary embrittlers and cohesion enhancers. Zhang found that Zn is a grain

boundary embrittler, confirming previous experimental works2. This finding has significant

implications in designing the alloy by eliminating Zn as a potential precipitate former to reduce

SCC and by confirming with previous experimental results that Zn is the main perpetrator of

SCC in 7xxx aluminum alloys.

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II. Team Organization

For Spring 2010, the MSE390 FSW Aluminum team is composed of four MSE

undergraduate students Boping Liu, Thomas Yu, William Chang, Ryan Ginder and one graduate

student advisor, Ryan Glamm, who will provide guidance for all technical aspects of the project

and supervise the progress of the group. The specific areas of expertise were delegated to each

team member based on each member’s personal preference as well experience in each area.

Boping Liu will be the lead for compiling and editing the overall report. Thomas Yu, William

Chang and Ryan Ginder will be responsible for the strength model and Thermocalc simulations..

In addition to their primary responsibilities, team members also have a secondary responsibility.

The RAM chart below summarizes the specific distribution of tasks where 1 indicates primary

responsibilities and 2 indicates secondary responsibilities.

Tasks Boping Liu Thomas Yu William Chang Ryan GinderReport & Organization 1 2

Strength Modeling 2 1Thermocalc Modeling 1 2

Laboratory Work 2 1Table II. Distribution of responsibilities among team members

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III. Previous Design Approachesi. Performance Objectives

Strength

For the designed alloy to become applicable in the automotive industry, the strength of

the alloy must meet the impact and roof crushing standards. Ford has stipulated a minimum

requirement of 175 MPa for yield strength and 275 MPa for tensile strength after a 2% forming

strain and a 30 minute paint bake at 175°C2.

Weldability

To improve processing efficiency, the designed alloy must be compatible with Friction

Stir Welding (FSW). The designed alloy should resist deterioration of mechanical properties at

the weld caused by FSW. Under the FSW, three different zones form around the weld as a result

of highly irregular deformation and temperature cycles. Among the three zones in a heat

treatable alloy such as 7050, the heat affected zones (HAZ) can suffer as much as 30% reduction

in hardness compared to the base alloy2. Therefore the designed alloy should retain at least 85%

hardness compared to the base alloy at the weld, especially the heat affected zones8.

Stress Corrosion Cracking Resistance

For the automotive design, the minimization of the SCC susceptibility is a high priority.

The designed alloy should minimize the Zn component or completely exclude Zn from the alloy

composition and it should replicate the SCC lifetime of the 6061 aluminum alloy (Al-Mg-Si-

Cu)2.

Affordability

In order for the designed alloy to be effectively commercialized, the cost of the alloy

should be comparable to the materials currently being used by the industry. For automotive

applications, the affordability requirement stipulates that expensive alloying elements that

require special scrap recycling procedure cannot be included in the design, which effectively

eliminates Scandium and Lithium. To reduce processing costs, the elimination of quenching

after solution treatment is desired. Instead of quenching, constant cooling rates can be employed

as an alternative to achieve similar microstructures formed upon quenching.

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ii. Performance – Property Relations

This project will utilize a systems approach of material design2. Figure 5 shows the

mutual dependence of processing, structure, properties and performance. This flow chart was

first developed in the first design iteration in MSE 390 during Spring Quarter of 2009. The

ultimate goals of the project define the properties the designed alloy must possess, the properties

in turn translate into the necessary structures that can exhibit those properties, and the structures

will come to form by the various processing steps shown on the left of Figure 52. The FSW

aluminum group in MSE 390 in Spring Quarter 2010 will continue employ the design approach

outlined by Figure 5.

Figure 5. Processing, structure, properties and performance relations2

iii. Property – Structure Relations (Strength Modeling)

The strength model will be used to determine the composition and desired size of

precipitates in the matrix. For this project, the strength model is based on the work of Myhr et al.

for Al-Mg-Si alloys15. Similar models can also be found in other Al based alloy systems such as

12

Al-Mg-Si-Cu, Al-Zu-Mg, and Al-Zn-Mg-Cu16, 17, 18. This model is capable to accurately predict

the yield strength and hardness of the materials from modeling, which provides discretized

particle size distribution and solute concentration remaining in the matrix. This model is

described by Equation 1,

(1)

Where σy is the yield strength, σgb is the contribution from grain boundary strengthening,

τppt is the contribution from solid solution strengthening, τi is the intrinsic strength, and τD is the

strengthening based on the dislocation density. For the purpose of this project, values for the

intrinsic strength, dislocation density strengthening, and grain boundary strengthening will be

approximated using standard literature values for Al-systems3.

The grain boundary strengthening is described by the Hall-Petch equation below:

€

∆σ g b = k 1

d (2)

where d is the average grain diameter and k1 is an experimental constant.

For solid solution strengthening, the model used is shown by Equation 3.

€

∆τss = k jc jn∑ (3)

Where cj represents the atomic fraction of each element in solution and kj represents the

element’s strength contribution. For magnesium, kMg = 590 MPa and n = 1 will be used9.

To determine the content and desired size of the precipitates in the matrix, the

precipitates are approximated as spherical. The strength model shown in Equation 4 can then be

applied to determine the precipitation hardening based on the radius and volume fraction of the

particles.

€

∆τ p p t = Kf 0 . 5

r (4)

Where τppt is the strengthening contribution of the precipitates, K is the proportionality

constant, r is the radius size of the precipitates, and f is the phase fraction of precipitates in the

material.

The strengthening contributed by the dislocation density is described by Equation 5:

€

∆τD =αGb ρD (5)

Where a is a constant with a value around 0.3 for Al based alloys and ρD is the dislocation

density.

13

With the aid of microstructure simulation through experiments, the strength model can be

tested and compared to micro-hardness measurements. The parameters for the strength modeling

equations aforementioned are taken from other published work modeling 7xxx series alloys and

are shown in Table III. Dislocation strengthening was considered for this calculation using a

dislocation density an order of magnitude lower than the maximum dislocation density observed

in literature. A experimentally determined conversion of Y. S. (MPa) / 2.1 = VHN was used to

convert predicted yield strength to Vickers hardness number.

Parameter Value ReferenceTaylor Factor, M 3 13

Intrinsic strength, τi 15MPaHall-Petch constant,

k10.08 7

Average Grain size,

d

Nugget - 6μm

HAZ- 27 μm21

kMg 29 MPa/wt%2/3 20kCu 46.4 MPa/wt

%2/3 22

kZn 15 MPa/wt%2/3 22n 1 22

Burgers Vector, b 2.86Åβ 0.43 5

Table III. Parameters for strength modeling

Strengthening in Base Alloy and Heat Effected Zone (HAZ)

As Zn is avoided due to stress corrosion cracking concerns, Cu is the only suitable

candidate out of the traditional alloying elements that would produce strengthening precipitates.

From the vertical section of the Al-Mg-Cu system shown in Figure 6, it can be identified that as

Mg level is raised, the strengthening precipitates goes from θ to S and then to Al6CuMg4-T

phase.

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Figure 6. Vertical Section of Al-Mg-0.1at%Cu from Thermocalc2

The T-phase is the low temperature precipitate in figure 6 and thus a low temperature

precipitate is desired to strengthen the base material as well as dissolving and re-precipitating in

the high temperature regions around the weld2. It is advantageous for the low temperature

precipitate to re-precipitate at a fast rate and therefore the diffusivity of the precipitate should be

higher in Al than that of Cu2. Figure 7 shows the diffusivity of different alloying elements in Al.

Figure 7. Plot of diffusivities of different elements in Al2

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The strengthening efficiency of the T-phase has been shown to increase by microalloying

with fast diffusing Ag. With the addition of Ag, a new precipitate, Z-phase, would form. From

Figure 7, the relative diffusivity of Ag (5.7) is significantly more than the diffusivity of Cu

(0.14), and Ag additions are expected to facilitate GP zones precipitation and strength recovery

in regions affected by the welding process2. To further study the effect of Ag and improve the

accuracy of current modeling, Ag will be included as a potential alloying element in the design.

Another potential candidate for microalloying is Sn, because it is also a fast diffusing low

temperature precipitate in Al alloys. Figure 8 shows the strength hardening along with enhanced

aging kinetics as a result of Sn additions. Previous works have shown that upon quenching, Sn

forms nanoclusters that act as nucleation sites for refined dispersion in Al-Cu alloys2, but it is

unclear if the same mechanism will be present in Al-Mg-Cu alloys with a high Mg constituent.

As a result, Sn will be included in the design as well. The inclusion of Ag and Sn has the

potential to significantly improve the base alloy strength and recovery in the weld region, since

Ag and Sn precipitates precipitate at low temperatures with fast kinetics. For automotive

applications, the fast kinetics can potentially enhance the hardness of the alloy during the paint

bake step, which is performed at low temperatures during a short cycle.

Figure 8. Hardness-time plots to show effects of Sn additions.

A high temperature phase is desired to pin grain boundaries during processing The

precipitate should have a low diffusivity in Al and withstand temperatures up to ~270 °C. From

Figure 7, Sc and Zr are the most qualified candidates. Sc has low solubility and high cost makes

it inapplicable in the automotive design. Zr on the other hand is much less expensive and forms

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precipitate in the form of Al3Zr. At the solutionizing temperature, Al3Zr is a stable phase that

pins grain boundaries and thus control grain growth2.

iv. Structure – Processing Relations (Thermocalc Modeling)

As introduced above, the structure-processing relations for this material are vital in

achieving the performance objectives. In order to investigate the structure-processing relations,

Thermocalc will be used. Thermocalc is a software package that performs thermodynamic

calculations using an extensive database of thermodynamic data. Thermocalc is capable of

computing phase diagrams of equilibrium binary, ternary, and multi-component systems.

Calculations for the driving force of nucleation can also be made. Users of Thermocalc are

capable of inputting specified elements, phases of interest, and components. Once the regions of

interest are inputted, conditions such as temperature, pressure, and component mass fraction

must be satisfied.

Thermocalc modeling is heavily dependent on prior thermodynamic data compiled in the

databases. Aluminum alloys are usually strengthened via metastable phases. By doing

experiments on the alloys by the proposed modeling techniques, essential data can be acquired to

calibrate and improve the Thermocalc models for the desired metastable phases.

A starting point for this friction stir weldable Al design will be a 5xxx series weldable

alloy composition. Low temperature precipitates are necessary to maintain alloy strength after

welding in the heat affected zone. By adding small quantities of elements such as Ag and Sn, a

fast precipitating strengthening dispersion can be formed. By further optimizing the processing

temperatures, such as the solution heat treat temperatures, or the precipitate aging temperatures,

the optimal precipitate size can be achieved, allowing the precipitates to act as efficient

strengtheners.

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IV. Current Resultsi. Summary of Objectives

To achieve the performance objectives, the following criteria have been established to

guide the initial design process. Previous design work had resulted in an untested prototype, from

which many of the current experimental work was performed on.

• The strength of the alloy should exceed the stipulated yield strength of 175 MPa after a

2% forming strain and a paint-bake of 180 °C.

• For the purposes of this design, the strength increase delivered by the 2% forming strain

will be approximated by literature values of comparable Al alloys.

• The formation of the beta-Mg phase in the Al alloy at the aging temperature should be

prevented. As the alloy will be solutionized for a period of time and then heat-treated, the

microstructure after the designed heat treatment should be free of beta-Mg. Therefore, it

is undesirable to have beta-Mg at the aging temperature, as it forms at the grain

boundaries and embrittles them..

• The paint bake used by Ford should also be integrated into the thermal cycle that the

precipitates undergo. By incorporating slightly under-aged precipitated into the material

before the paint bake, a paint bake at 180 °C would lead to the precipitates reaching their

peak size and the alloy reaching its peak hardness after the paint bake.

• Zirconium will be added at 0.1 mole percent to precipitate Al3Zr, which is a high

temperature precipitate. The solid solutionizing temperature should be high enough to

have only Al3Zr and the FCC phases. Al3Zr will pin the grain boundaries at the

solutionizing temperature, preventing grain growth.

• In order to lower processing costs, the quenching step should be removed from the

processing of this alloy. After the solution treatment, the temperature will be lowered to

age precipitates without the intermediate quenching step

ii. Phase Fraction Determination

Phase Diagrams of Cu & Mg in Al

First, the phase diagram for the binary system of Mg-Al was calculated using

Thermocalc. For this phase diagram, copper stabilizes the T phase which in turn destabilizes the

18

beta-Mg phase. The figure below shows the effect of copper on the formation of the beta-mg

phase at the aging temperatures of 463K. Phase diagrams calculated at 463 K, 483 K and 500 K

did not show significant changes from temperature effects. Figure 9 displays the effects of Cu on

the formation of beta-mg for an aging temperature of 463K.

Figure 9 is a phase diagram at a constant temperature, with varying mole fractions of Mg

and Cu. It shows that after 1 mole % of Cu in the alloy, Mg can exceed 10 mole percent without

forming the beta-mg phase at an aging temperature of 463K. However, this is a very large

amount of alloying elements in Al, and while may lead to an extremely high yield strength,

would most likely result in poor ductility. The goal is to reach the yield strength with minimal

precipitation strengthening necessary to retain ductility and formability for welding purposes.

Figure 9. Phase diagram of Cu & Mg in Al at aging temperature 463K

Phase Diagram of Mg-Al with Fixed Cu

Following the above analysis, a solutionizing temperature and appropriate composition

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must be determined. A priority for the alloy design is to ensure that at the solutionizing

temperature, only the FCC and Al3Zr phase are present, and in the aging regime the FCC and tau

phase must be present. The solutionizing temperature will be limited by the melting point of the

alloy. Since higher solutionizing temperatures will result in higher costs, keeping the

solutionizing temperature as low as reasonably possible can lead to considerable.

The following figures depict the effects of copper on the pseudo-binary phase diagram

between magnesium and aluminum. The x-axis depicts the mole fraction addition of magnesium,

and the y-axis is temperature. For the designed alloy, a suitable aging temperature is determined

to be between 460-500K. The suitable solutionizing temperature is determined from the

following graphs.

Figure 10 shows the effect of 0.25 mole % on the pseudo-binary phase diagram

calculated. The addition of copper introduces the formation of the tau-phase along with the FCC

region. The FCC region is present above 700 K for an alloy with 10 mole percent of Mg. The s-

phase to the left of the FCC + tau phase is also a strengthening phase. After solution treatment,

the alloy will be passed through the s-phase+ tau+ FCC region quickly. Again, an aging

temperature of approximately 500K results in approximately 3 mole percent of Mg in solid

solution.

Figure 10. Al-Mg pseudo-binary phase diagram with 0.0025 molar fraction Cu

20

Figure 11 shows an increase the copper content in the alloy to 0.005 molar fractions,

effectively double the copper content from the previous example. The amount of Mg that can be

added increases since Mg is entering the T phase. This insures that the alloy stays in a FCC

regime during solutionizing the material. The FCC region at high temperature turns into a S-

phase region. Therefore, there are limits in how much Cu can put into the alloy. From this

exercise, it can be determined that less than 0.25 mole percent of combined Ag and Cu should

allow enough Mg to enter solid solution for strengthening. However, 5 mole percent of Mg is a

high addition of Mg, and 3 mole percent should be sufficient. The aging temperature would be

170 - 190 °C, and the solutionizing temperature can be as low as 473 °C. Again, the solutionizing

temperature is constrained by the melting point of the alloy, and the fact that the higher the

solutionizing temperature, the higher the overall production cost of the alloy.

Figure 11. Al-Mg pseudo-binary phase diagram with 0.5 molar percent Cu

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Phase Diagrams of Ag & Mg in Al

Figure 12 shows how the addition of silver to the previous Al-Mg-Cu-Zr alloy

compositions changes the solubility of Mg. This phase diagram shows the effect of 0.25 mole

percent of both silver and copper. The addition of silver stabilizes the tau precipitate phase field,

destabilizing beta-Mg formation. The addition of silver can generate more precipitates and

without the formation of beta-Mg when using 3 mole percent Mg content.

Figure 12. Vertical section of Al-Mg-Cu-Ag with 2.5×10-3 mole fractions of Cu and Ag

Figure 13 shows that at high copper and silver mole percents, the beta-Mg phase is

entirely suppressed in the region of interest. The combined mole percent of Cu and Ag is 0.01.

This is a very high addition of silver and copper into an aluminum alloy, and shows the strong

effects of adding either to increase the solid solubility of Mg in Al.

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Figure 13. Vertical section of Al-Mg-Cu-Ag with 0.5 mole percent of Cu and Ag

Phase Fraction Calculations

From the previous section a viable composition for the alloy was calculated to be 0.25

mole percent Ag, 0.25 mole percent Cu, 3 mole percent Mg, 0.09 mol percent Zr with the

balance being Al. The aging temperature was determined to be 170-190 °C , and the

solutionizing temperature can be as low as 473 °C to lower processing costs. Local Electrode

Atom Probe (LEAP) work was done by Glamm to experimentally determine the composition of

the matrix and precipitate phases in a previous prototype with the same components. These

experimental values were input to Thermocalc to determine the theoretical solute levels and

phase fractions for further modeling.

Phase Cu (mole fraction) Ag Phase FractionBulk 0.0025 0.0025

Matrix 0.0002 0.0002 0.973Precipitate 0.015 0.12 0.027

Table IV. Phase fractions of precipitates formed by Cu and Ag

ii. Modeling the Yield Strength of Solutionized Alloy Prototypes

To begin modeling the precipitation strengthening, a conversion between the yield

strength and hardness of the prototype alloy is tested. As mechanical tensile tests are more time

23

and material consuming than micro-hardness measurements, this conversion could be very

useful. Figure 14 relates the magnesium content and yield strength of non-precipitation

strengthened 5xxx series alloys. The prototype that was used to calibrate the precipitation

strengthening model had a magnesium content of 2.99 weight percent. Using the linear

regression equation from the experimental data in the figure below, a yield strength of 113.3

MPa was calculated. This prototype alloy was then solutionized and then aged for 24 hrs to the

optimal precipitate size. The hardness of the solutionized alloy had a hardness of 62 HV. The

hardness difference caused by precipitation hardening was 15.24 HV. By assuming a linear

relationship between the Vickers hardness and the yield strength, the proportionality constant

was determined to be 2.1.

Figure 14. Yield strength (MPa) vs. Mg content (wt%) of solutionized and age 5xxx alloy

The strength model, as mentioned before, has five components. The five components are

the intrinsic strength of aluminum, precipitation strengthening, grain boundary strengthening,

dislocation strengthening and solid solution strengthening. By determining the prototype yield

strength without precipitation strengthening, and comparing it to the experimental hardness of a

solution treated sample in the above figure, a proportionality constant between strength and

hardness could be calculated.

24

0 1 2 3 4 5 60

50

100

150

200

250

300

350

400

450

1 23

4569

12

f(x) = 27.8x + 30.24R² = 0.95

f(x) = 24.97x + 176.49

Mg wt%

Yiel

d St

reng

th (M

Pa)

Solutionized

Peak-agedCold-worked

Annealed

In the strength model without precipitation hardening, the prototype has a yield strength

of 127.4 MPa. Experimentally, the micro-hardness was measured to be 62 HV. Thus, the

proportionality constant was found to be 2.05. This is in approximate agreement with literature

values, in which the proportionality constant between micro-hardness and yield strength is

between 2-3.5. In order to determine the precipitate strengthening contribution, a simplified

equation of the precipitation hardening was used.

Δτ = K*f0.5/r

Δτ is the increase in shear strength, k is the proportionality constant, r is the peak radius size of

1.4 nm, and f is the phase fraction of precipitates in the material. To determine both the peak

radius size and the phase fraction to calibrate K in this model, atom probe tomography was used.

Since the atom probe can experimentally determine both the phase fraction and precipitate size,

it is ideal for calibrating the precipitation strengthening model.

Strengthening

Mechanism

σ Contribution (MPa)

Intrinsic τ of Al 15Grain Boundary 32.66

Dislocation 14.7Solid Solution 5

Total (Solutionized) 127.4Table V. Strength contribution from the five components of the strength model

iii. Experimental Phase Fractions & Precipitate Size

Atom probe results were obtained by graduate coach Ryan Glamm using the local

electrode atom probe (LEAP) located at Northwestern University. The atom probe allows the

experimental determination of single atoms in the material. This allows for the determination of

the compositions of the precipitate, the matrix and the overall alloy. Also, the atom probe allows

the precipitate phase fraction and radius to be determined. The sample tested had a nominal

alloy mole percent composition of 3.9% Mg, 0.1% Ag, 0.1% Cu, 0.02% Sn, and 0.04% Zr. The

balance is aluminum. However, the atom probe calculated a different composition, shown in the

table below. The bulk composition of magnesium is much lower; this is due to magnesium burn-

off in the alloy casting process.

25

Phase Al (mol%) Mg (mol %) Ag (mol %) Cu (mol %)PPT 62.35 20.32 16.41 0.92

Matrix 97.35 2.61 0.04 N/ABulk 97.16 2.66 0.08 0.1

Table VI. Prototype Alloy composition obtained from atom probe

A 3-D reconstruction of the alloy is shown below. The average precipitate radius is

approximately 1.4 nm. The precipitates are composed of Al, Mg, Ag, and Cu. The exact

compositions are indicated in the table above. The atom probe was unable to find any Sn clusters

in the alloy, likely due to such small concentration.

Figure 15. 3-D reconstruction of the alloy by atom probe

Two samples with different aging times were run to acquire two different data points to

calibrate and test the precipitation strengthening model. The 2 hrs sample was of prototype 4

composition; the 24 hrs was of prototype 2 composition. Both compositions are similar;

prototype 4 has a higher solute level than prototype 2. The composition of prototype 4 is given in

Table VI. The important observations are included in the table below.

Aging Time &

Temperature (hrs)

Precipitate Radii

(nm)

Phase

Fraction (%)

Hardness increase

(HV)

Predicted Hardness

increase (HV)2 hrs @ 170 °C 1.24 0.28% 17.27 17.2224 hrs @ 170 °C 1.4 0.26% 15.24 Used to calibrated

Table VII. Experimental parameters of alloys with different aging times

26

Using these experimental parameters, it is determined that Kloop = 2.69×10-7 MPa/m.

Using the calibrated model, the properties of the sample with an aging time of 2 hrs at 170 °C

can be predicted. The expected hardness increase is very close to the actual hardness increase,

providing some validation for the method used to calibrate the model. Now that the precipitation

strengthening method has been established, a necessary phase fraction must be calculated to

bring the non-precipitated alloy to the final goal of 175 MPa, keeping in mind that the welding

and subsequent processing will lead to certain deterioration of mechanical properties. Using the

model, a phase fraction of 0.7 percent of precipitate was deemed adequate for sufficient

strengthening. The predicted yield strength increase of such a phase fraction is 54.45 MPa. This

relationship is shown in the Figure 25. The indicated phase fraction would give the necessary

strength increase to reach the yield strength goals.

Figure 16. Change in yield strength vs. precipitate phase fraction

Now the question remains, how much solute must be included into the alloy to achieve

the desired phase fraction and strengthening? The Thermocalc plot below shows the effect of

adding solute, in this case copper and silver, to the material. Figure 16 was constructed by adding

capillary energy to the thermodynamic calculations to match the peak hardness and phase

fraction measurements from the LEAP data. The LEAP data showed that the phase fraction was

0.258 percent with 0.05 mole percent Cu and 0.05 mole percent Ag. The capillary energy that

Thermocalc uses for thermodynamic calculations was adjusted until the phase fraction output by

Thermocalc matched experimental results. The Cu and Ag solute levels were set to be

27

0.00E+000 2.00E-003 4.00E-003 6.00E-003 8.00E-003 1.00E-0020

10

20

30

40

50

60

70

Precipitate Phase Fraction vs Change in Yield Strength

Precipi tate Phas e Fraction

Cha

nge

in Y

ield

Stre

ngth

equiatomic, so that there is a one to one ratio of Cu and Ag in this alloy composition. The

magnesium content is at 3.5 mole percent, while the temperature is 190 °C. The liquid in this

calculated diagram is an artifact. The Al3M_D023 is the Al3Zr grain refiner. From the figure,

0.1 mole percent of Ag and 0.1 mole percent of Cu is expected to achieve the required phase

fraction.

Figure 17. Phase fraction as a function of Cu and Ag levels

iv. Processing: Solutionizing, Welding and Aging

Processing is important in determining the size of the final precipitates. Additionally, the

alloy must be weldable, and will undergo a paint-bake treatment (180 °C for 30 mins),

processing conditions must be determined to maximize hardness and thus yield strength after

these processing steps. Two final recommendations will be made: one processing routine with

the conventional quench and age, and another without the quench and aging step, but replaced by

a predetermined cooling rate. In order to determine the homogenization temperature, a Scheil

model of the phase fraction solid versus temperature was constructed. Figure 18 shows that

homogenization should take place below 470 °C to prevent incipient melting. After

28

homogenization, the equilibrium phase diagram can be used to determine the solutionizinging

temperature of 470 °C. The liquidus temperature of the alloy is shown to be 650 °C.

Figure 18. Scheil Model of temperature vs. phase fractions

The following graph shows experimental data for two aging treatments. The 170 °C and

190 °C data set corresponds to an alloy that was quenched and then aged at 170 °C and 190 °C

respectively. The alloy aged at 190 °C has a lower peak aging time at 5 hours, while the alloy

aged at 170 °C material must be aged for 24 hours to reach maximum hardness. These plots also

show the effects of over-aging to the hardness of the material; over-aging does not decrease the

hardness substantially, thus the peak hardness is robust.

29

Figure 19. Alloy hardness with different aging treatments

The welded data shows the effects of putting the material through a welding cycle

simulated in the dilatometer. A thermal profile is shown in Figure 20. The thermal profile

generated by the dilatometry tracks the literature welding thermal profile well23; the peak

temperature has some fluctuations, which can lead to increased stresses in the material. This is

the first instance of simulating a FSW thermal cycle in a dilatometer that we are aware of.

Figure 20. Alloy thermal profile during weld simulation

30

0.1 1 10 10060

65

70

75

80

85

90

95

100

105

Effect of Different Aging Temperatures on Alloy Hardness 170C 190C

Log(Aging Time) (Hours)

Har

dnes

s (H

V)

Aging Time (hrs)

0 200 400 6000

50

100

150

200

250Dilatometry Welding Thermal Cycle

ActualIdeal

Time (s)

Tem

pera

ture

(C)

The effects of the weld and the subsequent paint-bake step are shown in Figure 21. Here,

the effects of the weld and then the paint-bake are shown for different samples aged at different

times. Initially, with solutionized samples, the change in hardness is positive. The weld treatment

causes precipitation within the solutionized material, raising the hardness. The paint-bake further

causes further precipitate growth, leading to an additional increase in strength. At the peak aging

time of 5 hrs, it is observed that there is a significant decrease in hardness of 13 HV. This

corresponds to a decrease of 25 MPa. This is likely due to dissolution and coarsening and then

insufficient re-precipitation of the material. The paint-bake step shows an increase in hardness

compared to the welded samples, since the precipitates have grown. The over-aged precipitate

exhibits stable hardness in heat-affected zone; for this reason, the design should attempt to over-

age the precipitates since the strength decrease is not substantial from the peak-aged state and are

more stable in the heat-affected zone.

Figure 21. Various welding and paint bake treatments on alloy hardness

31

0.1 1 10 10060

65

70

75

80

85

90

95

100

105

Effect of Different Welding and Paint Bake Treatments on Alloy Hardness 190C W/ Weld W/ Weld, Paint Bake

Log(Aging Tim e) (Hours )

Har

dnes

s (H

V)

Solutionized

Peak-agedOver-aged

Aging Time (hrs)

To investigate the feasibility of eliminating the need for a quench and aging step,

different cooling rates were selected to bring the material out of the solutionized state. The three

cooling rates are demonstrated below; there is a fast cooling rate of 1.8 °C/second, a medium rate

of 0.9 °C/second, and a slow cooling rate of 0.45 °C/second. It is expected that the formation and

final size of the precipitates upon reaching room temperature will be a function of the amount of

growth experienced in the ideal temperature range. Ideally, a cooling rate will be determined

which would give the same hardness after processing as a quenched and peak-aged alloy. This

would lead to energy efficiency, power saving and quicker alloy fabrication time, which are all

driven by cost.

0 500 1000 1500 20000

100

200

300

400

500Varying Cooling Rates to Eliminate Quench and Aging

Medium - 0.9 C/sSlow – 0.45 C/sFast – 1.8 C/s

Time (s)

Tem

pera

ture

(C)

Figure 22. Various cooling rates in place of quench and aging

The plot below shows the end hardness for these three alloys under different cooling

rates. It is interesting to note that the medium cooling rate of 0.9 °C/second delivered the highest

hardness. This plot can be explained in the following manner, using two extremes. At a slow

cooling rate, precipitates would form in the high temperature region. However, the number of

precipitates will be low due to the low driving force since the temperature is not far removed

32

from the solvus temperature. The precipitates will then be allowed to coarsen beyond the peak

radius, leading to a small number of large precipitates. This will lead to ineffective

strengthening, according to the precipitation strengthening model. By using a very fast cooling

rate, the time window by which the precipitates have to form and grow is small, and so

insufficient precipitation and growth will occur with a fast cooling rate. While the optimal

cooling rate may not have be achieved, it has been demonstrated that replacement of the quench

and age step is a viable option with the correct alloy composition. The paint bake treatment has

been shown to lower alloy hardness after controlled cooling. With cold work, controlled cooled

alloys should be able to meet the strength goal of 175 MPa. Future precipitate kinetic modeling

is planned in order to determine this optimal cooling rate for peak hardness.

Figure 23. Alloy hardness with different cooling rates and paint bake treatment

v. Conclusion and Final Recommendations

Based on extensive experimental data and modeling, the design is as follows. The final

composition is given in the table below. This final composition was determined using the

strength model, and was over-designed so that after the decrease in hardness caused by welding

would still allow realization of the yield strength goal of 175 MPa. The end product should have

33

60

65

70

75

80

85

Effect of Controlled Cooling Rate on Prototype 4 Hardness

No Paint Bake 0.45 C/s

PB 0.45 C/s No PB 0.9 C/s

PB 0.9 C/s No PB 1.8 C/s PB 1.8 C/s

Cooling Rate (Cels ius /sec)

Har

dnes

s (H

V)

a phase fraction of 0.007 or 0.7% of T-phase precipitates. The radius size after processing should

be close to 1.4 nm. The bar graph summarizes the strengthening mechanisms and contributions.

The material has been over-designed to compensate for the decrease in strength caused by

welding (~5 MPa decrease). The optimal theoretical prototype is very similar in composition to

the previously experimentally designed prototype used for experimental calibration; this should

be taken as a theoretical confirmation that this prototype is the optimal alloy to meet automotive

goals.

Figure 24. Strength contributions and the final alloy strength

Two different processing steps have been designed: one with a conventional quench and

aging, and another with the quench and aging step replaced with a gradual cooling from the

solution state. The parameters for the conventional quench and aging are as follows: the material

should be solutionized at a temperature of 470°C until all solute is in solution. The material

should then be quickly quenched to room temperature. The material should then be aged at 190

34

0

50

100

150

200

250

Strengthening Contributions and Final Goal Strength

Intrinsic Strength of Al Precipitate Strengthening

Dislocation Strengthening

SS Strengthening

Mechanism

Yiel

d S

treng

th C

ontri

butio

n (M

Pa)

-15 MPa (weld+ bake)

GOAL

-5 MPa (weld + bake)

°C for 10 hours. This will result in an alloy that is slightly over-aged, with precipitates stable in

the HAZ generated by friction stir welding. The material can then be welded and paint-baked at

180 °C for 30 minutes to reach its peak hardness and yield strength. The yield strength should

reach the goal of 175 MPa, factoring in the drop in hardness due to welding.

The alternative process of steady cooling from the solution state should use a cooling rate

in the vicinity of 0.9 °C/ second, which yields by far the highest hardness. However, the material

hardness was still a substantial fraction of the hardness achieved with a conventional quench and

aged process. Nevertheless, with additional cold work, the strength goal of 175 MPa is possible

after the paint-bake and cold work treatments. It has been demonstrated that with just controlled

cooling, the strength was within 10% of the strength goal. More cooling rates should be tested in

order to find the optimum which can potentially lead to complete replacement of the quenching

and aging step, which would result in significant savings to the manufacturers the alloy.

Element Al Mg Cu Ag Zr Sn

Mole % 96.7 3 0.1 0.1 0.09 0.01Table VIII. Recommended final alloy composition

Quench and Age Processing ConditionSolutionizing Time & Temperature 470 °C for 6 hours

Aging Time & Temperature 190 °C for 10 hoursPaint Bake Time & Temp 180 °C for 0.5 hoursPrecipitate Phase Fraction 0.7 mol %

AL3Zr Phase Fraction 0.01 mol%Predicted Strength 195 MPa

Table IX. Recommended processing conditions for conventional quench and ages

Processing Step ConditionSolutionizing Time & Temperature 470 °C for 6 hours

Alternative Controlled Cooling Rate 0.9 °C/ secondPaint Bake Time & Temp 180 °C for 0.5 hours

Predicted Strength (w/ cold work) <175 MPa Table X. Recommended processing conditions for controlled cooling rate processing

35

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