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NBS: Processing/Microstructure/Property

Relationships in 2024 Aluminum Alloy

Plates

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

National Measurement Laboratory

Center for Materials Science

Washington, DC 20234

Phone: (301) 921-2891

January 1983

Technical Report

Joint Sponsorship by NBS/NASA

NASA Government Order W-14,637; Modification 3

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

NBSIR 83-2669

MBS: PROCESSING/MICROSTRUCTURE/PROPERTY

RELATIONSHIPS IN 2024 ALUMINUM ALLOY

PLATES

L. K. Ives, L. J. Swartzendruber, W. J. Boettinger, M. Rosen,

S. D. Ridder, F. S. Biancaniello, R. C. Reno, D. B. Ballard and

R. Mehrabian

U S. DEPARTMENT OF COMMERCE

National Bureau of Standards

National Measurement Laboratory

Center for Materials Science

Washington, DC 20234

Phone: (301) 921-2891

January 1983

Technical Report

Joint Sponsorship by NBS/NASA

NASA Government Order W-14,637; Modification 3

U.S. DEPARTMENT OF COMMERCE, Malcolm Baldrige, Secretary

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director
'
PROCESSING/MICROSTRUCTURE/PROPERTY

RELATIONSHIPS IN 2024 ALUMINUM

ALLOY PLATES

ABSTRACT

Nondestructive evaluation (NDE) using eddy-current conductivity

and hardness measurements form an essential part of the quality control

of aluminum alloy plates used for aerospace vehicles. The relationships

between the NDE measurements and the important mechanical properties

are affected by a large number of variables including: chemical composition, cast structure, ingot scalping, solution heat treatment and

quenching, mechanical working, and aging treatment. At the request of

the National Aeronautics and Space Administration, a number of these

relationships has been explored for 2024 aluminum alloy. This work is

a continuation of our previous efforts on 2219 aluminum alloy and is

motivated by a concern that a number of improperly treated plates with

"soft spots" may have been incorporated into aerospace structures. A

major result of our research has been a delineation of which alloy

tempers and plate thicknesses are most likely to contain "soft spots"

due to specific processing errors.

The investigation included the following:

• Studies on as-received material. This included a 15.24 cm (6 in.)

thick plate of 2024-T851 , an 0.635 cm (1/4 in.) thick plate of

2024-F, and a direct chill cast ingot for 2024, all obtained from

i ndustry.

t Investigation of the phases and inclusions present in cast 2024

aluminum alloy with the aim of determining the degree of microand macro-segregation and identifying the inclusions present in

the as-cast ingots.

• Determination of a set of C-curves which can predict the mechanical

and NDE properties for any type of quench following the solution

heat treatment. Two tempers, T851 and T351 , were investigated.

• Transmission electron microscope studies of the stable and metastable phases present and an attempt to relate the observed microstructural changes to the measured changes in mechanical properties

and NDE measurements.

• A study of the ultrasonic wave propagation as a function of thermomechanical treatment of the alloy with the objective of establishing

a correlation between ultrasonic data and mechanical properties

and of providing an additional NDE method to improve the characterization of the material.

• A nondestructive evaluation of the age hardening sequence by means

of dynamic eddy-current conductivity measurements.

• Use of a heat flow model to calculate almost all conceivable heat

flow conditions anticipated during the quench of 2024 aluminum

alloy plate from the solution heat treatment temperature of 495 °C.

The calculated time-temperature data were then coupled to the

C-curves and the variations in properties across different thickness

plates for the "worst case" heat flow conditions were predicted.

Predictable macrosegregation was obtained in laboratory ingots

of 2024 aluminum alloy. It was found that macrosegregation of copper

and other alloying additions in direct chill cast ingots of 2024

aluminum alloy cannot be completely eliminated by chill face scalping

and subsequent thermomechanical treatment. Although good scalping

practice should maintain compositions within specified limits with no deterioration in mechanical properties, the macrosegregation remaining

in the finished plate product will contribute to the scatter observed

in NDE measurements. Further, because of the large copper content

variation near the chill face, surface hardness and eddy-current

conductivity measurements are necessarily very sensitive to scalping

depth in their ability to evaluate the condition of finished alloy

plates.

A large number of samples were taken from a 0.635 cm thick

plate of 2024 aluminum alloy in the F temper and processed to the

T851 , T351 or T4 temper. During processing to these tempers, the

quench following solution heat treatment was varied, giving a series

of samples with a wide range of microstructures and hence mechanical

properties. Two types of "pre-aging" treatment, labeled sequence A

and sequence B, were used. The hardness, eddy-current conductivity,

yield strength, ultimate tensile strength, elongation, and area reduction of these samples were measured. The accumulated data were

used to establish a set of approximate C-curves from which the alloy

properties can be established for any time-temperature cycle of the

quench following heat treatment. The C-curves can also be used to

generate correlations between mechanical and NDE properties. It was found that eddy-current conductivity alone cannot be used as a reliable

predictor of the mechanical properties of 2024-T351 or 2024-T851. It

must be combined with other information such as hardness and yield

strength measurements on the same lot (same ingot or plate) of material.

The C-curves were combined with time-temperature data from a computerized heat flow model to predict the variations in properties

across plates of different thicknesses for both sequence A and

sequence B type "pre-aging" heat treatments. It was found that the

T851 temper is quite "quench sensitive" in the sense that the

ultimate tensile strength falls below specification for the "worst

case" quench conditions for rather thin plates 1 cm thick).

TEM studies were carried out on a large number of specimens

in an 120 kV instrument equipped to operate in the scanning

transmission (STEM) mode as well as in the conventional transmission

mode. The instrument was equipped with an x-ray energy dispersive

spectrometer. A major objective was to identify those microstructural

changes that were responsible for the mechanical and physical properties
delineated in the C-curve representations. Although the microstructure

as a function of "pre-aging" treatment is complex, it appears that

"pre-aging" induced precipitates can contribute to the strength and are subject to overaging during further heat treatment. This provides a mechanism for the loss of strength which can account for the fact that

the strength of the T851 temper is actually reduced below that of the

T351 temper given the same "pre-aging" heat treatment (rather than

being raised as occurs for properly quenched material).

The objective of the ultrasonic studies was to determine the extent

to which a correlation exists between the mechanical and ultrasonic

properties. For this purpose, the absolute, rather than the relative,

values of the sound velocity and ultrasonic attenuation are required.

The absolute values of sound velocity and ultrasonic attenuation were determined to within ± 1 ms 1 and ± 0.02 dB, respectively. A parabolic

relationship was found between hardness and sound-wave velocity, whereas

ultrasonic attenuation decreases with increasing hardness.

IV
TABLE OF CONTENTS

Pa^e

I. INTRODUCTION 1

II. STUDIES ON AS-RECEIVED PLATES OF 2024 ALUMINUM ALLOY 4

III. SOLIDIFICATION-SEGREGATION STUDIES 7

1. Microsegregation in Case 2024 Aluminum Alloy 7

(a) Calculation of Microsegregation for

Al-Cu-Mg-Mn-Fe-Si Alloy System 7 (b) Second Phase Particles Formed During

Solidification of 2024 Aluminum Alloy 11

2. Macrosegregation in Cast 2024 Aluminum Alloy 14

(a) DC Cast Ingot 15

(b) Laboratory Cast Ingot 16

3. Thermomechanical Treatment and Evaluation of

Laboratory Ingot 19

IV. C-CURVES AND NONDESTRUCTIVE EVALUATION 21

1. Thermomechanical Treatment 22

2. Mechanical and Electrical Measurements 23

3. Calculation of C-curves 24

4. Discussion and Comparison of C-curves 29

5. Correlation Plots 31

6. Comparison with Other Data 32

V. TRANSMISSION ELECTRON MICROSCOPY STUDIES AND THE

RELATIONSHIP BETWEEN MICROSTRUCTURE AND PROPERTIES 34

1. Introduction 34

2. Experimental Procedure 36

3. Microstructure of 0.635 cm Thick Plate Specimens .... 37

(a) Constituent Phases 39

(b) T4, T351, and T851 Microstructures 44

(c) "Pre-aged" Microstructures 46

4. Discussion: Relationship Between Microstructure

and Properties 58

(a) T4 Condition 59

(b) T351 and T851 Conditions 66

VI. ULTRASONIC CHARACTERIZATION 68

1. Introduction 68

2. Experimental Procedure 70

3. Results and Discussion 72

VII. EDDY-CURRENT CONDUCTIVITY VERSUS HARDNESS DURING AGING. ... 79

1. Introduction 79

2. Experimental Procedure 81

3. Experimental Results 82

4. Discussion 85

v
VIII HEAT FLOW-PROPERTY PREDICTIONS 90

1. Heat Flow Model 90

IX. CONCLUSIONS - 94

1. As-received Plate 94

2. Solidification-segregation Studies 94

3. C-curves and Nondestructive Evaluation 95

4. TEM Studies: Relationship Between Microstructure

and Properties 96

5. Ultrasonic Characterization 98

6. Eddy-current Conductivity vs Hardness During Aging ... 98

7. Heat Flow-Property Predictions 98

REFERENCES 100

TABLES 103

FIGURES 128
I. INTRODUCTION

This is a comprehensive technical report of our investigations in

the past two years on 2024 aluminum alloy. The overall aim of this

work was to develop specific relationships between process variables

during casting, working and heat treatment of the alloy and the resulting

microstructures, mechanical properties and various nondestructive

evaluations (NDE), eddy-current (electrical) conductivity, hardness,

and ultrasonic measurements.

This work is a follow-up to a similar study on 2219 aluminum alloy

(1). Both investigations were carried out at the request of the National

Aeronautics and Space Administration. The initial motivation was

government and industry concerns that substrength aluminum alloys may

have been used in aircraft and space vehicle structures 1

. The concerns

originated from the discovery of "soft" spots 2 in an anodized 2124-T851

aluminum alloy machined part in July 1979. The "soft" spots were

apparently due to improper processing of the plate (1). Furthermore,

it was established that the same plant had produced a variety of other

aluminum plates including 2024 and 2219 aluminum alloys. Serious

concerns were also expressed about the variability of test techniques

used for quality assurance or finding the suspect plates 1.

In our earlier work (1), we determined the effect of improper

quenching on the mechanical properties and NDE (eddy-current conductivity,

and hardness) measurements of 2219 aluminum alloy. The kinetics of

precipitation for two types of cooling sequences from the solution heat

1 Aviation Week and Space Technology, August 1980, p. 14, and

August 13, 1980, p. 17.

2 "Soft" spots denote areas of a plate with mechanical properties

below federal specifications.

1
treatment temperature were investigated. Emphasis was placed on the

reliability of eddy-current conductivity and hardness as NDE tools to

detect variations in microstructure, hence in mechanical properties,

introduced by the various cooling sequences. An important result from

this work was that under the "worst case" quench conditions plates of

2219 aluminum alloy with thicknesses less than about 5 cm (2 in.) will

not suffer yield strength degradation below levels in federal specification QQ-A-250/30. Qualitatively speaking, the results showed that 2219

aluminum alloy is much less sensitive to improper heat treatment than

are other aluminum alloys. This is largely the result of the high

copper content in this alloy that gives a "C curve" with a "nose" at a

relatively high temperature.

Very evident from the results of this work (1) was the need to

always double check eddy-current conductivity measurements by periodic

hardness measurements or better, by direct mechanical property measurements when possible. Furthermore, a need also exists for better NDE

techniques which can scan aluminum plates rapidly and observe the

entire plate thickness rather than a thin surface layer. Because of

this, in the work described in the present report, the feasibility of

using ultrasonic techniques were also investigated as a corollary

method to available NDE tools for quality assurance.

The specific aims of the investigations carried out on 2024 aluminum

alloy included the following:

• To establish processing conditions and mechanisms responsible for

the occurrence of "soft" spots;

2
• To establish correlations between process variables during solidification and thermomechanical treatment and the composition and

microstructures of the plates;

• To identify the constituent phases in the alloy and to determine

the kinetics of precipitation and relate the microstructures to

mechanical properties and NDE measurements;

• To develop correlations between heat flow during quench from

solution heat treatment temperature and time-temperature precipitation models in order to determine the ranges of possible

degradation in mechanical properties due to improper heat treatment.

In this report we describe the details of our work on the 2024

aluminum alloy under the following subheadings:

• Studies on as-received plates of 2024 aluminum alloy;

• Solidification-segregation studies, microsegregation and

macrosegregation in laboratory and commercially cast ingots;

• C-curves and nondestructive evaluation, time-temperature

precipitation diagrams and the relationships between mechanical

properties and NDE measurements;

• Transmission electron microscopy studies, and the relationship

between microstructure and properties;

• Ultrasonic characterization;

• Eddy-current conductivity characterization, the study of aging

process by means of dynamic eddy current measurements, and

• Heat flow-property predictions, property degradations due to

improper quench from the solution heat treatment temperature.

3
II. STUDIES ON AS-RECEIVED PLATES OF 2024 ALUMINUM ALLOY

Two plates of 2024 aluminum alloy were used in this investigation.

They were:

1. A 15. 24 cm (6 in. ) thick plate of 2024 in the T851 temper

obtained from the Reynolds Metals Corp. Average composition

of this plate3 was determined to be 3.95 wt.% Cu, 1.47 wt.% Mg,

0.57 wt.% Mn, 0.20 wt.% Fe, 0.05 wt.% Si, and 0.025 wt.% Zn.

2. A 0.635 cm (1/4 in.) thick plate of 2024 in the F temper also

obtained from the Reynolds Metals Corp. The chemical composition

of this plate3 was determined to be 4.3 wt.% Cu, 1.44 wt.% Mg,

0.57 wt.% Mn, 0.28 wt.% Fe, 0.11 wt.% Si and 0.1 wt.% Zn.

In addition, segregation in a direct chill (DC) cast ingot of

2024, obtained from the Reynolds Metals McCook plant, was studied. The

average composition of this ingot was determined to be 4.54 wt.% Cu,

1.5 wt.% Mg, 0.52 wt.% Mn, 0.32 wt.% Fe, 0.11 wt.% Si and 0.1 wt.% Zn.

Note that there is some variation in the composition of the three

pieces of 2024 aluminum alloy used in this study. In particular, the

two 2024 plates appear to meet the composition specification for 2124

(with respect to the Fe and Si contents), whereas for the DC cast ingot

the Fe content exceeds the 2124 specification by a small amount. The

The composition of 2024 aluminum alloy according to ASTM Spec. B209

is 3.8 to 4.9 wt.% Cu, 1.2 to 1.8 wt.% Mg, 0.3 to 0.9 wt.% Mn,

0.50 wt.% Fe max, 0.50 wt.% Si max, 0.25 wt.% Zn max, 0.10 wt.% Cr

max, 0.15 wt.% Ti max (0.20 wt.% Ti + Zr max), each other 0.05 max

(total other 0.15 wt.% max). The composition of 2124 aluminum alloy

according to ASTM Spec. B209 is 3.8 to 4.9 wt.% Cu, 1.2 to 1.8 wt.% Mg,

0.3 to 0.9 wt.% Mn, 0.30 wt.% Fe max, 0.20 wt.% Si max, 0.25 wt.% Zn

max, 0.10 wt.% Cr max, 0.15 wt.% Ti max (0.20 wt.% Ti + Zr max), each

other 0.05 wt.% max (total other 0.15 wt.% max). The Cu, Fe, Mg, and

Mn contents of the 0.635 cm plate and the Cu contents of the 15.24 cm plate and the DC ingot were determined by atomic absorption spectrometry.

4
effect of these compositional differences is not determined herein and

would require a more extensive study. Because of iron's ability to tie

up large amounts of Cu and Mn under certain conditions, results could

be expected to be affected by the iron content.

It was expected that across the thickness variations in properties

due to macrosegregation in the original DC cast ingot or due to a

cooling rate gradient, because of normal thermal resistance of the

plate during quench from the solution temperature, would be most pronounced in very thick plates. Therefore, the 15.24 cm thick plate in

the T851 temper was carefully examined for chemical, microstructural

and property variation across its thickness.

The data obtained for the composition, hardness, and electrical

conductivity of the 15.24 cm thick plate are shown in Figure 1. These

data essentially establish the variations in properties due to the

normal variation in cooling rate experienced during quench from the

solution heat treatment temperature, and the macrosegregation remaining

in the plate from the original DC cast ingot. The first plot in

Figure 1 shows that there is approximately 0.2 wt.% variation in copper

content across the plate. This was determined by molecular absorption

spectrometry (wet chemistry). The abrupt changes in copper content at

the edges of the plate are due to the depleted region (negatively

segregated region) next to the chill face in the original DC cast

ingot4

. The scalping apparently removed the positive chill face segregation leaving some of the depleted region intact which ended up in the

plate. The negative segregation at the plate centerline is due to the

4 Macrosegregation across the DC cast ingot of 2024 aluminum alloy produced at the Reynolds McCook plant is shown in a subsequent section.

5
same type of segregation noted in the DC cast ingot4

. Rockwell B hardness measurements were made according to ASTM E-18 on a Wilson bench

model Rockwell hardness tester5

. Eddy-current conductivity measurements

were made using a Verimet M4900B conductivity meter. This conductivity

meter was checked before each use with three standards with nominal

conductivities of 30, 40, and 50 percent IACS (International Annealed

Copper Standard). These standards had in turn been calibrated at a

temperature of 23.0 ± 0.2 °C using the NBS conductivity bridge with the

standards. During use, the standards and the sample to be measured were

placed on a large aluminum block at room temperature (20 to 25 °C) and

allowed to equilibrate to the same temperature within better than ±0.5 °C.

All conductivity measurements herein refer to the value at 23 °C. Repeatability at any location on a given specimen was ± 0.1 percent IACS.

Although no detailed statistical analysis was made, spot checks indicate

that they were accurate to ± 0.5 percent IACS.

All references to commercial equipment in this report are for

identification purposes only and in no way constitute any

endorsement or evaluation of the relative merits of such

equi pment.

6
III. SOLIDIFICATION-SEGREGATION STUDIES

A series of calculations and experiments have been performed to

determine the phases present in cast 2024 aluminum alloy with the

ultimate aim of determining the degree of micro- and macrosegregation

and identifying the inclusions present in the as-cast ingots. Such

segregation could have effects on the heat treatment response of this

alloy and its properties, including nondestructive measurements made

for quality control.

1. Microsegregation in Cast 2024 Aluminum Alloy

(a) Calculation of Microsegregation for Al-Cu-Mg-Mn-Fe-Si Alloy

System

Calculation of expected microsegregation for the n-component alloy

was performed using the assumption of local equilibrium at the interface,

complete diffusion in the liquid phase, no diffusion in the solid phase

and no fluid flow in the interdendritic "mushy" region. During solidification of primary oraluminum, the situation is governed by ( n- 1

)

differential equations (2)

df

L

dC

Li

l-ka

l

i = 1, ... n-1 [ 1 ]

where f^ is the weight fraction liquid, C^. is the liquid concentration

of the i-th alloying element, and is the i-th equilibrium partition

coefficient for the solidification of the oraluminum phase. In general,

k. is a function of C, , , C 10 , ... C .

,

, x , but because the tie lines of

the phase diagram are not known we have assumed that k^ are constant

and are determined from the binary diagrams of aluminum with each

alloying addition.

7
Solution to equation [1] in this case is:

^ ^

Vk“-1

01

[ 2 ]

where C^. is the original composition of the i-th component and f is

the weight fraction solid (lrf^). We calculate the solidification

"path" (C|j vs f

s

) to determine at what fraction solid the i nterdendri ti

c

liquid becomes saturated with respect to a second solid phase, i.e.,

when the solidification "path" encounters a multivariant eutectic (or

peritectic). After this point, the solidification is governed by a

different set of differential equations.

Segregation studies in 2024 aluminum alloy requires the examination

of the Al-Cu-Mg-Mn-Fe-Si senary system. Other alloying additions are

present in small quantities or act as grain refiners and have been

neglected. Values of the used are given in Table I. As an example,

Table II shows the calculated solidification "path" (concentration of

the interdendritic liquid as a function of f

g

) for an alloy Al-4.0 wt.%

Cu-1.4 wt.% Mg-0.65 wt.% Mn-0.2 wt.% Fe-0.1 wt.% Si. The maximum levels

in 2024 for Fe and Si are 0.5 wt.%. The relatively low values for Si

and Fe used here are close to the ingots and plates characteri zed in

this study, and are more typical of a 2124 alloy.

Determination of the solidification "path" in the six component

phase diagram is extremely difficult and can only be done approximately.

In general, many possibilities exist, depending on the initial alloy composition, for the formation of second phases in the interdendritic region.

If we examine the solidification "path" (Table II) in various ternary

system combinations of these different components the possibilities become

8
apparent. Figures 2 and 3 show the solidification "paths" plotted in

the Al-Cu-Mg, and Al-Cu-Fe ternary systems.

In the Al-Cu-Mg system, the solidification "path" intersects the

monovariant eutectic through L -» orAl+CuMgA^ at f

g = 0.91 at about

513 °C. Hence, in this ternary system CuMgA^ is the second phase to

form and the third phase to form would be CuAl^ by way of the ternary

eutectic L orAT»-CuAl

2+ CuMgAl^ at 508 °C.

In the Al-Cu-Fe ternary system, the solidification "path" would

intersect the monovariant eutectic through L -» a-Al+(Fe,Cu)(Al ,Cu)g at

f

g = 0.83 at about 600 °C. Hence, in this case (Fe,Cu)(Al ,Cu)^ is the

second phase to form. The third phase to form would be Cu

2

FeAl^ by way

of the ternary peritectic L + (Fe,Cu)(Al ,Cu)g a-Al + Cu

2

FeAl^ at

590 °C. During dendritic solidification it is unlikely that this

ternary peritectic occurs to any extent and hence, (Fe ,Cu)(Al ,Cu)g

would probably remain in the microstructure. The interdendritic liquid

composition would continue on the monovariant eutectic through

L -» orAl+Cu

2

FeAl7 until complete solidification at the ternary eutectic

point L -» a-Al+Cu

2

FeAl^+CuAl

2

at 548 °C. It should be noted that a

slight reduction in the initial Fe content or small increases in the

equilibrium partition coefficient used for Fe will cause the solidification "path" to first intersect the monovariant eutectic L -» a-Al+C^FeAl^

and hence the phase (Fe ,Cu)(Al ,Cu)g will not form.

Examination of the Al-Fe-Si ternary system indicates that the

solidification "path" would intersect the binary eutectic through L

-*

a-Al+FeAl^. Between 629 and 611 °C, solidification would continue as

L -» a-Al+Fe^Si

2AI

^ 2

* This phase is often designated a-Al(Fe)Si or

orAl (Fe ,Mn)Si when Mn is present.

9
From these ternaries, we note the eutectics involving FeAl^,

a-Al(Fe)Si and (Fe ,Cu)(Al ,Cu)g occur at temperatures above 600 °C.

This fact makes one of these phases a likely candidate for the second

phase to solidify in the six component alloy. Such an analysis is

consistent with the microstructural observations of Sperry (3) on cast

2024 aluminum alloy.

An examination of the quaternary phase diagram data of Phragmen (4)

was performed to attempt to make the above predictions more specific.

Figure 4 shows a projection of the quaternary tetrahedron Al-Cu-Mg-Fe

into the Al-rich corner (2). The coordinates of this diagram are based

on the relative percentages of Mg, Fe, and Cu. In this figure a-aluminum

is always present and hence regions represent the solidification of two

solid phases and lines represent the solidification of three solid

phases (ternary eutectic, single arrow and ternary peritectic, double

arrow). Examination of the compositions reached in the interdendritic

liquid shows that (Fe ,Cu)(Al .Cu)^ is the second phase to form followed

by Cu^FeAl^. During the subsequent freezing of a-aluminum and (^FeAl^,

the liquid composition moves toward the line representing the ternary

eutectic L -» a-Al+C^FeAl-^+CuMgA^. Solidification will be completed

at the quaternary eutectic point D L -» a-Al+C^FeAl

^ + CuMgAl

2 +CuAl

In the Al-Cu-Mg-Si system shown in Figure 5, after the solidification of a-Al , freezing continues with the binary eutectic L -> a-Al+CuMgA^.

Subsequently, the liquid composition encounters the ternary eutectic

L -> a-Al+CuMgA^+CuA^- Solidification is completed at the quarternary

eutectic point C, L -> a-Al+CuMgAl^* CuAl

2 +Mg2

Si.

In the Al-Cu-Fe-Si system, the solidification "path", Figure 6, is

far removed from the phase Fe^S^Al-^ also designated a-Al(Fe)Si.

10
However, as will be seen in the next section, this phase is commonly

found in both cast and solution heat treated 2024 aluminum alloy.

(b) Second Phase Particles Formed During Solidification of 2024

Aluminum Alloy

Two ingots of 2024 aluminum alloy were examined by optical, SEM

and TEM microscopy to determine the phases present in the as-cast

condition. The first ingot was laboratory cast from alloy taken from

the 15.24 cm plate. It was solidified in a graphite crucible with a

cooling rate of approximately 0.15 K/s. The second ingot was a section

of the DC (direct chill) cast 2024 aluminum alloy obtained from the

Reynolds McCook plant.

The determination of the phases present in as-cast 2024 aluminum

alloy and related multi-component aluminum based alloy systems has been

the subject of numerous studies. The most relevant work to this study

is that of Sperry (3) and more recently of Munitz et al. (5) which deal

specifically with phases present in as-cast 2024 aluminum alloy.

Sperry (3) in particular has identified phases in their order of appearance during solidification by means of optical and x-ray methods. The

reported composition of Sperry's alloy was 4.90 wt.% Cu, 1.40 wt.% Mg,

0.70 wt.% Mn, 0.30 wt.% Fe, and 0.15 wt.% Si which compares closely

with the compositions studied here. Sperry (3) investigated cooling

rates in the range 0.06 to 0.92 K/s. A list of the phases he observed

together with their crystal structures is given in Table III. An

important observation made by Sperry concerned the fact that the major

Fe bearing phase at the slower cooling rates was Cu^FeAl^ while at the

higher cooling rates cubic orAl (Fe,Mn)Si was the predominant Fe bearing

phase. The phases present in the slow cooled laboratory and DC cast

11
ingots were determined by optical and scanning electron microscopy

using the known etching response and morphologies of the phases. TEM

(transmission electron microscopy) was used to confirm the identification

of phases by means of electron diffraction and EDS (x-ray energy dispersive spectrometry) and to study submicrometer sized structures.

Laboratory Cast Ingot . As previously noted, the average cooling

rate during solidification of this ingot was ~ 0.15 K/s. The measured

segregate spacings, i.e., secondary dendrite arm spacings, were in the

range between 70 and 90 pm. Optical and SEM examinations indicated the

presence of small second phase particles within the primary a-Al dendrites

concentrated around the border. A TEM micrograph6 of these particles

is shown in Figure 7. They are readily identified as the S' phase 7 by

means of electron diffraction. These precipitates develop through a

solid state reaction during cooling. The higher concentration of

precipitates at the dendrite borders is a consequence of microsegregation.

The interdendritic regions in this ingot consisted typically of a

fine coupled eutectic structure along with other coarser phases.

Examples are shown in Figures 8 and 9. Three phases are present in the

coupled eutectic and are identified as a-Al, O-CuA^ and S-CuMgAl^ in

Figure 8. A tentative identification was first made by EDS and then

confirmed by electron diffraction. In addition to this ternary eutectic,

a blade shaped phase was also frequently observed, Figure 9. The shape

immediately suggests that it is Cu^FeAl^ and, indeed, this was found to

be correct by means of electron diffraction and EDS. The three secondary

6 Details describing TEM specimen preparation are given in Section V.

7 The S' phase has approximately the same composition and structure

as S-CuMgAl

?

. A more detailed discussion of this phase is presented

in the description of "pre-aging" microstructures in Section V.

12
phases, G-CuAl^, S-CuMgA^ and Cu^FeAl^ together with the primary a-Al

phase were the major constituents in this relatively slow cooled ingot.

Observation of these phases is in accord with Sperry's (3) results. In

particular, the presence of Cu^FeAl^ is consistent with Sperry's finding

that this phase is promoted by a slow cooling rate. Other phases

listed in Table III may have been present but with a much lower frequency.

An exhaustive search for them was not made.

DC Cast Ingot . Samples for detailed microstructural characterization were taken from the one-quarter position across the short transverse

direction. The measured secondary dendrite arm spacings in this location

were in the range between 30 and 50 pm. This indicates a substantially

higher average cooling rate during solidification than that measured in

the laboratory cast ingot. Primary a-Al dendrites in the DC cast ingot

contained S' precipitates as was found in the laboratory ingot. In

contrast to the laboratory ingot, the precipitates were present at a

higher density and were distributed more uniformly throughout the

dendrites rather than being located primarily at the borders. An

example of the S' precipitates in an a-Al dendrite is shown in Figure 10.

The interdendritic structures in the DC ingot shown in Figure 11

were predominantly of two types: a fine multi-phase eutectic structure

and a single phase constituent sometimes with a polyhedral shape that

often entirely filled the interdendritic gap. An example of the eutectic

structure is shown in Figure 12. Four phases were present. They were

identified as a-Al, G-CuAl^, S-CuMgA^ and Mg^Si. According to the

analysis of the solidification "path" in this report and also in Sperry

(3), this comprises the quarternary eutectic which is expected to be

the last structure to solidify. In many instances, the Mg

2

Si phase was

13
not found, possibly because it was present at a lower concentration

than the other phases. The eutectic would then be identified as the

ternary structure which was found in the more slowly cooled laboratory

ingot. Occasionally, a blade of the C^FeAl^ phase was also noted in

the structure, Figure 11(b).

An example of the polyhedral ly shaped single phase interdendritic

structure is shown in Figure 13(a). Apparently this phase forms a

divorced eutectic with a-Al. Electron diffraction patterns from the

phase could be indexed in a way that was consistent with a bcc unit

cell having a lattice parameter of a

Q

1.25 nm. On this basis and with

composition information obtained by EDS, Figure 13(b), the phase was

identified as cubic a-Al(Fe,M)Si where M in this case designates Cu and

Mn. With the exception of a somewhat lower Si concentration in the

present samples, the composition agrees closely with results obtained

by Munitz et al . (5). The presence of a-Al(Fe,M)Si and not Cu^FeAl^

agrees with Sperry's (3) finding that the former phase is favored by

higher cooling rates. None of the other phases listed in Table III

appeared to be present to a significant extent at this location in the

DC cast ingot.

2. Macrosegregation in Cast 2024 Aluminum Alloy

A series of experiments and measurements were performed to determine

the degree of macrosegregation in semi-continuous DC cast and laboratory

cast ingots of 2024 aluminum alloy. The former gives an indication of

the maximum composition variations expected in the final plate product

while the latter will be used in the preparation of control specimens

to establish the effect of composition on properties including the NDE

measurements.

14
(a) DC Cast Ingot

A sketch of the section of DC cast 2024 alumimum alloy ingot

used in this study is shown in Figure 14. It extends from chill face

to chill face of the casting in the short transverse direction and from

chill face to centerline in the long transverse direction. Measurement

of macrosegregation was performed initially in a direction perpendicular

to the chill face in the short transverse direction (position Bl) shown

in Figure 14. Chemical analysis was performed by the Center for

Analytical Chemistry at NBS by spectrophotometric measurement at 435 nm

of a copper complex. The chemistry data should be considered accurate

to ± 0.04 wt.%.

As seen in Figure 15, very high positive copper segregation

is noted at the chill face 12 wt.% Cu) followed by a region of

rapidly varying composition that extends almost 3 cm into the ingot.

Two minimums in composition 4.25 wt.% Cu) occur at ^ 0.2 cm and 2 cm

from the chill face. A relatively uniform composition region 4.6 wt.%

Cu) extends from ^ 3 cm to 15 cm from the chill face followed by a

region of negative segregation at the ingot centerline where the composition drops to ~ 4. 1 wt.% Cu. These variations are caused by interdendritic fluid flow during solidification (6). The inset in Figure 15

shows the drastic reduction in the volume fraction of eutectic which

occurs about 0.2 cm from the chill face. This reduction is the origin

of the rapid drop in copper content near the chill face. Figures 16

and 17 show optical and SEM views respectively, of the relative change

in the volume fraction of the eutectic present at various distances

from the chill face. These photos can be qualitatively correlated to

the measured copper contents at these positions shown in Figure 15.

15
The occurrence of the two minimums in copper content near the

chill face was considered to be rather unusual. In the previous study

of macrosegregation in DC cast 2219 aluminum alloy (1) only a single

minimum in copper content was observed near the chill face. Although

for the 2024 ingot these composition excursions failed to take the

ingot outside of the specification for copper, the second minimum at

2 cm is substantially further from the chill face than is the comparable

minimum for copper (0.8 cm) observed in the previous studies of 2219

aluminum alloy. To establish whether this behavior was an exception to

normal in this ingot, further macrosegregation profiles were performed

at different positions in the ingot section shown in Figure 14. Data

near the chill face for sections A1 and A2 along with B1 are shown in

Figure 18. Note that the magnitude of the minima have changed but

their existence is quite evident. Figure 19 shows data in the long

transverse direction labeled LT in Figure 14. Here it might be argued

that the first minimum is not significant, however, the second minimum

is present and exists at the 2 cm position.

One can only speculate about the origins of this unusual

macrosegregation profile. Depressions in average copper content are

caused by sudden air gap formations and subsequent reheating of the

ingot surface during solidification. For example, it may be postulated

that after the initial air gap formation, adequate cooling was

re-established and interrupted for a second time. It is expected,

however, that in most ingots a single "dip" in copper content near the

chill face will be observed.

(b) Laboratory Cast Ingot

Following the work of Mehrabian and Flemings on macrosegregation

in multicomponent systems (6), a special geometry unidirectional ingot

16
was cast with the 2024 aluminum alloy to demonstrate the mechanisms

responsible for the positive and the negative macrosegregation noted in

the DC cast ingot and to obtain controlled composition samples differing

from the nominal. These samples were used for thermomechanical treatments

and nondestructive evaluation.

The geometry of the casting is shown in the inset of Figure

20 and employs a reduction in area of approximately 10 to 1 to cause

macrosegregation. The bottom section of the casting is 9.2 cm square

and 9.0 cm high while the top section of the casting is 2.7 cm square

and 21 cm high. An investment mold of plaster was preheated to 500 °C

and placed directly on a water cooled chill block. The mold has an

open bottom so that molten metal comes into direct contact with the

chill. This, coupled with the preheated mold, guaranteed directional

solidification of the ingot. The mold was filled with molten 2024

aluminum alloy (obtained from the full cross section of the DC Ingot)

at about 720 °C after being degassed with chlorine. The casting was

analyzed for average composition variation (macrosegregation) in the

direction perpendicular to the chill. The variation in copper and

magnesium content determined by atomic absorption spectrometry as a

function of distance from the bottom chill in the as-cast ingot is

shown in Figure 20. As expected from previous studies of macrosegregation noted above, high positive segregation (5.35 wt.% Cu) is observed

at the chill face- while negative segregation (^ 3.1 wt.% Cu) occurs in

the region of the cross section change. To a lesser extent, magnesium

also shows a high level at the chill face, drops to a minimum near the

cross section change and rises again in the rest of the ingot. This

similarity of shapes of the composition profiles is an indication that

17
the equilibrium partition coefficients for Cu and Mg (0.17, and 0.30

respectively) are less than one. Because the partition coefficient for

Mg is larger, macrosegregation of Mg is less than Cu. The relative

levels of Cu and Mg are important because they are the major components

in the hardening phase, CuMgAl^, in heat treated 2024 aluminum alloy.

Chill face and cross-section-change segregation both result

from the flow of segregated interdendritic liquid to feed solidification

shrinkage. It should be noted that the laboratory ingot exhibits a

lower positive chill face segregation and no adjacent negative segregation compared to the DC cast ingot. The occurrence of these phenomena

near the chill face of the DC cast ingot can be readily ascribed to the

formation of an extensive air gap which results in the abrupt reheating

of the ingot surface during solidification and the exudation of the

solute rich interdendritic liquid from the adjacent region. On the

other hand, the negative segregation at the DC ingot center line and at

the section reduction of the laboratory ingot are due to the extensive

flow of interdendritic liquid from the hotter to the cooler regions of

the ingots at these locations.

A second laboratory ingot of 2024 aluminum alloy with a 10 to

1 reduction in cross section was also cast in a manner identical with

the first ingot with eight thermocouples inserted through the mold at

different distances from the bottom chill face. Figure 21 shows the

temperature-time curves for the eight thermocouples and the distance

from the chill face to each. The liquidus and solidus temperatures for

this alloy are approximately 638 (± 5) and 502 (± 5) °C respectively and

hence, from these curves the local solidification time as a function of

position has been determined. Additionally, secondary dendrite arm

18
spacings were obtained from metal 1 ographic sections at various distances

from the chill face. These spacings are correlated with the local

solidification time in Figure 22.

3. Thermomechanical Treatment and Evaluation of Laboratory Ingot

The effect of variations in alloy composition due to macrosegregation

on the heat treatment response of 2024 aluminum alloy and its properties

including nondestructive measurements used for quality assurance has

been evaluated using samples taken from the reduced-cross-section

laboratory ingot.

Samples, ~ 2.7 cm square and ^ 2 cm thick in the solidification

direction, were cut from the ingot at different distances from the

bottom chill face to obtain specimens with differing compositions. The

samples, of course, contained composition gradients through their

thickness, but were relatively uniform in perpendicular directions due

to the unidirectional solidification. As a reference point, eddy-current

measurements were conducted on as-cast samples. Measurements were made

with the coil on the surfaces which were perpendicular to the solidification direction. The conductivity along with the previously determined

copper content is shown in Figure 23. These data indicate that wide

variations in the conductivity can exist depending on composition and

structure unrelated to subsequent thermomechanical treatment.

The samples were then thermomechanical ly treated to the T851

condition as follows:

(a) homogenization heat treated for 48 hours at 495 °C,

(b) hot rolled at 495 °C to 1/6 of their initial thickness (0.32

cm, ^ 1/8"),

(c) solution heat treated at 495 °C for 75 minutes,

(d) stretched 2.5% and aged at 190 °C for 12 hours.

19
Hardness measurements were carried out on both surfaces of these

0.3 cm thick samples. In Figure 24 hardness is plotted as a function

of the original distance of these surfaces from the bottom chill face

along with the copper content.

The hardness levels obtained in the bottom section of the casting

are relatively insensitive to changes in Cu and Mg content between 4.4

to 5.3 wt.% Cu and 1.3 to 1.5 wt.% Mg. The hardness level of 77 HRB is

typical of correctly processed commercially cast 2024 aluminum alloy in

the T851 condition. Near the cross section change of the ingot, the

hardness drops to below 64 HRB, a value normally considered below

specification. The composition at this point in the casting was 3.1

wt.% Cu, and 1.1 wt.% Mg, values which are below the minimum specification for these alloying additions (3.8 wt.% Cu and 1.2 wt.% Mg). After

the cross section change, the hardness rises to ~ 78 HRB as the compositions rise to normal levels. The low hardnesses obtained at the very

top of the ingot are due to excess porosity which occurred there due to

shrinkage. Although the limited number of samples obtained preclude

precise conclusions, the ranges of compositions seen in the as-received

plate (3.85 to 4.00 wt.% Cu) should not be the major cause of the

hardness variation across that plate.

20
IV C-CURVES AND NONDESTRUCTIVE EVALUATION

A large number of samples taken from the 0.635 cm (1/4 in.) thick

plate of 2024 aluminum alloy in the F temper condition were processed

to the T851 and T351 tempers. During the processing to these tempers,

the quench following solution heat treatment was varied. This gave a

series of samples with a wide range of microstructures and hence

mechanical properties. The hardness, eddy-current conductivity, yield

strength, ultimate tensile strength, elongation, and area reduction of

these samples were measured. The data accumulated were used to establish

a set of approximate C-curves from which the alloy properties can be

estimated for any time-temperature cycle of the quench following solution

heat treatment. The data are also presented in a number of correlation

plots to show the relationship between the mechanical properties and

the nondestructive evaluation (NDE) parameters. The measured changes

in physical and mechanical properties are related to microstructural

observations in the next section.

Some comparisons are also made with data from other sources and

with our previous study of 2219-T87* aluminum alloy (1). In particular,

it appears that 2024 aluminum alloy is a more "complex" system than the

2219 aluminum alloy. This complexity has a number of consequences. In

particular, it is seen that the simple C-curve shapes used in the

analysis, which worked well for the case of 2219-T87* aluminum alloy,

give much more approximate results for the case of 2024-T851 and

2024-T351 aluminum alloys. Further, relationships between mechanical

properties and NDE measurements (i.e. hardness and eddy-current conductivity) have considerably more scatter, making the task of testing the

2024 aluminum alloys more difficult. This is especially true for the

21
case of the naturally aged T351 material, and should be taken into

account in the specifications for this material.

1 . Thermomechanical Treatment

Specimens 0.635 cm thick and measuring 2.5 cm by 18 cm were cut

from the plate, supplied by the Reynolds Aluminum Company, such that

the long dimension was aligned in the rolling direction. The plate was

received in an F temper condition. As previously noted, the chemical

composition of the plate represents approximately the middle of the

specified range for Cu, Mg and Mn contents for both 2024 and 2124

aluminum alloys (3.8 to 4.9 wt.% Cu, 1.2 to 1.8 wt.% Mg, and 0.3 to 0.9

wt.% Mn). The amount of Si and Fe are also within specification of

both alloys (0.5 max. wt.% Fe and 0.5 max. wt.% Si for 2024 aluminum

alloy; 0.3 max. wt.% Fe and 0.2 max. wt.% Si for 2124 aluminum alloy)

with the Fe content being, however, only slightly below the maximum

specified for the 2124 aluminum alloy. Since the presence of iron is

expected to have a deleterious effect on the yield and ultimate tensile

strengths of these alloys, the results obtained here might be considered

applicable in the sense of a lower bound for 2124 aluminum alloys with

mid-range Cu, Mg, and Mn contents and with similar other elemental

contents. The effect of compositional variations on the C-curves was

not investigated and the degree to which such variations would alter

the C-curve parameters and contribute to interlaboratory disagreement

on the effect of heat treatment on alloy properties is not fully known

at present.

The 0.635 by 2.5 by 18 cm specimens were subjected to a systematic

thermomechanical processing schedule involving a 75 minute solution

heat treatment at 495 °C followed by a specially designed "pre-aging"

heat treatment using a salt bath. The term "pre-aging" heat treatment

22
is being used in this report in place of "interrupted quench" or "slack

quench" because the definition of the latter two terms does not appear

to be fully established. We use "pre-aging sequence A" or simply "sequence

A" to refer to a monotonic quench from solution heat treatment and

"pre-aging sequence B" or simply "sequence B" to refer to a quench with

reheating to an intermediate temperature.

During the "pre-aging" treatment, thermal data were continuously

acquired and manipulated by an automated high speed data acquisition

system for later use in C-curve determination. Following the solution

heat treatment at 495 °C for 75 minutes and "pre-aging", the specimens

were stretched from 2.25% to 2.5% permanent strain and either artifically

aged at 190 °C for 12 hours (T851) or aged at room temperature (T351).

A schematic of this thermomechanical treatment is given in Figure 25.

The alloys thus produced consist of four sets: 2024-T351 sequence A;

2024-T351 sequence B; 2024-T851 sequence A; and 2024-T851 sequence B.

2. Mechanical and Electrical Measurements

Following final aging, the hardness and eddy-current conductivity

were measured for each sample. Approximately one-half the samples were

machined into tensile specimens of the shape shown in Figure 26. The

yield strength (0.2% offset), ultimate tensile strength, elongation,

and reduction in area, as well as hardness and eddy-current conductivity,

were determined for each of these machined specimens.

Hardness measurements were made on the Rockwell B scale according

to ASTM E- 1 8 on a Wilson bench model Rockwell hardness tester. Each

time hardness was measured, two measurements were taken on the sample

surface at a random location, excluding 5 cm from the sample ends where

the sample was gripped during the stretching operation and, for the

23
tensile test specimens, excluding the gage. The yield strength (0.2%

offset) and ultimate tensile strength were determined on a calibrated

Satec System Inc. Baldwin Model 60 CG Universal Testing System. Eddycurrent conductivity measurements were made using the instrument and

procedure described in the previous section.

Hardness and conductivity measurements were made both before and

after machining the samples into tensile test specimens. Approximately

0.5 mm was milled from each sample surface during the machining operation.

Comparison of these measurements gives an indication of any possible

effect of a machining operation on these measurements, as well as an

indication of the measurement scatter that might be expected due to the

combined influences of measurement precision and "normal" sample variations. The variations obtained for hardness and conductivity were

± 0.57% IACS (two sigma level) and ±3.7 HRB (two sigma level), respectively. This comparison is illustrated in Figures 27 and 28.

The time- temperature quench history and the measured properties

for the thermomechanical ly processed samples are summarized in Table IV

for both tempers and both "pre-aging" sequences. These data are used

below to calculate C-curves and compare the mechanical properties with

the hardness and electrical conductivity.

3. Calculation of C-curves

The data on hardness, eddy-current conductivity, yield strength,

ultimate tensile strength and time-temperature history have been used

to determine a set of C-curves for both sequence A and sequence B

alloys and for tempers T851 and T351.

The C-curves are a family of C-shaped curves used to characterize

the effects of quench rate on the final properties of the finished

material. Their use for aluminum alloys was pioneered by W. L. Fink

24
and L. A. Willey (7). These curves represent the effect of "pre-aging"

time at a given temperature (intermediate between the solution heat

treatment temperature and the quench water temperature) on the final

properties of the material. In using these curves to determine the

effect of an arbitrary time- temperature history during the quench from

solution heat treatment temperature, a "rule of additivity" developed

by Cahn (8) is used. Basically, this rule states that the degradation

in, for example, yield strength that occurs in a given temperature

interval is independent of that which occurred during previous temperature intervals passed through during the quench. The total degradation

is taken as the simple sum of the degradations occurring in each temperature interval during the quench. This rule is also used in determination

of the C-curve but is much less important there because the "pre-aging"

is made as nearly isothermal as possible.

The extent to which the "rule of additivity" is applicable to

aluminum alloys has not been established. However, previous experience

by Staley (9) indicates that it can be applied with good approximation

at least to some aluminum alloys. One might expect the rule to apply

approximately when the degradation in properties is small and when the

temperature decreases monotonical ly. The differences we find in the

present study between the sequence A and sequence B heat treatments

gives some indication as to how closely this rule is followed. As in

the previous study (1), we use the results from these two sequences in

section VIII to calculate limits on what properties might be obtained

under "best" and "worst" quenching conditions in 2024-T351 and 2024-T851

aluminum alloy plates.

25
Following Evancho and Staley (10) and Cahn (8,11), we assume that

the resulting value of a property, a, can be represented as

° = (om

' °o ) ex» ( - KlxV + °o [3]

where a

m

is the maximum achievable property, a

Q

is the minimum or

"intrinsic" value of the property achieved under given conditions, and

i is given by

i x

dt

Cx (T(t))

[4]

with t being the time at which the quench from solution heat treatment

temperature is started, t

Q

the time to achieve a temperature less than

about 120 °C, T(t) the temperature as a function of time, and C

x (T) the

C-curve. The C-curve is represented by the equation

C (T) = K, K0 exp

x

v J lx 2

v [5]

where 1^, K^, K^, and are constants to be determined, T is the

absolute temperature, and K^ x

, is a constant given the value

[ 6 ]

This value of K, is chosen so that for x > 1, a < a . Thus x becomes

a critical parameter

IX for achieving some given

XXX

value, a , of the property

in question. The curve C x (T) has the dimensions of time and, when

plotted as a function of temperature (normally with the temperature as

the vertical axis and C x (T) as the horizontal axis), represents the

26
isothermal holding time (or critical time) at any temperature needed to

degrade the property to the value of ct 3 , with ct > cr > a.

x mxo

To determine the parameters o

m

, ct q, K^, Kg, K^, and K^ the following

procedure was used:

(i) The temperature from 110 °C to 496.5 °C (1.5 °C above the solution

heat treatment temperature) was divided into 32 temperature

i nterval s

,

(ii) The time, t. , spent in each of the temperature intervals during

the "pre-aging" treatment was stored in a datalogger,

(iii) These time data were used to numerically calculate the integral

of Equation [4] according to

ti

C (T.)

x

v

i y

[7]

where T. is the average temperature in the interval, and

(iv) Using an iterative, non-linear fitting routine, values of the

parameters which minimize the least squares deviation between

measured and calculated values were obtained.

The computer program used in the least squares fitting has been

listed and described in a previous report (1). For each iteration in

the fitting procedure, an "estimated standard deviation", defined by

e. s. d. [ 8 ]

t h where ct. is the measured value of the property in question for the i

sample, ct^ is the calculated value of that property, and N is the total

number of samples, is determined. The program allows one or more parameters to be fixed while the remainder vary to establish a constrained

27
minimum value of the estimated standard deviation. This is useful

because the values of K^, K^, and can often be varied over a considerable range without significant changes in the estimated standard deviation. A "significant change" in the e.s.d. was considered to be a

change greater than 10 percent of the "normal scatter" as determined

previously for the hardness and conductivity. It was found that the

value of was indeterminate over a range of about 20,000 to 40,000

cal/mol in the sense that, for any value of in this range, the other

parameters could be adjusted to give an essentially equivalent value of

the estimated standard deviation. A value of 32,000 cal/mol was selected

for this parameter on the grounds that it is close to the known activation energy for the diffusion of copper in aluminum. Having chosen K^,

values of and were selected for each temper and "pre-aging"

sequence in a similar manner. Given these values of K^, K^, and K^,

the remaining parameters were varied for each property to arrive at a

final set of parameters with the constraint that the a

m of each property

must be identical for the A and B sequences of each temper. The final

set of parameters is listed in Table V. For the T351 alloys, samples

"pre-aged" below 250 °C were not included in the fit. "Pre-aging"

below 250 °C is in some ways equivalent to an artificial aging and, for

T351 , increased the values of yield strength and hardness above those

obtained for a direct quench.

The C-curves obtained from these parameters are displayed in

Figures 29-44. With each C-curve is a plot of the measured property

versus the value of that property calculated from the C-curve parameters

using the rule of additivity and the data of Table IV. In these plots,

the data have been divided into six groups: (1) those alloys given a

direct quench (no "pre-aging"); (2) those alloys "pre-aged" at an average

28
temperature between 200 and 250 °C; (3) between 250 and 300 °C; (4) between

300 and 350 °C; (5) between 350 and 400 °C; and (6) between 400 and 450 °C.

These groups are given the symbols, respectively, of asterisk, plus sign, x,

circle, square, and triangle.

4. Discussion and Comparison of C-curves

The C-curves presented in Figures 29-44 represent the isothermal

holding time for a property to degrade to a given value. The most

critical temperature, i.e. the temperature at which the most rapid

degradation occurs, is represented by the minimum or "nose" of the

C-curves. The parameterization of Equations 3-6 gives a curve of

approximately the correct shape for a single precipitation process.

However, as explained in more detail in section V, there are several

precipitation processes, with each process probably requiring a separate

and different C-curve.

The justification for using a single C-curve to describe what in

reality is a much more complicated process lies in the fact that this

empirical approach appears to give a fairly good approximation of the

effect of "pre-aging" on the properties of the material after it has

been fully processed. The estimated standard deviations for the hardnesses, see Table V, are only slightly greater than the single sample

deviation of 1.85 HRB (one-sigma level) determined on the machined

samples as discussed above. Much more data would be required to

sensibly utilize a model with two or more individual C-curves (and

hence a larger number of adjustable parameters). Inspection of the

measured vs calculated plots of Figures 29-44 shows the predictive

value of the single C-curve approach fails most noticeably in two

areas. First, the large degradation that occurs for long "pre-aging"

29
times around 350 °C and above are not suitably predicted. In this

case, the C-curves have been adjusted to slightly overpredict the

degradation that does occur. Second, for the T351 alloys, "pre-aging"

in the temperature range of about 200 to 250 °C actually increases the

yield strength, ultimate tensile strength and hardness, rather than

decreasing these properties as predicted by the C-curves. This is because

"pre-aging" these alloys, which later receive only a natural aging, in

this temperature range accomplishes some artificial aging which enhances

the hardness of the final product. Thus, in fitting the model C-curves

for the T351 alloys, data on "pre-aging" in the temperature range of 200

to 275 °C was ignored. For both cases, prediction of the mechanical

properties which result from a particular quench following solution

heat treatment made using the C-curves presented here should tend to

slightly underestimate these properties. Further, the predicted values

are expected to have their greatest validity at the upper end of the

property value range, i.e. between a

m and about (am

+ ct q )/2.

Figure 45 gives a comparison of the yield strength C-curves of

2024-T851 , 2024-T351 and 2219-T878 aluminum alloys for both sequence A

and sequence B "pre-aging" treatments. The curves are drawn to represent

isothermal times required for a reduction to 90% of the maximum attainable

yield strength. During the quench from solution heat treatment, degradation of properties occurs most rapidly at temperatures near the "nose"

of the C-curve. During normal quenching, the time spent reaching a

given temperature increases with decreasing temperatures. Hence, for

two C-curves identical except for the nose temperature, the one with

the nose at higher temperature will be less "quench sensitive".

8 This is a modified T87 treatment practiced at the Reynolds McCook

plant (1).

30
Comparing the A and B sequences of 2024-T351 aluminum alloy it is

seen that the B sequence nose is at a higher temperature. The B sequence

is, however, slightly shifted towards shorter times. This gives the A

and B sequences of T351 nearly the same yield strength quench sensitivities (compare the section below on heat flow property predictions).

Comparing the 2024-T851 aluminum alloy sequence A and B C-curves it is

seen that, at all temperatures, the B sequence C-curve lies at shorter

times than the A sequence curve and hence the B sequence has greater

quench sensitivity. Comparing the 2024-T851 and 2024-T351 C-curves it

is seen that in almost all cases T851 is more quench sensitive. Further,

comparing 2219-T87* with 2024 it can be seen that, unless large amounts

of time are spent above 400°C during the quench, 2219-T87* aluminum

alloy is generally less "quench sensitive" than either 2024-T851 or

2024-T351 aluminum alloy.

5. Correlation Plots

Plots showing the correlation of ultimate tensile strength and

yield strength with the NDE properties of hardness and eddy-current

conductivity are displayed in Figures 46-57. The solid and dotted

lines in each figure, denoting sequence A and sequence B, respectively,

were calculated from their appropriate C-curves. It is seen that the

correlation predicted by the C-curves fits the data well. Inspection

of the scatter bands (which were drawn independently of the C-curves

using only the data itself) shows the difficulty in using conductivity

as an independent NDE tool. The scatter bands are very broad, especially

for the T351 alloys.

In Figures 53 and 55, which plot ultimate tensile strength and

yield strength, respectively, vs hardness for 2024-T351 aluminum alloy,

it can be seen that many yield strengths and hardnesses lie above the

31
terminus of the C-curve correlation. These represent samples which

were "pre-aged" at temperatures around 250 °C and below. The scatter

in the yield strength vs hardness for T351 (Figure 55) is very large.

Much of this scatter represents a real difference between the sequence

A and sequence B specimens, with the sequence B specimens generally

lying higher in yield strength and having considerably more scatter.

This difference can be related to the difference in precipitation

mechanism for the naturally aged 2024-T4 alloy as discussed in section

V. For example, the 0 phase found in sequence A specimens was not

found in sequence B specimens of 2024-T4.

6. Comparison With Other Data

A large amount of data on the ultimate tensile strength, yield

strength, hardness, and eddy-current conductivity has been presented by

Petrak and Gunderson (12). These data were taken from specimens cut

from a (5-1/2 in.) thick plate of 2124-T851, from a (2-3/4 in.) thick

plate of 2124-T851, and from a (2 in.) thick plate of 2024-T351 which

had been found to have or were suspected of having improperly quenched

regions. Petrak and Gunderson (12) plotted the ultimate tensile and

yield strengths vs hardnesses and conductivities for both the 2124-T851

and 2024-T351 data. For each plot, they have least squares fitted the

data to provide a mean line and a lower 90% confidence line for each

data set. In Figures 58-65 we have superimposed their lines (heavy

solid and dashed, representing the least squares mean and the lower 90%

confidence, respectively) on our own data.

Inspection of Figures 58-61 shows generally good agreement between

the data of this report and the data of Petrak and Gunderson (12) for

strength vs hardness of 2024-T851. There is also some agreement, over

32
at least a portion of the conductivity range, in the strength vs conductivity of 2024-T851. For 2024-T351 (Figures 62-65) the agreement is

less satisfactory. In every case, the strength vs hardness plots of

Petrak and Gunderson (12) lie lower than our own data, with a truly

large difference in the yield strength vs hardness data for 2024-T351

(Figure 64). This large difference is difficult to rationalize.

Considering the much closer agreement on the T851 temper, it could not

be simply a measuring instrument calibration problem. It may be due to

a difference in composition of the materials represented. For example,

as previously mentioned, the Fe and Si content of the 2024 alloy we

used met the specifications for 2124. However, no compositional analysis

is given by Petrak and Gunderson (12). Further investigation of the

reasons for these differences would undoubtedly be useful.

33
V. TRANSMISSION ELECTRON MICROSCOPY STUDIES AND THE

RELATIONSHIP BETWEEN MICROSTRUCTURE AND PROPERTIES

1 . Introduction

Changes in physical and mechanical properties of 2024 aluminum

alloy brought about by an improper quench or a "pre-aging" treatment

subsequent to solution heat treatment are related to the development of

stable and metastable phases. These phases either do not form or do so

to a negligible extent during a "proper" quench from the solution heat

treatment temperature. The main objective of the TEM (transmission

electron microscopy) investigation was to reveal the nature of the

"pre-aging" induced phases. The results of that investigation will be

presented in this section. As a natural corollary to the TEM investigation, an attempt has been made to relate the observed microstructural

changes to the measured changes in physical and mechanical properties.

Part of this discussion, particularly as it relates to mechanical

properties, will be included in the present section. Some additional

considerations will be presented in the following sections describing

NDE evaluations.

TEM studies were carried out on a large number of specimens from

the 0.635 cm (1/4 in.) thick plate. Some of these specimens were

"pre-aged" and processed to T351 and T851 conditions as described in

Section IV. The majority of the specimens, however, were prepared

specifically for the TEM investigation. These specimens were subjected

to "pre-aging" treatments which in many cases were identical to those

given specimens prepared for physical and mechanical property measurements,

but subsequent processing to T351 and T851 conditions was not carried

out. In this way, microstructural changes associated with stretching

34
and final aging operations which might change or confuse the observation

of "pre-aging" induced phases was avoided.

Strictly speaking, the microstructure was not necessarily identical

to that immediately after "pre-aging". The 2024 aluminum alloy, especially when quenched directly into ice water from the solution treatment

temperature or when "pre-aging" does not result in a significant reduction

in solute supersaturation, is subject to fairly rapid aging at room

temperature. Therefore, examination was carried out after room temperature aging to an essentially stable condition, i.e. after what may be

referred to as a T4 temper treatment. The so-called natural aging

response in this alloy is believed to be a direct result of the formation

of G.P. zones, or G.P.B. zones as they are frequently referred to in

Al-Mg-Cu alloys (13). In principle, valuable information with respect

to the level of solute supersaturation and the effect of solute segregation could be obtained by studying G.P.B. zones, however, as will be

pointed out later, zones exhibiting the expected characteristics were

not imaged in the TEM. In any case, the presence of G.P.B. zones has

no effect on the observation in the TEM of the various precipitates

formed during thermal processing. Indirect evidence of G.P.B. zones

and, more importantly, the concentration of available solute can be

obtained from room temperature aging experiments. Such experiments were

carried out on many of the specimens used for TEM examination. The

results are presented in this section in connection with the discussion

of the relationship between microstructure and properties.

After a brief description of the experimental procedure, the major

portion of this section will be concerned with the identification in

the TEM of various precipitates and constituent phases found in the

35
"pre-aged" specimens. A discussion of the influence of "pre-aging"

induced precipitation on mechanical and physical properties will be

presented at the end of this section.

2. Experimental Procedure

Specimens for TEM examination were prepared by cutting with a low

speed diamond saw a thin slice approximately 0.25 mm thick from bulk

material. Disks 3 mm in diameter were then punched from the slice

and electropol ished with a dual jet polishing device to obtain electron

transparent areas. The electropolishing solution consisted of 70%

methyl alcohol and 30% concentrated nitric acid. Satisfactory results

were obtained when the solution was maintained in the temperature range

18 to 20 °C and the applied cell voltage was approximately 20 V. The

presence of relatively large constituent phase particles in all materials

and porosity in the as-cast materials described earlier had an adverse

effect on the quality of the foils obtained. Preferential attack at

these features often led to foils that were thicker than was desirable

so that several attempts were sometimes required before a satisfactory

specimen was obtained. For the as-cast ingot, where the principal

objective was the identification of constituent phases, preferential

chemical attack and the possibility of a systematic loss of a particular

phase species was most undesirable. Therefore, some specimens were

thinned by argon ion milling. With this method, it is often possible

to achieve the simultaneous thinning of specimen areas which differ

widely in composition and crystal structure. Ion milling was also

carried out on some specimens subsequent to electropolishing to remove

or minimize the presence of an electrochemically induced corrosion film

or otherwise chemically depleted or enriched layer (14) that might

36
interfere with the accuracy of energy dispersive x-ray analyses in the

electron microscope.

Most of the TEM studies were conducted at 120 kV in an instrument

equipped to operate in the scanning transmission (STEM) mode as well as

in the conventional transmission mode. The instrument was also equipped

with an x-ray energy dispersive spectrometer (EDS). With this system

it is, in principle, possible to determine the chemical composition of

a column of material as small as 10 nm in diameter traversed by the

electron beam. The analysis is limited to elements with atomic numbers

^ 11 (Na). In general, an accurate analysis requires a knowledge of

the foil thickness, lack of an interfering layer of material having a

different composition and an appropriate calibration of the EDS system.

In this investigation, the analyses were primarily qualitative or

semi -quantitative in nature for the purpose of identifying the phases

present. Identification of a particular phase generally involved the

use of both electron diffraction and elemental composition determination

by EDS.

3. Microstructure of 0.635 cm Thick Plate Specimens

A chart showing the "pre-aged" specimens that were studied by

means of TEM is presented in Figure 66. The "pre-aging" times and

temperatures indicated are nominal values in each case. Because of the

finite heating and cooling rates, some of the "pre-aging" time interval

was spent at temperatures other than those shown. Typical "pre-aging"

temperature vs time curves are shown in Figure 26. A departure from

the ideal "isothermal" cycle would probably have the most significant

effect on the microstructures observed in specimens given short "preaging" treatments, particularly 20 s, where a considerable fraction of

37
the time shown was spent at higher than indicated temperatures in the

case of sequence A specimens and at lower temperatures for the sequence

B specimens. Even with the longer "pre-aging" treatments any time

spent at a temperature which might be critical to a nucleation process

has the potential of producing a microstructure that is different from

that which would occur for a strictly isothermal treatment at the

designated target temperature. The specimens in Figure 66 were selected

to provide a survey of the microstructural changes that occurred over

# the range of "pre-aging" treatments covered in this investigation. A

major objective was to identify those microstructural changes that were

responsible for the mechanical and physical properties delineated in

the C-curve representations. In addition to the specimens shown in

Figure 66, several others were examined. These are listed in Table VI.

They include a specimen quenched directly into ice water from the

solution treatment temperature and aged at room temperature to provide

a reference microstructure against which "pre-aged" specimens could be

compared. Directly quenched specimens processed to the T351 and T851

tempers were also included to provide examples of "properly" quenched

and processed microstructures with those tempers. The remaining two

specimens included in Table VI were "pre-aged" and then processed to

the T851 temper.

The TEM specimens in each case were prepared from thin slices cut

perpendicular to the longitudinal plate direction. The specimen plane

was therefore perpendicular to the longitudinal direction. Since it

was usually necessary to tilt the specimen in the electron microscope

to obtain the desired crystallographic orientation with respect to the

electron beam, the normal to the electron micrographs shown in this

38
report may deviate from the longitudinal direction by up to 30°. The

region examined was always approximately 1 . 6 mm in from the plate

surface. Variations in the number densities, both relative and absolute,

of different precipitate types were also observed on the scale of a few

tenths of micrometers in a given specimen. This was taken as evidence

for composition fluctuations on the same scale, indicating that homogenization of the inti ally segregated ingot structure had not been completely

achieved. In the following two subsections, the nature of the constituent

phases and of the "pre-aging" induced precipitates in the plate specimens

will be described.

(a) Constituent Phases

Constituent phases found in wrought 2024 aluminum alloy

materials can generally be classified into two size ranges, large

particles having a major dimension > 1 pm in length and small particles

whose major dimension is less than 1 pm. The large particles are in

general from the as-cast ingot structures. They may be broken up and

redistributed during mechanical stages of processing and perhaps suffer

a phase change or diffusion related compositional modification during

thermal processing steps following casting. In contrast, the majority

of the small constituent particles are not present in the as-cast

structure and probably develop during thermomechanical processing as a

result of dissolution of interdendritic phases and subsequent precipitation. The small particles, so formed, are stable at the solution

treatment temperature and are usually referred to as dispersoids. They

play a major role in reducing grain growth during processing but also

have a critical effect on the nucleation of other phases during quenching

and aging treatments as will be discussed later.

39
The permissible concentrations of the various elements in

2024 aluminum alloy cover a sufficiently wide range that a rather large

variety of constituent phases are possible. Many of the interdendritic

phases listed by Sperry (3) (see Table III) will be retained even if

complete homogenization and equilibration are achieved at the solution

treatment temperature, 495 °C. This is especially true with respect to

phases containing Fe and Si because of the low solubilities of these

elements. Whether or not ©-CuA^ or S-CuMgA^ are present as constituent

phases will depend on the available concentrations of Cu and Mg at the

solution treatment temperature. With reference to the phase diagram

shown in Figure 67 it can be seen that if the concentration of Cu

exceeds approximately 4.0 wt.% and Mg is < 1.5 wt.% the excess Cu will

be taken up by the 0 phase. As the concentration of Mg is raised above

1.5 wt.% an increasing amount of the excess Cu will appear in the S

phase. The level of available Cu is particularly sensitive to the

concentrations of Fe and Mn, mainly because of the formation of ^FeAl^

and Mn^^A^Q and to some extent because of the solubility of Cu in

the phases a-Al(Fe,M)Si and MnAl^ (3,5).

During the course of examining a great many specimens listed

in Figure 66 and Table VI, large constituent phase particles were

frequently encountered. Most often, these were found to be the orAl(Fe,M)Si

phase. It may be recalled that this was also determined to be the

predominant Fe bearing phase in the DC cast ingot. An example of one

of these particles is shown in Figure 68(a) together with its diffraction

pattern in Figure 68(b). EDS analyses of several particles gave an

average composition of about 70% A1 , 12% Fe, 5% Mn, 5% Cu, and 4% Si

(atomic percent) similar to that obtained for this phase in the DC cast

40
ingot. There was no evidence to indicate decomposition of the a-Al(Fe,M)Si

phase. Sperry (3) reported that a-Al(Fe,M)Si and Cu^FeAl^ might serve

as a rough index to the degree of homogenization. Since there was no

evidence of decomposition of or-Al (Fe,M)Si and the phase, C^FeAl^, was not

found, it is apparent that other factors must be important. While large

constituent particles of the 0 phase were occasionally seen, particles

of the S phase were not found in directly quenched specimens studied in

the TEM. As will be described later, these phases were present after

certain "pre-aging" experiments. It should be emphasized, of course,

that due to the relatively small volume studied in the TEM, low concentrations of widely dispersed particles might not have been detected.

Dispersoid particles were observed in all the 0.635 cm thick

plate specimens examined. They were typically equi-axed in shape with

faceted sides, although rod shaped particles were frequently present.

The equi-axed particles ranged in size from approximately 30 to 80 nm;

however, much larger particles about 500 nm in diameter were often

observed at grain boundaries. Particle spacing was of the order of 10

nm. Typical distributions of dispersoid particles from a directly

quenched specimen are shown in Figure 69. In Figure 69(a) at low

magnification it can be seen that the particles are arranged in bands

with particle free regions between bands. This effect is almost certainly

associated with segregation in the original as-cast ingot. Particle

free zones were also present at some grain boundaries. No reference

was found in the literature giving a detailed identification of the

dispersoid phase. The fact that Mn does contribute to the formation of

dispersoid particles has been mentioned by several investigators in

connection with corrosion and mechanical properties studies of 2024

41
aluminum alloy (15-21). That Mn can be implicated in the formation of

dispersoid particles is demonstrated by the investigations of Robinson

and Hunter (18) and Stolz and Pelloux (19). Both examined Al-Cu-Mg

alloys with no measurable Mn or at a concentration of Mn below the

solubility limit and found that dispersoids were not present. Robinson

and Hunter mention the C^Mn^l^Q phase but do not indicate that the

dispersoid particles are that phase. Kaufman and Low (20) referred to

rod shaped precipitates 0.1 to 0.2 pm in size at grain boundaries in

2124 aluminum alloy as (^MngA^Q but did not indicate how the identification of that phase was made. Van Stone et al. (21) indicate that

dispersoid particles in 2000 series alloys have been identified as

Al

i

2 (FeMn)3

Si . Petri (22) and others (23,24) have determined that the

solid solubility of Mn in Al-Cu-Mn alloys is approximately 0.2 wt.% at

500 °C. At a composition of Al-4wt.% Cu-0.6wt.% Mn according to Petri's

phase diagram, which is confirmed by Day and Phillips (23) and Hofmann

(25), the two phases a-Al and (^Mn^A^g (referred to as "T" by Petri)

are present. The work of Phragman (26) indicates that the additional

presence of 1 . 5 wt.% Mg in 2024 aluminum alloy probably does not change

this result. It might also be noted that Robinson and Hunter (18) show

qualitatively similar electron micrographs of dispersoid particles in

Al-Cu-Mn and Al-Cu-Mg-Mn alloys.

With respect to the material studied here, an average Mn

concentration of 0.6 wt.% was determined by chemical analyses. However,

some of the available Mn (0.4 wt.%) is lost through the formation of

the insoluble a-Al(Fe,M)Si phase formed during solidification. It is

therefore reasonable to assume that a large fraction of the dispersoid

particles in 2024 aluminum alloy are the Cu

2

Mn

2Al

2 Q phase. Analyses of

42
the Cu^Mn^A^Q phase have determined that it may differ from the ideal

composition, 80 at.% A1 , 8 at.% Cu, and 12 at.% Mn. Petri obtained a

variable composition ranging from 80.1 at.% A1 , 7.8 at.% Cu, and 12.1

at.% Mn to 78.1 at.% A1 , 6.9 at.% Cu, and 15.0 at.% Mn. Day and Philips

(23) found a single composition, 11 at.% Cu and 15 at.% Mn.

EDS analyses were carried out on a number of dispersoid

particles in the directly quenched specimen as well as in several

"pre-aged" specimens. A composition of about 80 to 90 wt.% A1 and 5 to

10 wt.% each of Cu and Mn was obtained in each case. The uncertainty

in this analysis was due to the different thicknesses of the particles

examined and to the fact that some thickness of the ot-Al matrix probably

covered many of the particles. Electron diffraction studies of the

particles were also carried out. Petri determined that C^Mn^A^Q has

an orthorhombic unit cell. A detailed x-ray crystallographic study by

o o Robinson (27) gave unit cell dimensions of a

Q = 24. 2A, bQ = 12. 5A and

o

c q = 7.72A, quite consistent with results obtained by Petri. Robinson

also determined on the basis of systematically absent reflections that

the space group was either Bbmm, Bbm2 or Bb2m. Several characteristics

of the dispersoid particles and of the C^Mn^A^Q crystal structure

complicate its identification by means of electron diffraction. They

are the small particle size, twinning within individual particles, a

large unit cell size and the fact that most of the strongly reflecting

planes have very nearly the same spacings (27). These factors may

explain why previous investigators have not offered a specific identification of this phase. It was possible in this investigation, however,

to obtain microdiffraction patterns from individual particles and to

index these patterns consistent with the (^Mn^A^Q structure. An

example is shown in Figure 70.

43
(b) T4, T351 and T851 Mi crostructures

It is widely accepted that age hardening at room temperature

of supersaturated Al-Cu-Mg alloys having a Cu:Mg ratio of approximately

2.2:1 occurs by the formation of so called G.P.B. zones (13) which

presumably are clusters rich in Cu and Mg. Age hardening at room

temperature of the more complicated 2024 aluminum alloy is expected to

occur by the same means. Silcock (13) on the basis of x-ray measurements

has proposed that the G.P.B. zones are rods parallel to the <100>

o matrix direction with an estimated length of 40 A and diameter of 10 to

o 20 A. Although TEM has proved to be very valuable for the direct observation of zones in a number of systems (Al-Cu is an especially good

example (28)), no reference was found providing a detailed study of

G.P.B. zones formed at room temperature in Al-Cu-Mg alloys. In fact,

it was noted in several cases that G.P.B. zones could not be resolved

(29-31). G.P. zones in general can be revealed in the TEM by diffraction

contrast as a result of surrounding strain fields (coherency strains)

in the matrix, in dark field due to non-matrix diffraction contributions,

by structure factor contrast, or through lattice imaging (32). During

the normal course of examining a number of specimens in this investigation, conclusive evidence of G.P.B. zones, consistent with the description

of Silcock (13), was not obtained despite a marked room temperature age

hardening response. This was true both with respect to images obtained

under a variety of diffraction conditions, and particularly in diffraction

patterns. There was no obvious contribution in addition to the a-Al

matrix pattern to suggest the presence of zones which one might expect

on the basis of the x-ray diffraction results. It must be concluded

that zones in Al-Cu-Mg alloys differ significantly from those in the

44
binary Al-Cu system, especially with respect to the development of

coherency strains.

The microstructure of a directly quenched specimen in the T4

condition shown in Figure 69 has already been discussed. An additional

micrograph of this specimen is shown in Figure 71 with dislocations in

contrast. The dislocations were either retained after solution treatment

or were introduced during the quench. Migration of excess vacancies to

dislocations and subsequent climb of screw segments has led to the

formation of helices (33). As will be seen later, these dislocations

can act as sites for the heterogeneous nucleation of precipitates.

The T351 temper is similar to T4 with the exception of an

intermediate plastic deformation step after quenching in which the

material is stretched to a strain of ~ 2.5%. The resulting microstructure

is shown in Figure 72. The relatively high dislocation density is a

consequence of the stretching operation, otherwise, the microstructure

does not appear to differ from the T4 condition.

The effect of artifical aging of 2024 aluminum alloys and

similar Al-Cu-Mg alloys has been extensively studied (34,35). Three

different stages of precipitation are recognized. They are generally

represented by the scheme,

supersaturated a-Al -> G.P.B. zones S' -> S.

It does not follow however that G.P.B. zones necessarily preceed the

formation of S' or that S is preceeded by S'.

The S' phase is thought to be nearly identical to S in composio tion (CuMgAip and structure (orthorhombic with a

Q = 4.04 A, bQ = 9.25

o o A, and c

q = 7.18A) (36). The two phases differ mainly in that S' is

45
semi-coherent and S is incoherent. Both S' and S have the same orientation relationship with the matrix:

[100]

s [100]a

; [010]

s [021]a

; [001]

s [012]o

.

Individual S' precipitates grow in the form of laths on { 21 0}^ planes

with their long axis parallel to <001

>

a

; thus there are 12 different

orientation variants in the a-Al matrix. Adjacent S' precipitates are

often arranged to form corrigated sheets.

The effect of aging time and temperature on the development

of the various precipitate species has been studied by Silcock (13) in

the ternary alloy Al-3.15 wt.% Cu-1.52 wt.% Mg which closely corresponds

to 2024 aluminum alloy. At temperatures between 130 and 260 °C, G.P.B.

zones were the first species to be observed. After approximately 10

days at 130 °C and in much shorter times at higher temperatures, S'

precipitates were also detected. At 260 °C only the S' phase was

found.

The T851 temper is achieved by exposing T351 material to

artificial aging at 195 °C for 12 hrs. The resulting precipitates are

predominantly S'. Grain boundary precipitates form also and these are

considered to be the S phase on account of their non-lath like shape.

An example of the T851 microstructure is shown in Figure 73(a). The

diffraction pattern from this area is shown in Figure 73(b) and the

corresponding indexed pattern is shown schematically in Figure 73(c).

That the majority of the S' precipitates probably nucleate heterogeneously

at dispersoid particles and dislocations will be demonstrated later.

(c) "Pre-aged" Microstructures

During quenching from solution treatment it has been found

that a cooling rate of less than approximately 550 °C/s through a

46
critical temperature range extending from about 400 °C to 290 °C is

detrimental to many of the properties of 2024 aluminum alloy (37).

Corrosion is one of the most sensitive properties to quench rate, and

degradation in corrosion behavior is the first effect to be noted when

the rate is reduced below 550 °C/s. Mechanical properties are not

affected until the quench rate is reduced below about 10 °C/s--corresponding

to a dwell time of 8 s in the critical temperature range. Note that

the nose of the yield strength C-curve, Figure 30(a), occurs at about

8 s. Materials in the naturally aged condition are in general more

susceptible to corrosion than those subjected to artificial aging. The

major microstructural consequence of a reduced quench rate is to permit

the nucleation and growth of precipitates that not only do not improve

the mechanical properties but consume solute elements that otherwise

would contribute to the desired age-hardening G.P.B. zones and precipitates. Furthermore precipitates formed during a slow quench may produce

sites for enhanced corrosion activity.

As already described two different "pre-aging" treatments

were employed to develop C-curves which can serve to predict the effect

of an improper quench on properties. Because two separate processes

are involved, one being nucleation and the other growth, a single

"pre-aging" sequence is not adequate to explore behavior under conditions

where an improper quench treatment may result in a non-monotonic cooling

curve. Thus the sequence A "pre-aging" treatment consists of rapidly

cooling from the solution temperature to an intermediate "pre-aging"

temperature where the specimen is held for a prescribed length of time

before quenching into ice water. Nucleation and growth is then confined

to temperatures above or at the "pre-aging" temperature. In the sequence

47
B, "pre-aging" treatment, the specimen is first quenched to ice water

and then rapidly reheated to the "pre-aging" temperature. Here nucleation and growth occur mainly below or at the "pre-aging" temperature.

In what follows, the results of TEM studies of sequence A "pre-aged"

specimens will be described first followed by the presentation of

sequence B results.

Sequence A Microstructures . Representative examples of sequence A

microstructures from specimens "pre-aged" at 400 °C, 350 °C, and 300 °C

are shown in Figure 74 at a relatively low magnification and in Figure

75 at a higher magnification. In each case, structures corresponding

to "pre-aging" times of 20 s and 500 s are included. Specimens were

examined after 100 s and the microstructures were found to be intermediate

between those at 20 s and 500 s. Consider first the series at 400 °C.

After 20 s of "pre-aging", Figures 74(a) and 75(a), the only detectable

change from the directly quenched structure was the presence of slightly

larger, and perhaps more frequent, grain boundary precipitates. "Preaging" for 100 s and 500 s led to a substantial change. Not only were

the grain boundary precipitates very large but large precipitates were

also present in the grain interiors. An electron diffraction analysis

of the grain boundary precipitates indicated that they were S-CuMgA^-

An example of one of these precipitates with its electron diffraction

pattern is shown in Figure 76. This and other S phase grain boundary

precipitates examined exhibited the reported orientation relationship,

[100]

s [100]

a

; [010]

s [021

]

a

; [001

]

g [012]^,

with at least one adjacent grain at the boundary.

48
Precipitates present within grain interiors were of two types.

One type had an irregular but more-or-less equi-axed shape. The other

was disk shaped. An example of an irregular shaped precipitate is

shown in Figure 77. From its diffraction pattern it was determined to

be the S phase obeying the orientation relationship noted above. That

is, these precipitates were essentially identical to the grain boundary

precipitates.

The disk shaped precipitates were determined by electron diffraction

to be the 0 phase. Analysis of diffraction patterns indicated the

orientation relationship,

(i io)0 dii)

a

; [!l0]e

[011]

a

.

Furthermore, the face of the precipitate disks was (110)

Q

and was

parallel to (111)^. An example of a 0 precipitate with important

matrix and precipitate crystallographic directions indicated is shown

in Figure 78(a). The corresponding diffraction pattern is presented

in Figure 78(b). It is interesting that the orientation relationship

determined here was not one of the several relationships that have been

reported for the 0 phase in Al-Cu alloys. The latter relationships

have been summarized by Lorimer (36). All have a common feature in

that [001

]

Q

is either parallel to [001 (within ±4° in one case) or

parallel to [110] . The 0 phase observed in these "pre-aged" specimens

was oriented with [001

]

Q

parallel to [21 1]^.

It might be hypothesized that the presence of both 0 and S phases

after "pre-aging" at 400 °C is a manifestation of the approach to

equilibrium. The equilibrium conditions under which both 0 and S may

coexist in a ternary Al-Cu-Mg alloy comparable to 2024 aluminum alloy

49
are indicated in the phase diagram shown in Figure 67. On cooling, the

path indicated in Figure 67 is followed and is approximately coincident

with a 2.2:1 ratio of Cu to Mg. If Cu is available in excess of this

ratio the 0 phase will form until the excess Cu is consumed and thereafter

the alloy will behave as a pseudobinary of a + S. Observing both the 0

and S phases after "pre-aging" suggests there may in fact have been an

overbalance of Cu. To test this hypothesis, a specimen was "pre-aged"

at 400 °C for 7200 min. to allow time for equilibrium to be reached.

Examination of this specimen revealed only the S phase. It may be

concluded, therefore, that the presence of the 0 phase at short "preaging" times was not the result of the suggested equilibrium reaction

and that the 0 phase is metastable.

The series of specimens "pre-aged" at 350 °C were similar in many

respects to those exposed at 400 °C. For short "pre-aging" times, 0

was the major precipitate phase with increasing amounts of the S phase

appearing as the duration of "pre-aging" was increased. Some $' precipitates were also observed. The distinction between S and S' is based on

the shape of the precipitates. Thin rod or lath shaped precipitates

parallel to <100> directions are considered to be S' while the much

thicker precipitates without this characteristic shape are referred to

as S. A relatively high concentration of precipitates was already

present after 20 s as is shown in Figure 74(c). Most of the precipitates

in Figure 74(c) are 0.

The S phase precipitates here and in other examples can often be

recognized by the region of light contrast that surrounds them. The

light contrast is the result of a local reduction in thickness caused

by preferential attack during electro-chemical thinning in preparation

50
of the specimen. Figure 75(c) clearly demonstrates that the 0 precipitates are heterogeneously nucleated at dispersoid particles. Here the

precipitates have a distinct hexagonal shape with straight sides parallel

to <112>

a

- An example is shown in Figure 79 where the (111)^ plane and

a 0 precipitate in that plane are approximately parallel to the viewing

plane. The relationship between the S phase and the other precipitates

was not obvious although it appeared that they developed in association

with the dispersoid particles and perhaps the 0 phase precipitates.

As Cu is consumed by the 0 phase, the relative proportion of Mg increases

and, according to the phase diagram, the tendency for the formation of

the S phase should be enhanced. Nucleation of the S' precipitates was

at dislocations and possibly at dispersoid particles. After "pre-aging"

for 500 s, 0 was no longer the predominant precipitate phase. The

majority of the precipitates were found to be the S phase. "Pre-aging"

at 350 °C for 1000 min led almost entirely to the loss of the 0 phase

with only S and possibly S' being present. Therefore, it must be

concluded that the 0 phase is not an equilibrium phase at 350 °C.

"Pre-aging" at 300 °C also led to the formation of 0 phase precipitates at dispersoid particles. The dispersoid particles were also

identified as the sites for the formation of S

1 precipitates. After

"pre-aging" for 20 s, Figure 75(e), the precipitates have grown out

only a short distance from the dispersoid particles. When the duration

of "pre-aging" was extended to 500 s the precipitates were distributed

uniformly throughout the matrix, Figure 75(f). Prolonged "pre-aging"

for 1000 min as at 400 °C and 350 °C again led to the disappearance of

the 0 phase with only S and S' remaining. As a further observation it

was noted that S' was increasingly present at dispersoid particles as

the "pre-aging" temperature was reduced.

51
A single specimen "pre-aged" at 250 °C was studied. The duration

of "pre-aging" was 400 s. Both 0 and S' precipitates were observed at

dispersoid particles, similar to specimens "pre-aged" at 300 °C. In

addition, however, much smaller S' precipitates were distributed homogeneously in the matrix, with precipitate free zones in the vicinity of the

larger S' and 6 precipitates at dispersoid particles. These different

types of precipitates are shown in Figure 80. The observation of

homogeneously distributed precipitates is important here because it is

an indication that the "pre-aging" temperature, 250 °C, is in the

temperature range where G.P.B. zones may form. The assumption is made

that G.P.B. zones must preceed the development of homogeneously distributed S' precipitates. Beton and Rollason (38) have determined the

G.P.B. zone solvus for Al-Cu-Mg alloys by means of a reversion method.

At a composition corresponding to the 2024 aluminum alloy, they obtained

a dissolution temperature of about 260 °C for G.P.B. zones which tends

to support the observation made here.

The precipitation processes resulting from the sequence A "pre-aging"

treatments can be conveniently summarized by means of a nucleation or

transformation diagram. Such a diagram has been drawn in Figure 81.

Each curve represents the estimated time and temperature dependences

for the onset of a particular transformation. The representation shown

in Figure 81 should be regarded as largely schematic, since in no case

was an exact determination made to fix the position of a curve. The

first phase observed at elevated temperatures was S at grain boundaries.

In time, this was followed by 6 and then S'. Both 0 and S' nucleated

heterogeneously at dispersoid particles, and in addition, S' nucleated

at dislocations. One might speculate that in the absence of dispersoid

52
particles and dislocations neither of these phases would have formed.

At low temperatures, G.P.B. zones, possibly preceeded by spi nodal

decomposition, occur first and then transform to homogeneously distributed

S' precipitates. For longer "pre-aging" times the 0 phase disappears

to be replaced by equilibrium S. The precise nature of that transformation is not understood. Prolonged aging also results in the disappearance

of S' and an accompanying increase in the amount of S. Observations

also indicate that the homogeneous S' precipitates disappear in the

presence of larger heterogeneously nucleated S' precipitates. Whether

the latter transformation is simply a consequence of Ostwald ripening

or represents a more fundamental difference between homogeneous S' and

heterogeneous S' is not known.

The C-curves presented in Section IV can, of course, be related to

the nucleation curves. If a single time and temperature dependent

transformation process is involved and the property (hardness, electrical

conductivity, etc.) for which the C-curve is drawn is sufficiently

sensitive to that transformation process, then the first C-curve,

indicating an initial change in the property, should coincide with the

nucleation curve. When several interdependent processes are involved,

the determination of C-curves is essentially a curve fitting exercise.

If the processes are well separated in time and temperature, a single

family of C-curves will probably fit the data poorly. If the processes

overlap and are interdependent as is apparently the case here, a better

fit may be achieved.

Sequence B Microstructures . The microstructures of specimens

subjected to sequence B "pre-aging" treatments at 400 °C, 350 °C, and

300 °C for 20 s and 500 s are shown in Figures 82 and 83. In each

53
case, S' was the only precipitate phase that formed within the grains.

Precipitates of the S phase were observed at grain boundaries under all

conditions, much as was found in the sequence A specimens. The change

in precipitate size and concentration can be followed in Figures 82 and

83 as a function of "pre-aging" time and temperature. In general, the

size of the precipitates increases with increasing time at each temperature. After only 20 s of "pre-aging" at 400 °C, Figures 82(a) and

83(a), precipitates are uniformly distributed throughout the matrix.

Here and in all examples shown in Figures 82 and 83 the direction

normal to the micrograph is within a few degrees of [001 ] . Therefore,

eight orientation variants of the rod shaped S' precipitates are present

having their long axis parallel to the plane of the micrograph and four

variants are present with their long axis perpendicular to the plane of

the micrograph. When viewed end-on, it was evident that many of the

rods were nearly circular in cross-sectional shape. Other rods were

elongated with an aspect ratio of about 2:1 or less. At 400 °C, the

diameter of the rods determined in the end-on orientation is approximately

10 to 30 nm after 20 s and 20 to 90 nm after 500 s of "pre-aging".

Thus, the growth rate of the average precipitate varies approximately

as d a t where d is the precipitate diameter and t is the time. This

suggests that the process is diffusion controlled. Undoubtedly, factors

other than volume diffusion were involved in determining the growth

rate; for example, impingement of precipitates has certainly occurred

in many instances in Figures 82 and 83. Although the precipitate

concentrations were not measured, it was apparent that fewer precipitates

were present at 500 s than at 20 s, indicating that coarsening had

occurred. As the precipitates become larger and fewer in number with

54
longer "pre-aging" times due to coarsening, the obvious relationship

between precipitates and dispersoid particles is lost. For example,

the S' precipitates after "pre-aging" at 400 °C for 500 s, Figure

83(b), appear to be relatively independent of dispersoid particles.

"Pre-aging" at 350 °C led to a precipitate morphology which appears

similar to that at 400 °C. S phase precipitates are formed at grain

boundaries and S' within grains. Typical microstructures are shown in

Figures 82(c), 82(d), 83(c) and 83(d). The precipitates formed at

350 °C are smaller than those at 400 °C for the same "pre-aging" times.

After "pre-aging" for 20 s and 500 s, the end-on diameters of the S'

precipitates are 3 to 10 nm and 20 to 75 nm, respectively. Comparing

Figures 83(c) and 83(d) at 20 s and 500 s, respectively, the number

density of precipitates doesn't appear to be significantly different,

although the precipitates are considerably larger at 500 s. The increase

in size is probably due to growth from solution rather than coarsening

due to the dissolution of smaller precipitates as observed at 400 °C.

The heterogeneous nature of the precipitation process is very

clearly demonstrated by the specimen "pre-aged" for 20 s at 350 °C.

The rod shaped S' precipitates all seem to radiate out from each dispersoid particle along <200> directions. Precipitates clustered along

short dislocation lines can also be seen.

The trend in the precipitation process observed at 350 °C and

400 °C was also found at 300 °C. After "pre-aging" for 20 s the precipitates at dispersoid particles are very small, Figure 83(e). With 500 s

of "pre-aging" the precipitates have grown substantially, Figure 83(f).

In addition to precipitates at dispersoid particles, some very small

homogeneously distributed S' precipitates are present in the matrix.

55
These as well as precipitates at dispersoid particles and dislocations

are shown in the dark field image in Figure 84. The presence of homogeneously distributed precipitates at 300 °C seems not to be consistent with

the reaction G.P.B. zones -* S' since 300 °C is probably above the

G.P.B. zone solvus. However, the reaction may have taken place during

the temperature rise to 300 °C from the ice water quench. To support

this suggestion, it may be noted that homogeneously distributed S'

precipitates were not observed after the sequence A "pre-aging" treatment

at 300 °C.

One specimen was "pre-aged" at 250 °C for 400 s. This temperature

is below the G.P.B. solvus boundary and homogeneously distributed S'

precipitates were present in the matrix as might be expected. Much

larger S' precipitates were observed at dispersoid particles and dislocations. The microstructure was similar to that observed in the specimen

"pre-aged" at 300 °C for 20 s, Figures 83(e) and 84.

Specimens were also examined after prolonged "pre-aging" treatments.

At 400 °C after 7200 min large S phase precipitates were present along

grain boundaries. A few very large precipitates of either the S or S'

phase were observed within grains. Because of their rod- like shape

with their long axis parallel to [100]^, they are assumed to be S'.

The corresponding sequence A specimen exhibited a higher concentration

of S precipitates without the rod-like shape. Optical micrographs of

these two specimens are shown for comparison in Figure 85.

Nucleation curves can be drawn for reactions occurring during the

sequence B treatment as was done for the sequence A treatment. The

representation, shown in Figure 86, is much simpler than for the sequence

A case since the 0 phase was not detected. In common with the sequence

56
A treatment, the first phase to be observed after "pre-aging" at elevated

temperatures was the S phase at grain boundaries. Almost simultaneously,

however, the S' phase was found at dispersoid particles and dislocations.

Continued "pre-aging" led to the transformation, S

1 -* S. However, the

process was relatively slow even at 400 °C and was apparently a result

of the dissolution of S' precipitates within grains and the diffusion

of solute to S precipitates at grain boundaries. At lower temperatures,

the formation of G.P.B. zones and homogeneously distributed S' was

similar to that after sequence A treatments.

As might be expected on the basis of the different transformation

processes, the sequence B C-curves differ from those for the sequence A

"pre-aging" treatments. While the sequence A treatment resulted in the

nucleation of 0, S and S' phases within grains, only the S' phase was

observed after a sequence B treatment. Differences in the nucleation,

growth, coarsening and transformation kinetics of the various phases

explain the differences in the C-curves.

Effect of "Pre-aging" on T851 Microstructure . Two sequence A

specimens "pre-aged" at 350 °C were examined after being subjected to a

T851 temper treatment. The "pre-aging" treatments were carried out for

20 s and 109 s each; the microstructures are shown in Figures 87(a) and

(b), respectively. Large precipitates at dispersoid particles are

readily identified as being produced during the "pre-aging" treatment

by comparison with Figure 73 given a T851 treatment without "pre-aging"

and Figure 75(c) which was subject to "pre-aging" and not processed to

the T851 condition. The "pre-aging" induced precipitates were determined

to be the 0 and S phases. The small needle shaped precipitates are, of

course, the S' phase which was formed during final aging at 190 °C.

57
The low density of S' precipitates around "pre-aging" induced precipitates

is indicative of the extent to which solute was lost during "pre-aging".

4. Discussion; Relationship Between Microstructure and Properties

Property changes associated with "pre-aging" are a result of the

nucleation, growth, and coarsening of precipitates during the "pre-aging"

treatment. TEM studies have demonstrated that the precipitates may

differ significantly in type, size, concentration, distribution and

coherence depending on the nature of the "pre-aging" treatment. The

changes in properties are strongly time dependent indicating that

transformation kinetics are of critical importance. "Pre-aging" induced

precipitates may have both a direct effect on properties and an indirect

effect. The indirect effect is a result of the consumption of solute

elements (Cu and Mg) by the "pre-aging" induced precipitates. This

reduces or in the worst case precludes the formation of G.P.B. zones

and precipitates which normally develop during the final natural or

artificial aging treatment.

Although no attempt will be made to develop quantitative relationships between properties and microstructural quantities such as precipitate type, size and concentration, an effort will be made to provide a

qualitative explanation of the effect that "pre-aging" has on properties.

In addition to the TEM studies, the discussion will depend strongly on

the results of the room temperature aging experiments presented below,

on measurements of the age hardening behavior of Al-Cu-Mg alloys from

Hardy (35) and Beton and Rollason (38), and on phase diagram information,

also from Hardy. A major part of the discussion is concerned with the

T4 condition since most of the TEM specimens examined were in that

condition. Also, in the T4 condition, the effects of "pre-aging" can

58
be related directly to property changes without the need to account for

additional modifications brought about by deformation and elevated

temperature aging which accompanies processing to the T351 and T851

conditions. The effect of "pre-aging" on these temper conditions will,

however, be considered in the final part of this section.

(a) T4 Condition

In general, the T4 condition refers to materials that have

been allowed to age at room temperature after the quench from solution

treatment without intervening or subsequent treatments. In the present

investigation, the T4 appellation is also applied to specimens given a

"pre-aging" treatment prior to room temperature aging. The room temperature age hardening responses of many of the specimens prepared for TEM

study (Figure 66) are shown in Figure 88. For reference, the age

hardening behavior of a specimen quenched directly into ice water from

the solution treatment temperature is shown in Figure 89. The data

plotted in Figures 88 and 89 were obtained as follows. After quenching

in ice water at the end of the "pre-aging" cycle for Figure 88 or after

the direct quench for Figure 89, the specimen was immediately immersed

in liquid nitrogen for storage until room temperature aging was commenced

The specimen was brought quickly to room temperature by agitating in a

bath of water. Hardness measurements were made according to the Rockwell

"B" method noted earlier. The first determination was obtained at 1 to

2 mins after the specimen had reached room temperature. Hardness measure

ments were made periodically thereafter until no further change was

observed. In many cases, an initial increase in hardness was observed

during the first 2 or 3 mins followed by an incubation period and then

hardening to the final value. The initial change in hardness will not

59
be considered further here, however it is thought that it may be associated with spi nodal decomposition. The effect would be better studied

at lower temperatures.

Consider first the effect on hardness of the sequence A

"pre-aging" treatments. Final hardness, post quench incubation hardness

and the difference between final and incubation hardness values are

plotted in Figures 90(a), (b) and (c) respectively. The plotted results

were taken from room temperature aging curves in Figures 88 and 89

after converting from Rockwell "B" to the Vickers hardness scale by

utilizing a table from Mondolfo (39). The Vickers scale bears a more

nearly linear relationship to strength than the Rockwell "B" scale. To

explain the variation in hardness shown in Figure 90 the following must

be considered:

1. Hardening by G.P.B. zones

2. Solid solution hardening

3. Precipitation hardening

4. Grain size effects

Age hardening of Al-Cu-Mg alloys with a Cu:Mg ratio of 2.2:1

occurs at room temperature by the formation of G.P.B. zones. A maximum

increase in hardness for a given Cu and Mg composition is obtained when

the concentration of G.P.B. zones is maximized. This is achieved by

solution heat treatment above the S+a boundary and quenching at a

sufficiently high cooling rate to retain all the Cu and Mg in solution.

A reduced quench rate, or "pre-aging" treatment as was done in this

investigation at a temperature below the S+a boundary can lead to a

reduction in the solute concentration through the precipitation of the

S, S', and 0 phases and a diminished concentration of G.P.B. zones upon

60
subsequent aging. The effect of "pre- aging" on the incremental increase

in hardness associated with G.P.B. zone formation after room temperature

aging to a stable condition is clearly demonstrated in Figure 90(c).

On "pre-aging" at 400 °C, the hardness increment falls until a constant

value is reached after 100 s. The behavior is similar for 350 °C but

2 the constant level which is obtained after 500 s is only about 10 N/mm .

At 300 °C, extended "pre-aging" eventually eliminates any observable

room temperature age hardening effect. Since G.P.B. zone formation

probably requires some minimal level of solute supersaturation, Figure

90(c) demonstrates that aging for a sufficient time (> 500 s) at a

temperature between 300 and 350 °C reduces the supersaturation below

the minimum value. The TEM studies determined that the loss of solute

was primarily a result of the formation of the 0 and S phases within

grains and the S phase at grain boundaries. In effect, the material

behaves like an alloy with a lower solute concentration. For the Al-Cu-Mg

system according to Hardy (35) the a/(a+S) boundary is given by log^

[Cu][Mg] = 5. 603-3975/T, where [Cu] and [Mg] are concentrations in

atomic percent and T is the absolute temperature. Using this equation

the equilibrium concentration of Cu+Mg at 400 °C is approximately 2.2

wt.% while at 300 °C it is only 0.6 wt.%. The relationship between the

hardness change produced by G.P.B. zone formation and solute concentration for the Al-Cu-Mg system plotted from the data of Beton and Rollason

is shown in Figure 91. The relationship is not linear, particularly at

concentrations above ^ 3.5 wt.% (Cu+Mg). Between 1 wt.% (the minimum

value for which data was reported and 3.5 wt.% the change in hardness

is approximately proportional to the concentration. Temperatures

corresponding to the a/(a+S) phase boundary location are also indicated

61
in Figure 91. It can be seen that the hardness increment at concentrations below 1 wt.% (Cu+Mg) attained by equilibration at temperatures

< 326 °C is very small. Thus, it may be inferred that by "pre-aging"

2024 aluminum alloy for a sufficiently long period of time at temperatures

close to 300 °C the concentration of solute remaining in solution would

be reduced to a point where few G.P.B. zones would form. A direct

comparison between 2024 aluminum alloy and ternary Al-Cu-Mg alloys, of

course suffers from the fact that 2024 aluminum alloy contains additional

elements that can effect its aging behavior.

Figure 90(c) also provides information on the rate at which

equilibration occurs. The rate is determined by the nucleation and

growth kinetics of "pre-aging" induced precipitates. At 400 °C where

nucleation is comparatively difficult as is demonstrated by the absence

of precipitates after "pre-aging" for 20 s, there is little change in

solute supersaturation. After 100 s large precipitates were observed

throughout the matrix and according to Figure 90(c) an equilibrium

concentration of solute had been reached. There is a complication

here, of course, because the 0 phase, the major phase within grains,

consumes mainly Cu leaving a surplus of Mg. What effect this may have

on G.P.B. zone formations is not known, although it is not expected to

be large.

At 300 °C nucleation occurs more rapidly as evidenced by a

high concentration of 0 precipitates. The initial rate of approach to

equilibrium was also the most rapid for the temperatures studied.

The final hardness shown in Figure 90(a) can be regarded as

the sum of the incremental increase associated with G.P.B. zone formation

and an incubation hardness value that varies with "pre-aging" time and

62
temperature. At 400 °C there is no change in hardness before 20 s and

a logarithmic decrease thereafter. Initially a rapid drop in incubation

hardness is observed at 350 °C and 300 °C which is followed by a slower

logarithmic decay at about the same rate as at 400 °C. The difference

between the initial behavior at 400 °C and that at the two lower temperatures is consistent with differences in the microstructures. After

aging for 20 s at 400 °C there was no evidence that precipitation had

occurred except at grain boundaries where very small S precipitates

were observed. In contrast relatively large precipitates were present

after "pre-aging" for 20 s at 350 and 300 °C. The initial large rate

of decrease in incubation hardness can be attributed to a decrease in

solid solution hardening as a result of the nucleation and growth of

the 0 and S phases. A measure of the extent to which Cu and Mg in

solution can effect hardness is indicated by the as-quenched hardness

vs (Cu + Mg) concentration curve obtained by Hardy (35), Figure 92.

The period during which the rapid loss of solute from the matrix takes

place corresponds to the portion of the curves in Figure 90(c) showing

a rapid change in hardness increment due to G.P.B. zone formation. The

0 and S phase precipitates may increase the hardness but apparently not

nearly enough to offset the effect associated with the loss of solute.

The continued decay in hardness with increasing time at all temperatures

can be attributed in part to a further loss of solute from the matrix

due to the gradual approach to equilibrium. Eventually, this involves

the slow dissolution of 0 and the growth of S, where it will be recalled

that the S phase consumes both Cu and Mg and the 0 phase only Cu, i.e.

the change would mainly be associated with the loss of Mg from solution.

Loss of other elements in solution, for example Mn and Cu to the

63
Cu

2 Mn

3

Al

2 0 dlspersold phase may also contribute to the reduction In

plateau hardness. Part of the decay In hardness can also be attributed

to precipitate coarsening, a reduction In the number density of precipitates due to Ostwald ripening, and transformation of S' to S.

Finally, prolonged "pre-aging" especially at temperatures

close to the solution treatment temperature could lead to grain growth

and a reduction in hardness. However, grain growth was not observed

after 7200 min at 400 °C and is made unlikely by the closely spaced S

precipitates at the grain boundaries.

Turning to the sequence B treatment, a similar approach can

be taken to relate the effects of "pre-aging" induced phase transformation

to property changes. Figure 93(a), (b), and (c) summary the room

temperature age hardening response of specimens "pre-aged" at 300 °C,

350 °C and 400 °C. Figure 93(a) presents the final hardness values,

Figure 93(b) gives the incubation values and Figure 93(c) shows the

incremental change in hardness. The room temperature age hardening

response after sequence B treatment at 400 and 350 °C differs from that

observed for a sequence A treatment. After "pre-aging" for 20 s at

400 °C, the incremental change in hardness, Figure 93(c), has dropped

to ~ 200 N/mm from 330 N/mm , the directly quenched value. No

change was observed after a sequence A treatment for the same period of

"pre-aging". The initial rapid decrease in hardness increment after

the sequence B treatment is consistent with the presence of a high

concentration of S' precipitates within grains and S at grain boundaries

observed in the TEM. It may be recalled that precipitation was just

beginning at grain boundaries after the sequence A treatments at 400 °C

for 20 s. The low rate of nucleation at 400 °C is overcome in the

sequence B treatment during the quench-reheat portion of the cycle.

64
"Pre-aging" at 400 °C for 100 s and 500 s resulted in increasingly larger S' precipitates with little change in hardness increment.

Thus the process during this period was mainly one of coarsening and

did not involve a change in the volume fraction of the S' phase, i.e.

the concentration of solute in solution did not change appreciably.

"Pre-aging" for 7200 min led to an increase in the hardness increment

and resulted in a microstructure that consisted of only a few large S'

precipitates within grains and very large S precipitates at grain

boundaries. Thus, prolonged aging apparently resulted in the dissolution

of much of the S' phase within grains and a corresponding growth of the

S phase at grain boundaries. The reason for the increase in hardness

increment after 7200 min at 400 °C is not understood. The same effect

was not observed for the corresponding sequence A treatment. The

difference may be associated with the much lower incubation hardness

after the sequence B treatment.

The hardness increment vs "pre-aging" time curves at 350 °C

and 300 °C, Figure 93(c), are similar to the companion sequence A

curves, Figure 90(c), although the decrease occurs more rapidly at

350 °C than was obtained for the sequence A treatment.

The incubation hardness curves, Figure 93(b) for the sequence

B treatments at 400 °C and 300 °C show an approximate decrease in

hardness with logarithmic time from the onset of "pre-aging". At

300 °C, however, there is an initial rise during approximately the

first 100 s followed by a decrease at about the same rate as for 350 °C

and 400 °C. The incubation hardness over the range of "pre-aging"

times studied is greatest at 300 °C and smallest at 400 °C. This

differs from the sequence A treatment results where the highest hardness

65
was found at 400 °C and the lowest at 350 °C. On the basis of the TEM

studies, it appears that the incubation hardness behavior of the sequence

B specimens is strongly dependent on the formation and subsequent

coarsening of heterogeneously nucleated S' precipitates. The initial

rise in hardness observed at 300 °C can be attributed to the presence

of very small, homogeneously distributed S' precipitates. Since the

hardness increment curves, particularly at 300 °C and 350 °C, indirectly

indicate a decreasing concentration of (Cu and Mg) in solution until

"pre-aging" times of 500 s and 100 s, respectively, have been reached,

solid solution hardening must also contribute during that period to the

observed incubation hardness. Comparing sequence A and B final and

incubation hardness curves, Figures 90 and 93, at 350 °C and 300 °C, it

is seen that after a "pre-aging" time of 500 s the sequence A curves

fall more rapidly to considerably lower hardness valves. This is

apparently a result of the considerable stability of the S' phase in

the sequence B specimens against coarsening and transformation to the S

phase.

(b) T351 and T851 Conditions

The foregoing discussion of the influence of "pre-aging" on

T4 microstructure and properties is directly applicable to the T351

condition. It is only necessary to include the effect on properties of

an incremental increase in the dislocation density. To a first approximation, the difference between T4 and T351 properties should be constant,

independent of the "pre-aging" treatment. This does not necessarily

hold true for the T851 condition. The microstructure of a specimen

that has been directly quenched to ice water from the solution treatment

temperature and processed according to T851 specifications consists of

66
a uniform distribution of heterogeneously nucleated S' precipitates.

To obtain this structure, it is necessary that solute be available for

the nucleation and growth of the S' precipitates at dislocations during

final aging. The solute may be supplied either directly from a supersaturated solid solution or by G.P.B. zones and possibly from subsequently

formed homogeneously distributed S' precipitates. Because of the high

rate at which G.P.B. zones form, its unlikely that much solute is

supplied directly from the supersaturated solid solution.

Solute that has been consumed by the relatively stable "preaging" induced precipitates is no longer available for the formation of

S' precipitates during subsequent aging treatments. Moreover, the

"pre-aging" induced precipitates may continue to consume solute during

the final aging treatment. Finally, the "pre-aging" induced precipitates,

as already discussed, contribute to strength and are subject to overaging

during further heat treatment and provide an additional mechanism for

the loss of strength. It is therefore not surprising that the strength

of specimens given a T851 treatment was actually reduced below that of

the T351 condition as opposed to the properly quenched material.

67
VI. ULTRASONIC CHARACTERIZATION

Correlation of Ultrasonic Data with Hardness after

Aging to T4, T351, and T851 Conditions

1 . Introduction

The study of ultrasonic wave propagation can provide information

about the elastic properties and absorption characteristics of the

material in which the wave propagates. These properties are intimately

related to microstructure, hence the thermomechanical treatment that

the material underwent prior to the ultrasonic measurements, thus

enabling a nondestructive evaluation and characterization of the materials

properties.

The speed of sound in a solid is most often used in nondestructive

testing as a constant programmed into an ultrasonic thickness gauge.

However, in the realm of nondestructive characterization of materials,

the absolute measurement of longitudinal and shear wave velocities can

be used to calculate many useful material parameters, e.g., Young's

modulus (E), shear modulus (G), bulk modulus (B), and Poisson's ratio

(v). The determination of these effective moduli is based on the

assumption that the material is isotropic. The calculation of the

moduli requires knowledge of the value for the density of the material.

The energy loss, or ultrasonic attenuation, of elastic waves propagating in a solid, may be divided into contributions from geometrical and

intrinsic effects. Geometrical effects include reflection and refraction

at a free surface, grain or phase boundary, beam divergence due to diffraction, as well as waveguide effects due to multiple boundary surfaces.

Intrinsic effects include scattering of the ultrasonic wave at inhomogeneities, interaction with thermal phonons, dislocation damping, and conversion of sound energy to heat as a result of elastic deformation.

68
Unlike geometric effects, intrinsic effects may change in time if

the internal crystal lattice or defect structure undergoes change.

During precipitation hardening of 2024 aluminum alloy the main contribution to the attenuation is believed to be scattering of the ultrasonic

wave by newly precipitating phases. The magnitude of the absorption

will be determined by the magnitude of the stress field at the boundary

between the precipitate and the matrix which, in turn, is determined by

the precipitate size and the degree of coherency of the precipitate

with the matrix (41). An attenuation peak should be observed when the

size of the precipitate and associated stress field around it reach a

magnitude that produces the maximal interaction with the ultrasonic

wave. This value of attenuation will also depend on the wavelength of

the ultrasonic wave relative to the size of the scattering center. As

frequency increases, the wavelength decreases, and hence the size of an

effective scattering center will also decrease.

With continuous monitoring, the relationship between aging time

and change in sound velocity and ultrasonic attenuation can be determined.

The information may provide an insight into the kinetics and mechanism

of precipitate formation and growth in the age-hardening aluminum

alloys. Several investigations (41-43) have demonstrated that the

ultrasonic method is operationally feasible for monitoring the precipitation process in aluminum alloys in real time, over a wide temperature

interval. The calculated activation energies and characteristic kinetic

parameters could be related to the mechanism of the formation and

growth of the precipitates, and their effect on the ultimate mechanical

properties of the alloy. The extension of techniques for ultrasonic

NDE to industrial applications requiring on-line, real-time monitoring

69
during thermomechanical treatment would necessitate modifications of

the experimental approach for the specific technological application.

The objective of the present study was to establish the existence

of a correlation between ultrasonic data and mechanical properties

(e.g. hardness, yield strength) of 2024 aluminum alloy specimens after

they had been subjected to carefully controlled thermomechanical treatments. For this purpose, the absolute, rather than the relative,

values of the sound velocity and ultrasonic attenuation are required.

The data have to be determined, consistently, to a high degree of

accuracy so as to enable a comparison to be made between specimens of

unknown thermal history, and to correlate these data with the ultimate

mechanical properties of the material. Ideally, as an NDE technique

the ultrasonic measurement should provide the necessary information in

order to specify, nondestructi vely , the mechanical properties of the

alloy.

As will be shown, a correlation between the absolute values of the

sound velocity and ultrasonic attenuation, and the hardness data was

found. Regardless of the thermal sequence of the aging process, maximal

hardness (and strength) of the 2024 aluminum alloy was found to be

related to a definite range in the values of the sound velocity. A

consistent difference between the sequence A and B "pre-aging" heat

treatments was revealed. Furthermore, a difference was also noted in

the values of the sound velocity between the T851 and T351 tempers of

2024 aluminum alloy.

2. Experimental Procedure

A systematic ultrasonic examination was carried out on a total of

about 140 specimens that underwent the variety of controlled "pre-aging"

70
heat treatments reported in earlier sections of this report. The

samples were subjected to "pre-aging" sequences A and B, after solution

treatment at 495 °C, and prior to T4 or the stretching operation that

precedes the aging to T351 and T851. The specimens were flat and

parallel and had the following dimensions: 150 x 25.4 x 6.35 mm (in

thickness). The measurements were performed by means of conventional

ultrasonic equipment (MATEC pulse generator and receiver and delta-time

HEWLETT-PACKARD oscilloscope).

Specimens and ultrasonic sensors were immersed in a disti lled-water

tank to avoid complications caused by variations in couplant thickness.

Constant separation between the transducer and the specimen surface was

maintained. The specimen mounting stage was designed to allow adjustment

of the parallelism between the surface of the transducer and the insonated

surface of the sample. Before making velocity and attenuation measurements the ultrasonic pulse train was adjusted to obtain an exponential

decay of the successive echoes combined with a maximal number of echoes

in the pulse train. The pulse superposition technique was employed for

sound wave velocity determination. Attenuation was monitored by means

of a MATEC Automatic Attenuation Recorder which provides analog processing of the video output of the MATEC pulser/receiver.

The absolute values of the sound velocity and ultrasonic attenuation

were determined to within ±1 m.s

"* and ± 0.02 dB, respectively, It

should be pointed out, in this context, that both the precision and

accuracy of velocity and attenuation determination can be markedly

improved by increasing commensurately , the specimen thickness since the

accuracy in velocity and attenuation depends also on the accuracy of

the thickness measurement.

71
3. Results and Discussion

Figure 94 shows the variation of the ultimate sound velocity in

2024-T4 aluminum alloy as a function of "pre-aging" time for various

"pre-aging" temperatures, between 300 and 400 °C, for both sequences A

and B.

For long "pre-aging" times, exceeding 2000 min the sound velocities

converge at the same value. Apparently, the "pre-aging" temperatures

are sufficiently elevated to allow the operation of a diffusional

process for the formation of a precipitate morphology that yields a

similar average sound velocity.

Figure 94 shows that sequence A yields lower initial sound-wave

velocities than sequence B. The main reason for this behavior is the

fact that the sequence B samples exposed to low temperatures contain a

high concentration of S' precipitates. TEM evidence for 20 s "pre-aging"

at 400 °C following sequences A and B are shown in Figures 75(a) and

83(a), respectively. In the sequence A treatment, there is essentially

no difference in microstructure compared with the direct quench from

the sol utionizing temperature. In contrast, sequence B treatment

results in relatively high concentration of S' precipitates throughout

the matrix (Figure 83(a)). This difference in microstructure appears to

be the cause of the significant change in the initial sound-wave

velocity, Figure 94.

Figure 95 exhibits the change in ultimate hardness of 2024-T4

aluminum alloy as a function of "pre-aging" time for various "pre-aging"

temperatures, between 300 and 400 °C, for both A and B sequences.

"Pre-aging" at the temperatures indicated, for both sequences,

reduces the ultimate hardness values determined after completion of the

aging process at room temperature. In general, sequence A leads to a

72
greater reduction in hardness than sequence B. For long "pre-aging" times

the least detrimental "pre-aging" temperature was found to be 400 °C. It

can be reasoned that at 400 °C, much of the Cu and Mg is retained in solution in the aluminum matrix. When the alloy "pre-aged" at 400 °C is quenched

to room temperature, conditions are quite favorable for formation of

G.P.B. zones typical of natural aging. The expected microstructure is,

therefore, a mixture of low temperature precipitates (G.P.B. zones) as

well as some 0

1 , S' and S that formed at 400 °C prior to aging, Figures

75(a) and 75(b).

The lowest ultimate hardness values after prolonged "pre-aging"

times were obtained for the sequence A "pre-aging" treatment at 300 °C.

These specimens were not exposed to room temperature before aging.

Therefore, one should expect an extensive formation (by volume) of 0

and S precipitates that are responsible for the drastic decrease in

hardness. This line of reasoning is compatible with the observed high

values in the sound velocity (Figure 94). The incoherent intermetal 1 ic

solute-rich compounds 0 and S have a significantly higher specific

velocity than the aluminum matrix, thus contributing to a net increase

in the measured sound velocity (50).

Figure 96 shows the influence of the "pre-aging" time, for various

temperatures and heat treatment sequences, on the ultrasonic attenuation.

The highest attenuation value, for all "pre-aging" times, was

found for sequence A and 350 °C. As was mentioned earlier, 350 °C is

apparently the most efficient "pre-aging" holding temperature for the

formation of the incoherent 0 and S as well as of the semi-coherent S'

precipitates, Figures 75(c) and 75(d). The size distribution and

concentration of these precipitates is particularly conducive to an

73
enhanced attenuation effect obtained by attenuation measurements at

frequencies of about 10 MHz. In addition, dislocation loops formed

around precipitates, that have lost coherency with the aluminum matrix

(41), Interact most effectively with the propagating ultrasonic waves,

thus yielding a high attenuation value. As previously noted (41) this

behavior of the ultrasonic attenuation is compatible with the substantial

decrease in hardness, shown in Figure 95.

The lowest absolute values as well as the smallest changes of

attenuation, were observed In specimens that were "pre-aged" at 400 °C,

for both A and B sequences. The behavior of the ultrasonic attenuation

is in agreement with the behavior of the sound velocity and hardness

(Figures 94 and 95) in the sense that at 400 °C only a small volume

fraction of 0 and S and S' is formed. The final fully aged microstructure for this specific "pre-aging" holding temperature consists mainly

of G.P.B. zones that were formed upon subsequent aging at room temperature. These low temperature precipitates are apparently too small in

size to significantly contribute to the observed ultrasonic attenuation.

For zero "pre-aging" time, the attenuation curves seem to extrapolate

to a common low value. For prolonged "pre-aging" times, the attenuation

values are significantly different due to the differing microstructures

and size distribution of the precipitates.

Figure 97 exhibits the effect of "pre-aging" (sequence A and B) on

the velocity, attenuation, hardness and electrical conductivity when

specimens were held for 60 minutes at different "pre-aging" temperatures,

prior to the natural aging process.

Sequence A yields a lower hardness for each of the "pre-aging"

temperatures, between 300 and 400 °C. The lowest hardness values are

74
found at 300 and 350 °C. The reason that the B sequence yields consistently higher hardness values than A sequence is because the S' precipitates formed during sequence B "pre-aging" make an effective contribution

to hardness.

The sound velocity, for both A and B sequences, exhibits a maximum

for the 350 °C "pre-aging" temperature. This behavior is in agreement

with the fact that at 350 °C, a substantial volume fraction of 0, S'

and S is observed, Figures 82 and 83 thus contributing to the observed

increase in the sound velocity. The 400 °C "pre-aging" temperature is

too close to the solution temperature of 495 °C to form an appreciable

volume fraction of incoherent precipitates that may contribute to an

increase in the sound velocity. Therefore, the velocity values for the

400 °C "pre-aging" temperature are more typical of the low temperature

precipitates (G.P. zones) formed at room temperature aging, after

"pre-aging," since they constitute the major volume fraction of the

precipitated phase.

The difference in the observed sound velocities between B and A

sequences is consistent with the high concentration of S' precipitates

that are nucleated during the quench and reheat cycle in the B sequence

treatment. The nucleation process is less efficient in the A sequence.

The behavior of the ultrasonic attenuation is compatible with the

behavior of the hardness. Higher attenuation values are indicative of

lower hardness (Figures 98 and 99). The attenuation behavior confirms

the model that at this temperature the majority of the precipitates are

apparently of the semi-coherent and incoherent character surrounded by

dislocation loops that strongly interact with the propagating sound

75
wave, thus yielding high attenuation values. The higher attenuation

values of B sequence, relative to A sequence, as seen in Figure 98, may

again be explained by the larger amount of low temperature precipitates

that cause scattering of the ultrasonic waves.

The electrical conductivity (Fig. 97), as obtained from eddy

current measurements after completion of aging, shows that A sequence

treatment leads to higher electrical conductivity than the corresponding

B sequence treatment. Moreover, conductivity increases, for both A and

B, with increasing "pre-aging" temperature.

In section VII of this report, it will be shown that the electrical

conductivity increases with the "purification" of the aluminum matrix

by mean of segregation of the solute atoms (mainly Cu and Mg) and

formation of solute-rich precipitates. Also, the electrical conductivity

decreases by the presence of small precipitates (e.g. G.P.B. zones) that

are quite effective scatterers of electrons. However, this negative

contribution to the conductivity is offset by the "purification" process,

particularly at temperatures where an enhanced diffusion of the solute

species from the supersaturated aluminum matrix is possible.

The conductivities of A and B sequences converge at a value of

about 41.0% IACS, at 400 °C "pre-aging" similar to non "preaged" 2024 A1

alloy (51). This experimental observation may be explained in terms of

the equal amount of Cu and Mg atoms that left the aluminum matrix. This

does not necessarily imply that the microstructure of the precipitation

must also be similar or identical. In fact, TEM observations (Section V)

point quite to the contrary, as do sound velocity, attenuation, and

hardness measurements. The effect of microstructure on the measured

electrical conductivity is complex, and depends quite insignificantly

on the morphology.

76
Figure 98 shows the influence of the "pre-aging" sequence (A and

B) on the relationship between hardness and ultrasonic attenuation in

T4 heat treated 2024 steel. For each sequence, separately, a good

correlation was observed: high hardness is directly related with low

attenuation values.

The correlation between hardness and attenuation is very important

from an NDE point of view. This correlation suggests that within a

limited range of thermomechanical treatments the hardness can be uniquely

determined by means of ultrasonic attenuation. This correlation could

be quite useful for in-process monitoring where small deviations from the

norm of a wel 1 -specified thermomechanical treatment is encountered. On

the contrary, Figure 98 indicates that substantial deviations from the

norm, as obtained for example for sequence A relative to B, renders such

predictive capability rather difficult.

Figure 98 indicates that sequence B exhibits consistently higher

hardness values, than sequence A for a given value of attenuation.

Consequently, sequence B should contain a higher volume fraction of the

low temperature GP zones an intermediate (S') precipitates that are known

to contribute substantially to hardness and enhance the attenuation

of ultrasonic waves.

Figure 99 shows hardness versus ultrasonic attenuation for 2024-T351

temper, "pre-aged" by the sequence A scheme, for various "pre-aging"

temperatures between 300 and 400 °C. The observed change in hardness

is between 65 and 80 HRB. The ultrasonic attenuation varies linearly

and i nversely with hardness. For this range of hardness (between 65

and 80 HRB) an approximately linear decrease in attenuation was also

observed for 2024-T4 (Figure 98).

77
Figure 99, similar to Figure 98, does not differentiate between

"pre-aging" times and "pre-aging" temperatures. This figure represents

a strictly NDE correlation giving a qualitative change between two properties, hardness and ultrasonic attenuation . The behavior of these two

properties is a consequence of differing microstructures, precipitate

sizes, states of coherency etc., but the integral effect yields a linear

correlation that may be utilized in qualitative NDE characterization.

Figures 100-102 depict the correlation between sound velocity and

hardness for 2024 A1 alloy of different tempers, namely; T4, T351 , and

T851

.

Analogous with Figures 98 and 99, Figures 100-102 present an NDE

correlation. Neither "pre-aging" times nor "pre-aging" temperatures,

nor heat treatment sequences are specified. This emphasizes the potency

of the correlation between sound-wave velocity and hardness of the

completely aged alloy . The correlation (Figures 100-102) is parabolic,

where the maximal hardness (or yield strength) is attained for a narrow

window of sound-wave velocities. When maximal strength is required,

the NDE technique employing ultrasonic sound velocity measurements

would probe and monitor for a definite predetermined range of sound-wave

velocities. The absolute values of these velocities may change slightly

with the selected temper or state of plastic deformation prior to

aging. Plastic deformation prior to aging (e.g. T351 compared with T4

temper) increases the absolute value of the ultimate sound velocity, at

the peak hardness, 6369 versus 6363 m.s \ respectively. Larger volume

fractions of precipitated phases will also contribute to an increase in

the absolute value of the sound velocity.

78
VII. EDDY-CURRENT CONDUCTIVITY VERSUS HARDNESS DURING AGING

1 . Introduction

The kinetics and mechanism of the precipitation hardening process

in pseudo-binary aluminum alloys has been extensively investigated

(43). A discussion of the nature of the nucleation and growth of

precipitates in binary aluminum alloys was developed by Lorimer and

Nicholson (44). However, relatively little work has been reported on

studies involving nondestructive characterization of the technologically

important 2024 aluminum alloy during the aging process. This alloy

acquires high strength and hardness when subjected to controlled thermomechanical treatments involving natural or artificial age hardening.

Electrical conductivity measurements, employing a.c., d.c., and

eddy currents are often utilized to determine the kinetic behavior of

the precipitation process, or to evaluate nondestructively the progress

of the age hardening sequence (46-48). The variation of electrical

conductivity as a function of aging is complex and results from a

number of factors (49), a major contribution is associated with changes

in the scattering of conduction electrons. Such scattering can be

caused by:

a. vacancies quenched- in during rapid cooling following solution

heat treatment

b. alloying atoms present in the matrix

c. G.P. zones with dimensions comparable with the electron

wavelength

d. coherency strains and dislocations at boundaries of the zones

or clusters.

79
During aging at low temperatures, as for example, below 50 °C, two

competing processes are operative. There is a decrease in the number

of quenched-in vacancies, and a clustering of Cu and Mg, which are the

major additions to A1 in the 2024 alloy. These processes tend to

increase the measured electrical conductivity but simultaneously there

is a more significant decrease in the electrical conductivity due to

the appearance of newly formed precipitates which are small enough to

be effective as scatterers of electrons. At higher temperatures,

precipitates of larger size are formed, which are not as effective

scatterers resulting in an increase of conductivity to an ultimate

value determined by the solubility of impurity atoms at the specific

aging temperature. The absolute value of electrical conductivity

achieved after any given aging time is dependent on the aging temperature due to the fact that the geometry and size distribution, and type

of the precipitates are governed by a thermally-activated process (35).

Addition of trace elements to the Al-Cu-Mg system may have a

profound effect on the kinetics and mechanism of the precipitation

process. Some elements may suppress one stage of the aging process

while stimulating another. At low aging temperatures, zone formation

may be retarded, apparently due to a strong vacancy-trace element

interaction which prevents vacancies from enhancing the diffusion of Cu

atoms in the a-Al matrix. At higher aging temperatures these elements

may accelerate the formation of intermediate precipitates (0', S') by

reducing interfacial energies.

The influence of precipitation kinetics on eddy-current conductivity

and hardness during the aging of 2024 aluminum alloy was investigated

by means of dynamic eddy-current conductivity measurements. Aging

80
temperatures between 21 and 190 °C were used, and measurements were

made on both unstretched and plastically deformed samples (three percent

permanent strain). The two measurement techniques, electrical conductivity and hardness, respond in a complementary manner to the varying

microstructures that form during different stages of the aging process.

2. Experimental Procedure

The specimens used in this investigation were prepared from the

0.635 cm thick 2024 aluminum alloy plate supplied in the F condition.

Results of the chemical analysis were given in an earlier section.

Prior to the thermomechanical treatment, the specimens were machined

into rectangular bars 170 x 25.4 x 6.35 mm. Solution heat treatment at

495 °C for 75 minutes was followed by ice-water quenching. The specimens

were then stored in liquid nitrogen to protect them against changes in

microstructure. Hardness measurements, carried out at subzero temperatures, before and after storage, verified no natural aging had occurred

during storage.

Some of the specimens were plastically deformed to 3 percent

elongation in a tensile testing machine. This stretching process was

used to simulate the T351 and T851 tempers. The stretching was performed

while the specimens were kept at subzero temperatures in order to avoid

natural aging during the plastic deformation procedure.

The specimens were then aged in a thermostatic bath containing

heated oil kept in continuous circulation. The bath could be maintained

within 0.05 °C at any temperature between ambient and 200 °C.

Electrical conductivity measurements were performed on the specimens

by means of SUPERHALEC (England) and VERIMET (USA) eddy current monitors

and probes.

81
Each eddy-current conductivity measurement was preceded and followed

by calibration of the eddy current monitors against NBS secondary

standards. These standards are calibrated at 20 °C; hence all reported

conductivity values herein refer to the conductivity value at 20 °C.

The measured values were considered accurate to within ± 0.05 percent

IACS 9

.

Hardness measurements were made using a WILSON hardness tester.

The precision of the hardness values on the Rockwell B scale was approximately ± 1 units.

Room temperature measurements of hardness were made at prescribed

intervals throughout the aging treatment. The specimens were removed

from the isothermal bath, cooled rapidly to room temperature, and

subjected to electrical conductivity and hardness measurements. Whenever

necessary due to the time delays in making a measurement, the specimens

were temporarily stored in liquid nitrogen.

3. Experimental Results

Figure 103 shows the variation of the electrical conductivity as a

function of aging time of 2024 aluminum alloy at different isothermal

holding temperatures. This series of specimens did not undergo plastic

deformation (stretching) prior to aging. The value of electrical

conductivity, before aging, was found to be 31.5% IACS. The salient

features in Figure 103 are:

a. An initial decrease in conductivity at each of the isothermal

aging temperatures in the range between 21 and 190 °C. The

rate of initial decrease in conductivity depends on the aging

9 International Annealed Copper Standard

82
temperature, increasing with increased aging temperature. In

the temperature range between 150 and 190 °C, the rate of

initial decrease in conductivity is relatively large and

virtually identical within this set.

b. A sigmoidal decrease in conductivity with aging time was

observed for the 21, 35, and 50 °C isotherms. For these

aging temperatures, the ultimate value of the fully aged

alloy was found to be 29.5% IACS.

c. In contrast to the variation of electrical conductivity as a

function of aging time at low temperature (below 50 °C), the

conductivity above 150 °C i ncreases with increasing isothermal

aging temperature, after the initial decrease during the

first five minutes of the aging process. The change in

conductivity with aging time is extremely sluggish at 150 °C,

exhibiting a nearly constant value for about 30 hours of

aging.

d. For unstretched specimens aged at 170, 180, and 190 °C, the

most pronounced rate of increase in electrical conductivity

occurs at the highest aging temperature. The ultimate value

of the electrical conductivity, ^ 40% IACS, was attained after

about 50 hours at 190 °C, and after somewhat longer times at

the lower isothermal holding temperatures.

Figure 104 exhibits the variation of hardness as a function of

aging time for the unstretched specimens at aging temperatures within

the range of 150 and 190 °C. In general, the rate of increase of

hardness with aging time increases as aging temperature increases.

Overaging was observed at aging temperatures of 180 and 190 °C, in the

former case occurring at about 1,000 min and at approximately half that

83
value at the higher temperature. With decreasing aging temperatures,

the peak value of hardness increases, and occurs after a longer aging

time.

Figures 105 and 106 compare the variation of hardness and electrical

conductivity of unstretched and stretched specimens at 190 °C (Figure

105) and 150 °C (Figure 106). The initial hardness of the stretched

specimens is higher by about 5 units of the Rockwell B scale. The

change in hardness with aging time is quite similar for both sets, and

the initial difference in hardness is essentially preserved. In Figure

105, the stretched specimens show the same response to overaging as the

unstretched ones.

The initial electrical conductivity, before aging, of the 3 percent

stretched specimens is lower by about 0.4% IACS than that of the

unstretched specimens. However, at 190 °C, (Figure 105), the rate of

change of conductivity with aging time after about 100 hours is higher

in the stretched specimens but the two curves begin to converge after

an aging period of about 100 hours.

At 150 °C (Figure 106), the initial shape of the conductivity

curves for stretched and unstretched specimens is in general accord

with the curves shown in Figure 105, but at the lower temperature the

electrical conductivity values remain about the same over an aging span

of 1000 minutes.

The hardness values, before aging of the 3 percent stretched and

unstretched specimens, were 60 and 55 units on the Rockwell B scale,

respectively. Thus, while higher initial values for hardness were

obtained on stretched specimens prior to aging, as compared to

unstretched specimens, the opposite relationship was observed in electrical conductivity measurements. Table VII gives the hardness values

84
attained after aging at different temperatures. It should be noted

that continuous aging, somewhat beyond 6 hours, at temperatures of 150,

170, and 180 °C yielded higher values than those tabulated. After 16

hours of aging at 150, 170, and 180 °C, the average hardnesss values of

the unstretched specimens were found to be 76, 78.5, and 82 HRB, respectively, indicating an increasing trend in hardness values with time.

However, after 30 hours of aging, well beyond the overaging point on

the 180 and 190 °C isotherms, the hardness values for the unstretched

specimens aged at 150, 170, 180, and 190 °C were found to be 79, 80,

77, and 71 HRB, respectively. The increased values at 150 and 170 °C

reveal that overaging has not yet set in at these temperatures in the

specified time period. Furthermore, it is noteworthy that the overaging

phenomenon at 180 and 190 °C present on the hardness isotherms was not

at all apparent in the electrical conductivity.

4. Discussion

The initial decrease in conductivity of the alloy, during the

first few minutes of the aging process at all isothermal holding temperatures between 21 and 190 °C (Figure 103) is apparently due to the

formation of G.P.B. zones. The kinetics of their formation is governed

by the mobility of Cu and Mg atoms which may be enhanced by the presence

of quenched- in vacancies. The interaction is particularly strong with

Mg atoms (47). Under favorable nucleation conditions, these zones

rapidly nucleate in a homogeneous fashion throughout the lattice. The

formation of zones is a thermally activated process; therefore, their

rate of formation increases with increasing temperature and the hardness

of the lattice increases accordingly (48). However the zones rapidly

formed at temperatures above 100 °C cannot persist in the presence of

85
more stable precipitates (35). Consequently, they dissolve back into

the matrix by a diffusional process or become nuclei for the formation

of intermediate, semi-coherent, precipitates of well-defined crystallographic identity.

The initial decrease in the electrical conductivity of the alloy

is due mainly to the appearance of zones which are effective scatterers

of electrons. The zone formation process segregates the Cu and Mg

atoms, as well as the frozen-in vacancies, thus purifying the aluminum

matrix of potential scatterers of electrons. However, the purification

achieved by the formation of zones is not sufficient to compensate for

the opposite effect, namely the effective scattering of electrons by

zones. This very mechanism continues to operate at 21, 35, and 50 °C

(Figure 103) thus contributing to the decrease in measured electrical

conductivity. The sigmoidal behavior is suggestive of a thermally

activated process. The logarithmic aging-time dependence of the conductivity is qualitatively similar to the kinetics in other ternary aluminum

alloys (35), except for the absence of a conductivity minimum in 2024

aluminum alloy. The minimum is generally related to the appearance of

critical-size precipitates (44). Noteworthy is the fact that the

sigmoids (Figure 103) for 21, 35, and 50 °C achieve an identical electrical conductivity value at long times. From the isothermal reaction

kinetics, using these three isotherms, the activation energy of the

process can be calculated. The time to achieve 50 percent of the aging

process is indicated by the intersection of the tangent and each sigmoid

curve at the inflection point. In this manner, the aging time for 50

percent reaction at 21, 35, and 50 °C, are 135, 25, and 7.5 minutes,

respectively. Plotting 1000/T, where T is the absolute temperature of

86
these isotherms, as a function of the aging times for 50 percent reaction,

yields a straight line whose slope is proportional to an activation

energy for the diffusional process governing the kinetics of formation

of both G.P. and G.P.B. zones in 2024 aluminum alloy at relatively low

temperatures. The activation energy obtained by this procedure is 20

kcal/mol. Formation of both zones requires the simultaneous transport

of Cu and Mg atoms as Cu-Mg-vacancy groups by a reshuffling process

(47).

At temperatures above 100 °C, intermediate precipitates, rather

than zones, begin to appear with the approximate composition CuMgA^.

They are predominantly S' with S along the grain boundaries, Figure 75.

The precipitation process of the S' particles is thermally activated.

The process is rather sluggish at 150 °C (Figure 103). Since the

hardness increases at 150 °C, though more slowly than at the higher

temperatures, (Figure 104) it is conjectured that the contribution to

hardness is due to G.P.B. zones, rather than nucleation and growth of

S'. As aging temperature increases, the rate of increase in both conductivity and hardness increases in accordance with the rate of formation

of S' particles. The conductivity at 180 and 190 °C levels off (not

shown in Figure 103) at a value of about 40% IACS. This asymptotic

value was achieved after isothermal holding for 60 hours at 180 °C.

Hardness, Figure 104, increases at a higher rate as aging temperature

increases. However, the 190 and 180 °C isotherms reveal overaging

after 5.5 and 16 hours of aging, respectively.

The increase in electrical conductivity, after the initial decrease,

at aging above 150 °C, can be explained in terms of purification of the

aluminum matrix. Above 150 °C, G.P.B. zones desolve while S' and S

87
precipitates nucleate and grow. In contrast to zones, the relatively

large particles, S' and S precipitates, scatter electrons to a lesser

degree. The net effect is an Increase In conductivity.

Plastic deformation, induced by stretching the 2024 aluminum alloy

prior to age hardening, is meant to have a beneficial effect on the

ultimate mechanical properties. The stretching deformation induces a

high dislocation density, homogeneously distributed. Consequently,

nucleation of S' precipitates can be controlled and made predominantly

homogeneous as was demonstrated in the TEM studies. Figures 105 and

106 show the effect of plastic deformation on hardness and electrical

conductivity of 2024 aluminum alloy aged at 190 and 150 °C, respectively.

The additional hardness of the stretched specimens, about 5 units on

the Rockwell B scale, is preserved during the entire precipitation

hardening process. This indicates that the dislocations are retained

during the aging process. As expected, the increased population of

dislocations in the lattice of the stretched specimens causes the

initial conductivity to be lower than in the unstretched ones. The

dislocations in the stretched matrix provide nucleation sites for S'

precipitates, thus enhancing the precipitation process. Figure 105

depicts this behavior where it is seen that the curve for electrical

conductivity for the stretched specimen crosses over the curve for the

unstretched specimen at a time period of less than 100 minutes.

The results of the present study illustrate that in 2024 aluminum

alloy a given property, such as hardness, can be associated with a wide

range of eddy current conductivities in the final aged condition. The

differing conductivities are a reflection of differences in the microstructure of the material. Hence, a single nondestructive measurement,

88
such as hardness or electrical conductivity alone, does not itself

determine the state of material. However, because eddy current conductivity is sensitive to the alloy microstructure, it could, when used in

combination with hardness and/or ultrasonic measurements, be incorporated

into a sensitive test procedure for detecting variations in microstructure

in 2024 aluminum alloy.

89
VIII. HEAT FLOW-PROPERTY PREDICTIONS

A heat flow model developed earlier for 2219 aluminum alloy (1)

was used here to calculate almost all conceivable heat flow conditions

anticipated during the quench of 2024 aluminum alloy plates from the

solutionizing temperature of 495 °C. The calculated time-temperature

data was then coupled to the C-curves established in section IV. The

variations in properties across different thickness plates for the

worst and the best heat flow conditions were thus predicted.

1 . Heat Flow Model

The heat flow calculations carried out for the various "pre-aging"

treatments during cooling of a flat plate from an initial temperature

of T

q were identical to those previously described (1). They included

the following:

(a) Asymmetric cooling of plates, where heat is withdrawn from

only the top surface for all times.

(b) Symmetric cooling of plates from both top and bottom followed

by an abrupt variation in the heat transfer coefficient on the bottom

surface of the plate, at different times.

Time- temperature data from the computer program was then combined

with equations (3) to (5) for the determination of C-curves using the

values of the constants reported in Table V. The numerical procedure

for the determination of a given property, e.g., yield strength, was as

follows. Equation (4) is integrated, using the calculated timetemperature data and equation (5) for a given position in the plate,

and the quantity is determined. Using the values of c

m and aQ

from Table V, the value of o, in this case yield strength, is established.

These computations are carried out numerically and simultaneously with

the heat flow calculations.

90
Finally, for a given sequence (A or B) and a given heat treatment,

the values of the constants K

3

, and K

5 are identical for all properties

in question, while a , a and K0 differ according to Table V. Under

m o 2

these conditions, it follows from equations (3) to (5) that the quantity

K

2

ln[(a-CT

o )/(am-ao )] remains the same for all properties, e.g. hardness,

tensile strength, and conductivity, for a given sequence. Thus, from

the yield strength results one can readily obtain all the other properties

without further heat flow calculations.

The thermophysical properties used in the calculations were:

Initial temperature T

Q = 495 °C

Water temperature T^ = 40 °C

Thermal conductivity k = 1.2 W/cm K

2 Thermal diffusivity a = 0.5 cm /s

Heat transfer coefficient 10 h = 0.8 W/cm^K

It was found that minimum strength properties predicted were

always at locations near the bottom surface (where quench interruptions

were induced) for the cooling conditions described under (b) above.

The data for the worst properties (e.g. lowest yield strength), were

established using a computer model and the trial and error methodology

described earlier (1).

The predicted “worst possible case" yield strength data for both

"pre-aged" sequences in the T351 and T851 conditions are plotted versus

plate thickness in Figure 107. Noted on the same plot are the ASTM

B209 specifications (which are the same as those given in Military

10 As before (1), this heat transfer coefficient was deduced from

simulation of temperature- time data on the computer and comparison

of same with actual data obtained in laboratory and commercial

practice. This value approximates the normal condition during

water quench from the solution heat treatment temperature.

91
Specifications QQ-A-250/4E and QQ-A-250/29A) for 2024 and 2124 aluminum

alloys, respectively. The 2124 specifications are listed because the

iron and silicon contents of the 0.635 cm (1/4 in.) thick plate, 0.27 wt.%

Fe and 0.11 wt.% Si, are within the 2124 composition specification.

The minimum specification for yield strength in the T351 condition is

quite low and is not available for the 2124 alloy in the T351 condition

or for the complete range of plate thicknesses (up to 6 inches) considered.

The predictions show that under the "worst" heat flow conditions, the

T851 plates subjected to a sequence B "pre-aging" treatment suffer

significant deterioration in yield strength. T851 plates as thin as

1.2 cm (1/2 in.) thick can fall out of specification. This behavior might

be expected when one considers the relative location of the "nose" of

the C-curves in Figures 30 and 34. Unlike the T851 , the minimum yield

strengths for the T351 do not differ significantly for the A and B

sequences. It should be noted, however, that the precision of the T351

predictions (compare the least squares deviations shown in Table V) is

less than the precision of the T851 predictions so that the T351 results

are less reliable than the T851 results.

Figure 108 is similar to Figure 107 except that the "worst possible

case" ultimate tensile strength is shown as a function of plate thickness.

It is seen that the T351B curve can be slightly below the minimum

specification for 2024-T351. Thus a plate of T351 might, according to

Figure 107, meet the yield strength specification yet fail the ultimate

tensile strength specification. As for the yield strength, T851 plates

as thin as 1.2 cm (1/2 in.) can fall out of specification for the "worst

case" of one-sided quench water flow interruption considered here.

92
In Figures 109 and 110, the "worst case" hardness and eddy-current

conductivity respectively, are plotted as a function of plate thickness.

Typical minimum hardness and allowable eddy-current conductivity ranges

from current industrial practice are indicated on these two plots. It

can be seen that, for the T351 temper, these results predict that for

thin plates the hardness and conductivity fall out of range while the

yield strength and ultimate tensile strength are within specification.

93
IX. CONCLUSIONS

1 . As-received plate

1.1 A 15.24 cm (6 inch) thick 2024-T851 aluminum plate was found

to contain moderate variations in composition, hardness and

eddy-current conductivity across its thickness. Macrosegregation present in the original ingot is responsible for the

composition variations. The variations in hardness are

mainly due to changes in cooling rate across the plate during

the quench from solution heat treatment and are probably

influenced to some extent by inhomogeneous mechanical deformation during processing.

2. Sol idifi cation-Segregation Studies

2.1 The phases present in cast 2024 aluminum alloy due to

solidification-microsegregation, determined by metallography,

electron diffraction and x-ray energy dispersive spectroscopy,

were a-Al solid solution, 0-CuA 1 2

, S-CuMgAl

2

, Cu

2

FeAl

7

, Mg

2

Si

,

and orAl(Fe,M)Si where M in this case designates Cu and Mn.

The formation of the latter two phases is a function of alloy

composition and it is favored by higher cooling rates during

solidification.

2.2 Macrosegregation of copper and other alloying additions in DC

cast ingots of 2024 aluminum alloy cannot be completely

eliminated by chill face scalping and subsequent thermomechanical treatment. Macrosegregation does remain in the finished

plate product. However, good scalping practice should maintain

compositions to within specified limits for 2024 with no

deterioration in mechanical properties.

94
2.3 Positive and negative macrosegregation was obtained in a

unidirectional ly cast ingot. The negative segregation was

intentionally introduced by an abrupt cross-section change in

the ingot which resulted in excessive flow of segregated

interdendritic liquid.

2.4 Eddy-current conductivity of cast 2024 aluminum alloy is

approximately inversely related to copper content in a complex

manner. This fact complicates the relationship of conductivity

to mechanical properties used for nondestructive evaluation

of the finished plate product.

2.5 Because of large copper content variation near the chill face

of DC cast ingots, surface hardness and eddy-current conductivity measurements may be very sensitive to scalping depth

in their ability to evaluate the condition of finished alloy

plate.

2.6 Based on a limited number of samples obtained from the laboratory cast ingot designed to cause macrosegregation, the

hardness of properly heat treated 2024-T851 is below specification when the average composition is below 3.1 wt.%Cu and 1.1

wt.%Mg.

3. C-Curves and Nondestructive Evaluation

3.1 The C-curves developed in this report provide a good description

of the effect of time- temperature quench history on the

mechanical and NDE properties of 2024-T351 and 2024-T851

.

3.2 Eddy-current conductivity alone cannot be used as a reliable

predictor of the mechanical properties of 2024-T351 or 2024-

T851. It must be combined with other information such as

95
hardness and yield strength measurements on the same lot

(same ingot or plate) of material.

3.3 For any given thermomechanical treatment of 2024 aluminum

alloy, there will be a large amount of scatter in the final

mechanical properties and NDE measurements, both within a

given lot of material and between different lots of materials.

This scatter is much greater for 2024-T351 than for 2024-T851.

3.4 Comparison of the C-curves for 2024-T351 , 2024-T851 , and

2219-T87* shows that: (1) 2024-T851 is more quench sensitive

than 2024-T351, (2) 2024-T851 is more sensitive to a B-sequence

quench (i.e. "pre-aging" with reheating of the material) than

to an A-sequence quench (i.e. one in which no reheating

occurs), and (3) both tempers of 2024 are more quench sensitive

than 2219-187*.

4. TEM Studies: Relationship Between Microstructure and Properties

4.1 Constituent phase particles in the 0.635 cm thick 2024 aluminum

alloy plate, in agreement with other investigations, consist

of two types; large particles > 1 pm in size retained from

the cast structure and small dispersoid particles < 1 pm.

4.2 The majority of the large particles retained from the cast

structure were found to be the cubic phase a-Al (FeCuMn)Si

.

Other constituent phase particles would probably be present

in 2024 aluminum alloy plate material having a composition

different from that studied here.

4.3 Dispersoid particles are not present in the as-cast ingot and

are formed during thermomechanical processing treatments

subsequent to casting. These particles were determined by

96
means of electron diffraction and energy dispersive x-ray

analysis to be the orthorhombic phase Cu

2

Mn

3

Al

2

Q. They are

insoluble at the solution treatment temperature of 495 °C.

4.4 Quenching from the solution heat treatment temperature (495 °C)

and holding at an intermediate temperature above approximately

300 °C (sequence A "pre-aging" treatment) results in the

heterogeneous nucleation and growth of ©-CuA 1 2

, S-CuMgAl

2

and

S' phases at dispersoid particles and S-CuMgAl

2 precipitates

at grain boundaries.

4.5 An ice water quench from the solution heat treatment temperature followed by aging at a temperature between 300 °C and

495 °C (sequence B "pre-aging" treatment) leads to the formation

of the S' phase heterogeneously nucleated at dispersoid

particles and dislocations and S phase precipitates at grain

boundaries.

4.6 Prolonged "pre-aging" results in the disappearance of all

phases initially induced by "pre-aging" with the exception of

S-CuMgAl

2

which was apparently the equilibrium phase in the

2024 aluminum alloy plate studied.

4.7 Precipitate phases formed during "pre-aging" or alternatively

as a result of an abnormal or slow quench from solution heat

treatment consume Cu and Mg solute, required during final

aging for the development of zones and precipitates, resulting

in the degradation of properties.

4.8 C-curves which indicate the degradation of properties due to

an abnormal quench are also a measure of the amount and type

of second phase constituents formed during quenching.

97
5. Ultrasonic Characterization

5.1 A parabolic relationship was found between hardness and

sound-wave velocity in 2024 aluminum alloy for T4, T351 , and

T851 tempers. Maximum hardness (or yield strength) values

are related to an intermediate, well determined, range of

values of sound velocity for each temper.

5.2 Ultrasonic attenuation decreases as hardness increases. A

linear relationship was found for the range of hardness

values between 60 and 80 HRB.

5.3 "Pre-aging" at 350 °C for 60 minutes yields the greatest

reduction in hardness, and corresponding maxima in sound

velocity and ultrasonic attenuation.

6. Eddy-Current Conductivity vs Hardness During Aging

6.1 Eddy-current conductivity and hardness measurements during

aging of 2024 aluminum alloy show that the two measurements

are not a single valued function of one another.

7. Heat Flow—Property Predictions

7.1 Yield strength, tensile strength, hardness, and eddy-current

conductivity for 2024-T351 and 2024-T851 were calculated from

the appropriate C-curves for interrupted (abnormal) cooling,

in which the heat transfer coefficient at the bottom plate

surface changes from the same value as at the top surface to

a zero value. The "worst case" properties occur near the

bottom surface of the plate. These calculations are subject

to the uncertainties found in the determination of the C-curve

parameters.

98
7.2 For the T851 alloy, sequence B "pre-aging" treatment gives

considerably poorer "worst case" properties than sequence A

for all properties (yield strength, tensile strength, hardness,

and eddy-current conductivity). For example, for a 5.08 cm

(2 in.) thick plate, T851-sequence A has a "worst case" yield

strength of 58.4 ksi compared with 38.4 ksi for the sequence

B "pre-aging" treatment. Under the "worst" heat flow conditions,

T851 plates subjected to sequence B "pre-aging" treatment

suffer significant deterioration in properties and fall below

ASTM specifications.

7.3 For the T351 alloy, there is little difference in yield

strength, tensile strength, and hardness between sequence A

and B "pre-aging" treatments. The sequence B "pre-aging"

treatment gives higher values of the "worst case" conductivity.

Even the "worst case" yield strength lies above minimum

specifications and the "worst case" tensile strength lies

above or only slightly below the minimum specifications.

99
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100
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101
41. M. Rosen, C. L. Friant, S. Fick and E. Horowitz, NDE of the Precipitation

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49. G. Panseri and T. Federighi, Acta Met., 8 , p. 21 7 (1 960).

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Mater. Sci. Engr., !53, p. 191 (1982).

102
TABLE I

Equilibrium Partition Coefficients for

Solidification of Primary a-Aluminum Phase

1

Al-Cu 0.17

Al-Mg 0.30

Al-Mn 0.95

Al-Fe 0.02

Al-Si 0.13

103
TABLE II

Solidification "Path" for Al-Cu-Mg-Mn-Fe-Si

f

s wt% Cu wt% Mg wt% Mn wt% Fe wt% Si

0.0 4.00 1.40 0.65 0.20 0.10

0.1 4.36 1.51 0.65 0.22 0.10

0.2 4.80 1.64 0.66 0.25 0.12

0.3 5.36 1.79 0.66 0.27 0.14

0.4 6.12 2.00 0.67 0.33 0.16

0.5 7.12 2.27 0.68 0.39 0.28

0.6 8.52 2.66 0.68 0.49 0.22

0.7 10.8 3.25 0.69 0.65 0.30

0.8 15.2 4.33 0.70 0.97 0.40

0.85 19.3 5.28 0.72 1.28 0.52

0.90 27.0 7.01 0.73 1.91 0.74

0.91 29.3 7.55 0.73 2.12 0.81

104
TABLE III

Phases Identified in As-Cast 2024 Aluminum Alloy Ingots by Sperry (3)

a-Al Face Centered Cubic

Mg

2

Si Cubic

CuAl

2

Tetragonal

CuMgAl

2

Orthorhombic

MnAlg Orthorhombic

FeAl

3

Orthorhombic

a-Al (Fe,Mn)Si Body Centered Cubic

Cu

2

FeAl^ Tetragonal

105
TABLE IV

Data for sequence A and sequence B heat treatments are used to determine

C-curves for 2024-T851 and 2024-T351 in this report. The range between

the solution heat treatment temperature (495°C) and 110°C was divided

into intervals. First given is the list of the bracketing temperatures

for these intervals. Next, sample numbers are given followed, on the same line, by the Rockwell B hardness, the conductivity in % IACS, the

yield strength (0.2% offset) in ksi, the ultimate tensile strength in ksi,

the % elongation and the % reduction in area, respectively for the fully

processed alloy. For each sample, the following four lines give a list of

times, in seconds (± 0.1 second), spent between each of the temperatures

listed, respectively. For sequence A alloys, the cooling cycle from the

solution heat treatment temperature was performed by a direct transfer from

the solution heat treatment furnace to a salt bath and then to an ice water

quench. For sequence B alloys, a quench into ice water from the solution

heat treatment temperature was followed by an immediate transfer to a salt

bath at elevated temperature and then by another ice water quench.

2024-T851 SEQUENCE A

1 # 496.5 485.6 473.6 461 .

S

456 .6 438 .0 426.0 414.0 402.5 396.5

16 379.6 367.6 355.6 343.6 331 .0 319 .0 307.0 295.0 283.0 270.5

16 258.5 246.6 234.6 221.5 269 .0 197 .0 184.0 172.0 159.0 147.0

3 134.6 122.6 116.6

END

1669 83.2 38.4 .6 .6

• 8 *0 *0 t 0 *0 .0

.0

.0

.0 .0 .0 .0

•0 #0 *0 *0 *0 .6 .0 .0 .0 .0

• 0 *0 *0 #0 • 0 .0 .0 .0 .0 .0

• 0 *0 1010 85.1 38.8 69.3 72.5

•0 .0 .0 .0 *0

12.0

.0

23.0

.0 .0 .0 .0 .0 .0 .0 .0 • 0 .0 .0 .0 .0 .0 .0 .0 .0 .0 *0 .0 .0 .0 .0 .0 .0 .0

1011 81.7 38.4 .0 .0

• 0 .0 .0 • 0 *0

.0 .0

.0 .0 .0 .0 .0

.0 .0 .0 .

0

*0 .0 .0 .0 .0 .0 .0 .

0

.0 .0 .0 .0 .0 .0 .0 .0

• 0 .0 1012 84.9 38.2 68.1 71.9

•0 .0 #0 .0 *0 12.0

.0 22.0

.0 .0 .0 .0 0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

• 0 .0 1015 83.9 38.7 67.9 72.8

•0 .0 .0 .0 .0 13.0

.0 24.0

.0 .6 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 «

0

1016 * 85 . 0

* 38.2 68.6 71.9

.0 .0 .0 .0 .0 11.0

.0 22.0

.0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

•0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0

1C6
1017 84.5 71.5 11.0 20.0

.0 .0

.0 .0

.0 .0

.0 .0

1018 78.9

.3 .1

.1 .3

.0 .0

.0 .0

1019 77.5

.3 .1

.1 .3

.0 .0

.0 .0

1020 78.0

.3 .1

.1 .3

.0 .0

.0 .0

1021 79.8

.3 .1

.1 .3

.0 .0

• 0 .0

1022 81.0

.4 .1 .3 .3 .0 .0

.0 .0

1023 79.7

.4 .1 .3 .3 .0 .0

.0 .0 1024 79.0

.4 .1 .3 .3 .0 .0 .0 .0 1025 80.5

.4 .1 .3 .3 .0 .0

.0 .0

1026 79.3

.4 .1 .3 .3

.0 .0

.0 .0

1027 77.0

.4 .1 .3 .3 .0 .0 .6 .0 1028 77.5

.4 .1 .3 .3 .0 .0

.0 .0 1029 78.8

.4 .1 .3 .3

.0 .0

.0 .0 1031 42.3

.2 .1 .4 .4 .0 .0 .0 .0 1032 43.6

.2 .1 .4 .4 .0 .0

.0 .0

38.4

.0

.0.0

40.6

.1.3

.0 40.0

.1.3

.0

39.8

.1.3

.0

39.7

.1.3

.0

39.2

.2

.6

.0

39.5

.2

.6

.0

39.3

.2

.6.0

39.2

.2.6

.0 39.6

.2

.6

.0 40.0

.2

.6

.0 39.9

.2

.6

.0 39.8

.2

.6

.0 44.2

.1.6.0

43.7

.1.6

.0

67.9

.0 .0

.0 .0

.0 .0

59.0 66.1

.2 .2

.3 .3

.0 .0

.0 .0

.2 .2

.3 .3

.0 .0

.0 .0

.2 .2

.3 .3

.0 .0

59.9 67.0

.2 .2

.3 .3

.0 .0

63.7 68.1

.1 .2

.6 1.2

.0 .0

.0 .0

.1 .2

.6 1.2

.0 .0

.0 .0

.1 .2

.6 1.2

.0 .0

62.8 67.7

.1 .2

.6 1.2

.0 .0

60.9 66.1

.1 .2

.6 1.2

.0 .0

.0 .0

.1 .2

.6 1.2

.0 .6

.0 .0

.1 .2

.6 1.2

.0 .0

59.9 65.2

.1 .2

.6 1.2

.0 .0

.0 .0

.2 .2

.5 1.7

.0 .0 34.8 46.7

.2 .2

.5 1.7

.0 .0

.0 .0

.0 .0

.0 .0

12.0 21.0

.2 .1 .6 .6

.0 .0

.0 .0

.2 .1 .6 .6

.0 .0

.0 .0

.2 .1 .6 .6

.0 .0

13.0 22.0

.2 .1 .6 .6

.0 .0

13.0 23.0

.2 .1 2.0 2.0

.0 .0

.0 .0

.2 .1 2.0 2.0

.0 .0

.0 .0

.2 .1 2.0 2.0

.0 .0 12.0 21.0

.2 .1 2.0 2.0

.0 .0

12.0 22.0

.2 .1 2.6 16.0

.0 .0

.0 .0

.2 .1 2.6 16.0

.0 .0

.0 .0

.2 .1 2.6 16.0

.0 .0

13.0 21.0

.2 .1 2.6 16.0

.0 .0

.0 .0

.2 .4 4.0 194.4

.0 .0 17.0 29.0

.2 .4 4.0 194.4

.0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0

.1 .1 .1 .1 2.0 93.1

. 0 .0 .

0

.1 .1 .1

.1 2.0 93.1

.0 .0 .0

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.0 .0 .1 .1 .0 .0

.0 .0 .0 .0 .0 .0

.0 .1 1341 88.6 38.6 66.0 73.5 .0 1.8 .3 .5 .5 .6 .7

.0 .0 .1 .1 .0 .0 .0 .0 .0 .0 .0 .0

.0 .1 1343 77.3 38.9 62.2 67.6 .0 3.2 .3 .3 .4 .5 .7

.0 .0 .0 .0 .0 .9

.0 .0 .0 .0 .0 .0

.1 1344 76.3 38.8 61.8 67.1 .0 2.3 .3 .3 .4 .5 .7 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.1 END

.4.0

.1

1.4

.0

.9

6.6

.6

.0

1.4

.9.9

6.6

.6

.9

1.8 43.4

.6 .9

.0 .0

.0

1.2 43.4

.0 .0

.6

.6

.0

10.6

.0

.0

10.5

.0

.0

53.2

.0

.0

53.2

.0.0

.6

.0

.1.0

.0

.1.0

.0

.1.0

.0

.1.0

.0

8*84 -T851 SEQUENCE B

TtHP MNQC0

10 496 , 6 488.0 473.0 481 .6 466 .0 438 .0 488.0 414.0 408.6 390.6

10 379.0 387 .0 366.0 343.0 331 .0 319.0 307.0 295.0 233.0 870.6

10 868.6 848 .0 834.0 881 .6 809.6 187.0 184.6 178.0 169.0 147.0

3 134.

CND

8 182 .0 110.0

1009 83.3 38.4 .0 .6 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0

1010 85.1 38.2 69.3 72.5 13.0 83.0

.0 .0 .0 .0 .0 .9 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 1011 81.7 38.4 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 1012 84.9 38.2 68.1 71.9 12.0 22.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 1015 83.9 38.7 67.9 72.8 13.0 24.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 1016 85.0 38.2 68.6 71.9 11.0 22.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 1017 84.5 38.4 67.9 71.5 11.0 20.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0A

.0A

.0 .0 .0 .0 .0 .0 .0 .0

112
1*75 68.8 41.8 51.9 59.9 14.0 23.0

• 0 .0 .0 .0 .0 94.9 8.7 1.2 .8 .6

• 6 • 5 .4 .4 .3 .3 .4 .4 .3 .3

• 5 .7 .6 .5 .5 .5 .2 .5 .2 .2

• 2 • 2 1676 72.0 41.5 .0 .0 .0 .0

.0 .0 .0 .0 .0 94.9 2.7 1.2 .8 .6

• 6 • 5 .4 .4 .3 .3 .4 .4 .3 .3

• 6 .7 .6 .5 .5 .5 .2 .5 .2 .2

• 2 .2

1079 72.3 41.3 54.1 64.1 12.0 22.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

8.7 1.5 .7 .7 .7 .5 .4 .4 .5 .4

.3 1.0 .6 .3 .9 .4 .4 .3 .2 .2

.2 .4

10S0 73.5 40.9 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

2.7 1.5 .7 .7 .7 .5 .4 .4 .5 .4 .3 1.0 .6 .3 .9 .4 .4 .3 .2 .2 .2 .4 1089 56.3 43.2 39.4 52.9 15.0 27.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 12.5

2.0 .9 .8 .6 .5 .3 .4 .3 .3 .3

.1 .2 .2 .1 .2 .1 .1 .1 .1 .1 .1 .1 1098 56.0 43.2 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 12.5

2.0 .9 .8 .6 .5 .3 .4 .3 .3 .3

.1 .2 .2 .1 .2 .1 .1 .1 .1 .1

.1 .1 1091 51.8 43.7 36.5 49.6 15.0 28.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 86.4

1.8 1.4 1.0 2.0 1.8 1.4 1.2 1.2 1.0 .8 .8 .6 .4 .4 .2 .2 .2 .1 .1 .1

.1 .1 1092 50.7 43.8 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 86.4

1.8 1.4 1.0 2.0 1.8 1.4 1.2 1.2 1.0 .8

.8 .6 .4 .4 .2 .2 .2 .1 .1 .1

.1 .1 1096 80.0 39.5 63.2 70.2 23.0 25.0

.2 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 79.8 3.2 1.9 1.2 2.3 2.6 2.6 2.6 2.0

2.0 2.0

1097 80.2 39.5 .0 .0 .0 .0

.2 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 79.8 3.2 1.9 1.2 2.3 2.6 2.6 2.6 2.0

2.0 2.0

1103 63.5 42.7 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 58.4 25.0 1.6

1.2 1.0 1.1 .5 .4 .6 1.2 1.2 1.9 .3

.1 .3

1104 62.5 42.7 42.3 56.4 13.0 28.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 58.4 25.0 1.6

1.2 1.0 1.1 .5 .4 .6 1.2 1.2 1.9 .3

.1 .3

1109 80.0 39.5 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 12.1 1.9 1.9

.9 .9 .4 .4 .4 .4 .4 .4 .4 .4 .3 .3 1110 79.3 39.4 62.9 69.7 12.0 23.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 12.1 1.9 1.9

.9 .9 .4 .4 .4 .4 .4 .4 .4 .4 .3 .3

1111 83.2 38.9 .0 .0 .0 .0

.2 .1 .1 .1 .1 .1 .1 .0 .0 .0

.0 .0 .0 .0 .6 .0 .0 3.5 1.2 1.2

.5 .5 .5 .3 .3 .3 .3 .3 .3 .3

.3 .4

113
1112 82.1 39.2

.2 .1 .1 .9 .9 .9

. S .5 .5

.3 .4

1116 59.3 43.5

.1 .1 .9

.9 .9 72.9

.5 .3 .3

.5 .9

1126 76.8 49.4

.9 .9 .9

.9 .9 .9

.5 .4 .6

.2 .5 1127 77.5 49.3

.9 .9 .9

.9 .9 .9

.5 .4 .6 .2 .5 1128 81.9 38.9

.9 .9 .9

.9 .9 .9

.7 .7 .7

.2 .2 1129 81.5 38.9

.8 .8 .8 .9 .9 .9 .7 .7 .7 .2 .2 1142 79.9 39.8

.9 .0 .0 .0 .0 .0 .0 .0 .0 .8 .5 1143 77.8 39.8

.0 .0 .0

.0 .0 .0

.0 .0 .0

.8 .5 1146 81.5 39.3

.0 .0 .1 .9 .0 .0

.0 .0 .0

1.2 1.2

1147 79.8 39.4

.0 .0 .1 .0 .0 .0 .0 .0 .0

1.2 1.2

1170 78.5 39.6

.1 .1 .0

.0 .0 .9

88.7 190.0 4.9

.4 7.1

1171 78.6 39.9

.1 .1 .0

.0 .0 .0

88.7 100.0 4.9

.4 7.1

1176 78.2 39.8

.1 .9 .0

.9 .0 .9

29.1 252.2 50.6

•3 36.7

1177* 79.0* 39.8

.1 .0 .0

.0 .0 .0 29.1 252.2 50.6

.3 36.7

1178 77.5 39.9

.8 .2 .0

.0 .0 .0

.9 .9 .5 .2 15.1

66.3 71.8 12.9

.1 .1 .1 .0 .9 .0

.3 .3 .3

34.2 49.3 15.0

.0 .0 .9

18.7 2.0 2.0

.3 .3 .3

58.4 67.4 12.0

.0 .8 .9

.0 8.7 3.5

.7 .7 .9

.0 .0 .0

.0 .0 .0 .0 8.7 3.5

.7 .7 .9 66.7 71.8 13.0

.0 .0 .0 .0 .0 .0 .7 .7 .7

.0 .0 .0 .0 .0 .0 .0 .0 .0

.7 .7 .7

.0 .0 .0 .0 .0 .0

.

0

. .0 .0 771.9 4.1 2.1

60.3 69.3 13.0

.0 .0 .0 .0 .0 .0 771.9 4.1 2.1

.0 .0 .0 .0 .0 .0 .0 .0 .0 277.1 4.7 4.6

61.8 69.9 13.0

.0 .0 .0

.0 .0 .0

277.1 4.7 4.6

.0 .0 .0

.9 .0 .9

.9 .9 .9

1.2 .6 .9 60.6 69.0 13.9

.0 .0 .0 .9 .9 .0

1.2 .6 .9

.0 .9 .0

.0 .0 .0

.9 .9 .0

1.4 1.3 .9 58.6 68.4 13.0

.0 .0 .0

.0 .0 .0 1.4 1.3 .9

.0 *0 .0

.0 .0 .0

.0 .0 .0 .5 .5 .4

23.0

.1 .0 .0

.9 3.5 1.2

.3 .3 .3

29.0

.0 .0 .0

.7 .7 .5 .3 .3 .3

22.0

.0 .0 .0 1.5 .8 .7 .9 .8 .7

.0

.0 .0 .0 1.5 .8 .7

.9 .8 .7 23.0

.0 .0 .0

.0 .0 3.1

.7 .3 .2

.9

.0 .0 .0

.0 .0 3.1

.7 .3 .2

.0

.0 .0 .0

.0 .0 .0 1.3 1.2 1.9

26.9

.0 .0 .0

.0 .0 .0 1.3 1.2 1.9

.0

.0 .0 .0 .0 .0 .0 1.2 1.2 1.2

25.0

• & .0 .0 .0 .0 .0

1.2 1.2 1.2

.0

.9 .9 .9 .0 .0 .0 .6 .4 .5

.0

.9 .0 .0

.9 .9 .9 .7 .6 .6 31.0

.0 .0 .0

.0 .0 .9 .7 .6 .6

.0

• 0 .0 .0

.9 19.0 15.5

.4 .3 .3

.9 1.2

.3

.0

.5

.4

.0.5

.8

.0

.5

.8

.0

.9

• 2

.0

.9

.2

.0

.0 1.0

.0

.0 1.0

.0

.0 1.2

.0.0

1.2

.0.0

.4

.0

.9

.4

.0.0

.2

.0

.0.2

.0 2.0

.2

114
1179

.8

.0

.9

.2

1186

1.3

.0

.4

1882

.2

.8

.6

.2

1283

.2

.0

.6

.2

1286

.8

.0

.5.3

1207

.8

.0

.5

.3

1210

.0

3.3

.1

.1 1211

.6

3.3

.1

.1 1238

1.1

3.9

.2

.2

1231

1.1

3.9

.2

.2

1310

.1.0

.5

.2

1311

.1.0

.5.2

1305

.2

.0

.8

.2 1308

.2.0

.8.2

78.0

.2.0

.9 15.1

81.5

.0.0

.4 3.4

81.8

.0.0

.4 3.4

75.4

.2

.0

.4 28.5

76.7

.2

.8

.4 28.5

66.5

.6

.0

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3.9

66.8

.6

.0

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3.9

55.9

.0

1.5

.2

16.8

50.5

.0

1.5

.2

16.8

52.8

.1.8

.2

6.4

54.7

.1.8

.2

6.4

47.8

.8

.0

.4.2

48.0

.8

.0

.4

.2 63.8

.0

.0

.6

.3 64.4

.0

.0.6

.3

40.1

.0

.0

.5 39.5

.0.0

.3

38.8

.0.0

.3 40.1

.2

.0

.4 40.0

.2

.0

.4 41.8

.0

.0

.5 42.1

.0

.0

.5 43.7

.0

1.9

.2

43.4

.0

1.9

.2

43.6

o 0

.6

.2

43.3

.0.6

.2

43.6

.0

.0

.5 43.2

.0

.0

.5 41.9

.0

.0

.3 41.9

.0.0

.3

60.2 68.3

.0 .0

.0 .0

.5 .5

.0 .0

.0 .0

.0 .0

.3 .3 65.1 72.2

.0 .0

.0 .0

.3 .3 58.7 67.9

.2 .1 .0 4.0

.4 .4

.0 .0

.2 .1 .0 4.0

.4 .4

.0 .0

.0 .8

.0 .0 .3 .3 45.2 58.9

.0 .0

.0 .0

.3 .3

38.9 50.9

.0 .0

.7 .3

.2 .2

.0 .0

.0 .0

.7 .3

.2 .2

38.1 50.0

.0 .0

.5 .4

.2 .2

.0 .0

.0 .0

.5 .4

.2 .2

34.4 48.7

.0 .0

.0 434.7

.5 .3

34.5 49.4

.0 .0

.0 434.7

.5 .3

43.9 57.6

.0 .0 .0 5.2

.3 .3

44.4 57.7

.0 .0 .0 5.2

.3 .3

12.0 25.0

.0 .0

.0 .6

.4 .4

.0 .0

.0 .0 1.2 2.2

.2 .2 13.0 26.0

.0 .0 1.2 2.2

.2 .2 12.0 25.0

.0 .0

6.8 2.0

.3 .4

.0 .0

.0 .0

6.8 2.0

.3 .4

.0 .0

.0 .0

7.5 26.4

.3 .3

13.0 26.0

.0 .0

7.5 26.4

.3 .3

14.0 27.0

.1 .1 .4 .3

.1 .2

.0 .9

.1 .1 .4 .3

.1 .2

15.0 27.0

.0 .0

.4 .3

.1 .1

.0 .0

.0 .0

.4 .3

.1 .1

.0 .0

.0 .0

15.4 2.0

.3 .2

.0 .0

.0 .0

15.4 2.0

.3 .2

.0 .0

.0 .0

24.0 2.6

.3 .3

.0 .0

.0 .0

24.0 2.6

.3 .3

.0 .0 .0 19.0 15.5 2.0

.3 .3 .2

.0 .0 * 0

1.9 1.0 .5 .3 .2 .2

.0 .0 .

0

1.9 1.0 .5 .3 .2 .2

• 0 .0 • 0

1.1 .7 .8 .3 *2 .3

.0 .0 .0

1.1 .7 .8 .3 .2 .3

.0 .0 .

0

1.0 .9 .7

.3 .3 .5

.0 .0 .0

1.0 .9 .7

.3 .3 .5

6.2 104.3 79.0

.4 .4 .4

.1 .0 .0

6.2 104.3 79.0

.4 .4 .4

.1 .0 .0

.0 181.4 310.0

.4 .2 .3 .2 .2 .2

.0 181.4 310.0

.4 .2 .3

.2 .2 .2

0 .0 .0

.0 .0 .8

.3 .2 .2

.0 .0 .0

.0 .0 .8 .3 .2 .2

.0 .0 .0

2.1 1.5 .7

.3 .3 .3

.0 .0 .

0

2.1 1.5 .7

.3 .3 .3

115
1386 68.6 41.5 48.4 61.8 .9 .9 13.7 5.9 .9 .9 .9 .9 .9 .9 .9 .9 .9 .9 .5 19.4 2.4 1.1 .9 .9 • 6 .4

.6 •8 .3 .3 .3 .2 .2 .2 9 8 4 3

.6 .3 1387 67.8 41.3 59.8 68.3 .9 .9 13.7 6.9 .9 .9 .9 .9 .9 .9 4 0 .9 .6 .9 .5 19.4 8.4 1.1 .9 .9 .6 .4 .6 .8 .3 .3 .3 .2 .8 .2 .8 . 3

.6 .3 1339 54.6 48.7 37.6 58.5 .9 .9 8.6 1.9 .9 .9 .9 .9 .9 .9 .9 .9 .6 .1 81.6 11.1 1.7 .9 .7 .8 <6 4 6 .4m

.8A

.3 .3 .8 .3 .8 .5 .4 .3

9 9

1331

• 1

66.8 48.6 37.8 58.8 .9 .9 8.6 1.9 .9 .9 .9 .9 4 6 .9 .9 .9 ,6 .1 81.6 11.1 1.7 .9 .7 .8 46 4 6

.4 .2 .3 .3 .8 .3 .8 .5 .4 .3

.3 .4

1334 43.6 44.9 38.4 47.9 .9 .9 3.6 .9 .9 .9 .9 .9 .9 .9 .9 .9

.8 .9 848.6 849.7 1.4 1.9 .8 .8 0 fe 4 S

.3 .3 .4 .2 .2 .2 .3 .3 .2 .4

.1 .3

1336 43.8 43.8 32.6 47.3 .9 .9

3.5 .9 .9 .9 .9 .9 .9 .9 4 0 • 0

.• .9 242.5 249.7 1.4 1.9 .8 .8 • 6 .5 .3 .3 .4 .2 .2 .2 .3 .3 .2 .4

.1 .3

1347 55.2 42.9 38.4 51.8 .9 .9 .2 .9 .9 .9 .9 .9 .9 .9 39.6 3.9

1.6 .7 .7 .7 .8 .8 .6 .2 .1 .1

.1 .1 .1 .1 .1 .1 .1 .1 .2 • 2

.2 .4 1348 55.7 42.8 39.4 52.5 .9 .9

.2 .9 .9 .9 .9 .9 .9 .9 39.6 3.9

1.6 .7 .7 .7 .8 .8 .6 .2 .1 .1 .1 .1 .1 .1 .1 .1 .1 .1 *Z • 2 .2 .4 END

2024-T351 SEQUENCE A

Igy RANGES

19 496.5 48S.9 473.9 461.5 459.9 438.9 426.9 414.0 492.5 399.5

19 379.9 367.9 355.9 343.9 331.9 319.9 397.9 295.9 283.0 279.5

19 258.5 246.9 234.9 221.5 299.9 197.9 184.9

3 134.9 122.9 119.9

END

172.0 159.0 147.0

1995 76.7 28.4 .9 .9 .9 .9

.9 .9 .9 .9 .9 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0

.9 .0

.0 .0 .0 1996 79.4 29.5 55.1 68.9 23.0 26.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

•0 .0 .0 .0 .0 .0 .0 .0 .3 .0

.0 .0 .0 .0 .0 .0 .0

.0 .0

.0 .0 .0 1997 79.9 29.0 55.5 69.9 22.9 28.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .9 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 1008 76.0 28.7 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0

.0 .0 .0 1013 79.1 29.4 54.2 68.0 24.0 28.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 •

0

.0 .0 .0

•0 .0 .0 .0 .0 .0 .0

.0 .0

.0 .0 .0

116
1914 79.3 29.2

.9 .8 .0

.0 .0 .0

.0 .0 .0

.0 .0

1036 76.1 31.4

.1 .1 .2 2.4 6.0 .0

• 0 .0 .0 1037*

0

7S . 0* 0

31.3

.1 .1 .2

2.4 6.0 .0 .0 .0 .0

.0 .0

1065 67.5 35.6

.1 .2 .2

.3 .3 .5 .0 .0 .0

.0 .0 1066 66.0 35.2

.1 .2 .2

.3 .3 .5

.0 .0 .0

.0 .9

1069 69.4 32.5

.1 .3 .4

.0 .0 .0

.0 .0 .0

.0 .0

1070 69.0 31.8

.1 .3 .4

.0 .0 .0

.0 .0 .0

.0 .0 1073 77.1 29.6

.6 .4 1.4

.0 .0 .0

.0 .0 .0

.0 .0 1074 76.5 29.7

.6 .4 1.4

.0 .0 .0

.0 .0 .9

.0 .0 1101 63.0 36.5

.2 .2 .2

.4 .4 .6 .0 .0 .0 .0 .0 1102 65.0 36.9

.2 .2 .2

.4 .4 .6 .0 .0 .0 .0 .0 1134 71.8 32.6

.3 .3 .3

2.6 8.9 3.3

.0 .0 .0 .0 .1 1135 71.5 32.6

.3 .3 .3 2.6 8.9 3.8

.0 .0 .0

.0 .1 1156 81.7 33.1

.2 .2 .2

.2 .2 .2

1.9 4.6 727.0

.3 .3

1157 81.5 32.8

.2 .2 .2

.2 .2 .2

1.9 4.6 727.0

.3 .3

54.8 68.4 24.0

.0 .0 .

0

.0 .0 .

0

.0 .0 • 0

53.8 67.0 20.0

.2 .5 .4 .0 .0 .0

.0 .0 .0 .0 .0 .0 .2 .5 .4 .0 .0 .0 .0 .0 .0 48.1 60.6 16.0

.2 .1 .2

.8 1.4 2.2

.0 .0 .0

.0 .0 .0

.2 .1 .2

.8 1.4 2.2

.0 .0 .0

56.5 64.2 11.0

.4 .6 .6

.0 .9 .9

.0 .0 .0 .0 .0 .0

.4 .6 .6 .0 .0 .0 .0 .0 .0

55.2 68.6 22.0

2.6 99.0 .0

.0 .0 .0

.0 .0 .0

.0 .0 .0 2.6 99.0 .0

.0 .0 .0

.0 .0 .0

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.6 1.2 2.9

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.0 .0 .0 51.6 65.2 19.0

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.0 .0 .0

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.0 .0 .0

.2 .2 .2

.4 .4 .4

.3 .3 .3 60.5 73.2 21.0

.2 .2 .2

.4 .4 .4 .3 .3 .3

28.0

.0 .0 .

0

.0 .0 . .0 .0 .0 20.0

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.0

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94.0 .0 .0 .0 .0 .0

.0

.1 .3 .2

94.0 .0 .0

.0 .0 .9

19.0* 1.2 52.0 50.8

.0 .0 .0

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.0

1.2 52.0 50.8

.0 .0 .0 .0 .0 .0 28.0

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.2 .2 .4 104.5 .0 .0 .0 .0 .0 23.0

.5 .5 1.0

.0 .0 .0

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.0

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.0 .0 .0

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.0

• 2 .2 .2

.4 .6 .6

.3 .3 .3

.0

.2 .2 .2

.4 .6 .6 .3 .3 .3

.0

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.0

1.2

.0

.0

1.2

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.9

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117
1168 75.7 35.3

.7 .6 .2 •3 .3 .3 315.9 68.2 .2

• • .1 1161 76.8 35.2

•7 .6 .2 .3 .3 .3

316.9 68.2 .2

.8 .1 1162 88.5 32.2

.1 .1 .1 .2 .3 .3

2.8 5.8 337.9

.1 .1 1165 81.8 31.9

.1 .1 .1 .2 .3 .3 2.8 5.8 337.9

.1 .1 1168 78.5 33.8

2.8 .6 .1 .2 .3 .3 132.3 56.3 .0 .8 .1 1169 78.8 33.5

2.8 .6 .1 .2 .3 .3 132.3 56.3 .8 .8 .1 1184 73.2 33.1

.3 .3 .3

.5 .5 .8

.8 .8 .8 .8 .8 1185 71.3 33.1

.3 .3 .3

.5 .5 .8 .8 .8 .8 .8 .8 1192 73.9 32.8

.1 .1 .1 .7 .7 1.8

.0 .0 .8

.8 .8 1193 73.2 32.3

.1 .1 .1 .7 .7 1.8

.3 .8 .8 .8 .8 1196 66.8 35.1

1.5 .9 .3 .9 1.8 1.3

.8 .8 .8

.8 .1 1197 65.8 34.9

1.5 .9 .3 .9 1.8 1.3

.8 .8 .8 .8 .1 1288 75.8 38.9

1.8 .6 .3 .6 .8 1.3

.8 .8 .8

.8 .1 1281 75.7 31.8

1.8 .6 .3

.6 .8 1.3

.8 .8 .8 .8 .1

.0 . 8

.8 .1 .5 .5 .8 .8 59.4 65.8

.8 .1 .5 .5

.8 .8

.8 .8

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.4 .4

.1 .1 59.3 72.5

.2 .2

.4 .4

.1 .1

.8 .8

.1 .3

.3 .3

.8 .8 61.5 78.8

.1 .3

.3 .3

.8 .8

.8 .8 .3 .3 1.0 2.6

.8 .8

51.3 64.8

.3 .3 1.8 2.6

.8 .8 52.5 66.8

.1 .2 1.4 12.6

.8 .8

.8 .8

.1 .2

1.4 12.6

.0 .8 47.1 60.3

.3 .3 28.6 5.2

.0 .8

.8 .8 .3 .3 28.6 5.2

.8 .8 53.9 67.7

.2 .2 3.2 1.4

.8 .8 .8 .8 .2 .2

3.2 1.4

.8 .8

.8 .8

.1 .2

.6 .9 .8 .8 16.8 23.8

.1 .2

.6 .9 .8 .8

.8 .0

.1 .1 .3 .7

.1 .1 19.8 19.8

.1 .1 .3 .7

.1 .1

.8 .0

.2 .2

.6 .5

.0 .8 18.8 21.0

.2 .2

.6 .5

.8 .0

.0 .8 .3 .3 25.2 8.3

.0 .8 11.0 16.8

.3 .3 25.2 8.3

.0 .0 10.8 15.8

.1 .2 2.1 .3

.0 .0

.0 .8

.1 .2 2.1 .3

.8 .0 17.0 23.8

.2 .4

.8 .8

.0 .8 .0 .0

.2 .4

.8 .8

.8 .8 28.8 25.8

.2 .3 1.1 1.1

.0 .8

.0 .0

.2 .3

1.1 1.1

.0 .0

.2 .2 .4 1.1 1.6 9.9

.8 .0 .0

.2 .2 .4 1.1 1.6 9.9

.0 .0 .0

.2 .4 .3

.7 1.0 1.5

.1 .1 .1

.2 .4 .3

.7 1.8 1.5

.1 .1 .1

.2 .2 .2

.8 1.8 5.7

.0 .0 .0

.2 .2 .2

.8 1.8 5.7

.8 .0 .8

.3 .3 .4

.1 .0 .0

.0 .0 .0

.3 .3 .4

.1 .0 .0 .0 .0 .0

•3 .2 .2

.1 .0 .0

.0 .0 .0

.3 .2 .2

.1 .0 .0

. 0 .0 .0

.5 .6 .6

.1 .0 .8

.0 .0 .0

.5 .6 .6

.1 .0 .0 .0 .0 .0

.3 .3 .3

.1 .0 .8

.8 .0 .8

118
1216 66.1 33.3

.3 .5 .6

.6 .6 .6

.6 .6 .6

.1 .1 1217 67.2 33.3

.3 .5 .6 .6 .6 .6

.6 .0 .0

.1 .1 1226 59.5 35.3

1.2 .6 .5

.1 .1 .1 .0 .0 .0

.0 .1 1227 64.2 34.2

1.2 .6 .5

.1 .1 .1 .0 .0 .0 .0 .1 1301 36.8 41.3

3.5 2.9 .1 .5 .4 .4

.1 .0 .0

.1 .1 1304 37.8 41.6

3.5 2.9 .1 .5 .4 .4

.1 .0 .0

.1 .1 1314 68.8 33.5

1.0 1.1 1.0

2.1 8.8 6.5

.0 .0 .0

.1 .1 1315 68.7 33.4

1.0 1.1 1.0

2.1 8.8 6.5

.0 .0 .0

.1 .1 1318 49.3 37.6

1.5 1.0 .3

1.1 2.2 3S.8

.0 .0 .0

.0 .1 1321 50.0 37.5

1.5 1.0 .3

1.1 2.2 35.8

.0 .0 .0 .0 .1 1322 38.4 40.3

1.7 .4 .1 .9 2.2 346.5

.0 .0 .0

47.4 61.5 14.0

1.3 2.6 3.2

.0 .0 .0

.0 .0 .0

.0 .0 .0

1.3 2.6 3.2

.0 .0 .0

.0 .0 .0

43.2 56.3 10.0

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.0 .0 .0

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39.4 44.1 .0

.1 .6 .1 .4 1.2 4.2

.0 .0 .0 39.9 44.2 .0

.1 .6 .1 .4 1.2 4.2

.0 .0 .0

52.1 62.2 .0

1.0 .6 .3 .3 .2 .1 .0 .0 .0 51.6 62.3 .0

1.0 .6 .3 .3 .2 .1 .0 .0 .0 39.7 51.0 .0 .3 .3 .4 47.5 12.7 .0

.0 .0 .0

39.4 50.9 .0

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47.5 12.7 .0

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36.5 43.8 .0

.1 .3 .4

148.5 .1 .1 .0 .0 .0

7.0

6.4 18.5 92.3

.0 .0 .0

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.0

6.4 18.5 92.3

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5.0

1.2 6.0 443.1

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.01.2 6.0 443.1

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.0 .0 .0

.0

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490.2 20.1 1.6

.0 .0 .0

.0

.1 .3 .3

490.2 20.1 1.6

.0 .0 .0

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.0 .0 .0

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.0

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.0 .0 .0

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.0 .0 .0

• 0 .0 .0 .0

.4 .5 .6 .0 .0 .0

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.0

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.0 .0 .0

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77.8

.0

.1

77.8

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48.3

.0.0

48.3

.0

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.7

.6

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.7

.6

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1.1

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.1

1.1

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.1

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119
1326

1.7

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.9

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1339

1.2

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.9 1349

1.2

.9

.9

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2.2

.0

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2.2

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.0 END

36.6

.4

2.2

.9

.1 77.8

.3

.0

.0

.1 78.1

.3

.0

.0

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.3

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.1 70.1

.3

.0

.0

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49.3 37.1

.1 .1 346.5 148.5

.9 .9

29.6 57.9

.5 .5

.1 .1 .0 .6

29.6 57.7

.5 .5

.1 .1 .0 .0 31.6 51.3

.3 .4

.0 .0

.0 .0 31.9 50.6

.3 .4

.0 .0

.0 .0

43.9

.3

.1.0 69.0

.6

.0

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68.7

.6

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.5

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.4 .4

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.0 .9

.0 .0

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.0 .0

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.0 .0

.0 .0

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.0 .0 .0 .0

.0 .0

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.0 .0

.0 .0

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.0 .0 .0 .0

6.5 10.5

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.0 .0

6.5 10.5

.0 .0

.0 .0

43.4 53.2

.0 .0

.0 .0

43.4 53.2

.0 .0

.0 .0

.6

.0

.0

.1.0

.0

.1.0

.0

.1.0.0

.1.0

.0

2024-T351 SEQUENCE B

TEMP RANGES

19 496 .5 485 .0 473.0 461 .5 450 .0 438 .9 426 .0 414 .0 462. 5 390

10 379 .9 367 .9 355.0 343 .0 331 .0 319 .0 307 .0 295 .0 283. 0 270

19 258 .5 246 .0 234.9 221 .5 209 .0 197 .0 184 .0 172 .0 159. 0 147

3 134 .0 122 .0 110.0

END

1095 76.7 28.4 .0 .0 .0 .0

.9 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .9 .0 .0 .0 .0 .0 .0 .0 .0 .0

1006 79.4 29.5 55.1 68.9 23.0 26.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .9 .0 1007 79.9 29.0 55.5 69.9 22.0 28.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 1008 , 76.0 28.7 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 1013 79.1 29.4 54.2 68.0 24.0 28.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 1014 79.3 29.2 54.8 68.4 24.0 28.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 1077 69.8 34.3 48.3 62.6 19.0 20.0

.0 .0 .0 .0 .0 94.9 2.7 1.2 .8 .6

.6 .5 .4 .4 .3 .3 .4 .4 .3 .3

.5 .7 .6 .5 .5 .5 .2 .5 .2 .2

.2 .2

120
1078 70.0 34.4 .0 .0 .0 .0 .0 .0 .0 .0 .0 94.9 3.7 1.3 .8 .6 6 • 5 .4 .4 .3 .3 .4 .4 .3 .3

5 .7 .6 .5 .5 .5 .3 .5 .3 .3 .a .3

1081 76.3 31.9 55.3 67.3 19.0 34.0

*0 .0 .0 .0 .0 .0 .0 .0 .0 .0

3.7 1.5 .7 .7 .7 .5 .4 .4 .5 .4 .3 1.0 .6 .3 .9 .4 .4 .3 .3 .3

.3 .4

1083 76.0 31.7 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

3.7 1.5 .7 .7 .7 .5 .4 .4 .5 .4 .3 1.0 .6 .3 .9 .4 .4 .3 .3 .3

.3 .4

1087 64.9 36.4 47.7 59.3 9.0 13.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 13.5

3.0 .9 .8 .6 .5 .3 .4 .3 .3 .3

.1 .3 .3 .1 .3 .1 .1 .1 .1 .1 .1 .1 1088 65.0 36.3 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 13.5

3.0 .9 .8 .6 .5 .3 .4 .3 .3 .3

.1 .3 .3 .1 .3 .1 .1 .1 .1 .1 .1 .1 1093 59.5 37.7 43.4 56.1 10.0 8.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 86.4

1.8 1.4 1.0 3.0 1.8 1.4 1.3 1.3 1.0 .8 .8 .6 .4 .4 .3 .3 .3 .1 .1 .1 .1 .1 1094 56.3 37.6 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 86.4

1.8 1.4 1.0 3.0 1.8 1.4 1.3 1.3 1.0 .8

.8 .6 .4 .4 .3 .3 .3 .1 .1 .1 .1 .1 1095 77.5 31.1 .0 .0 .0 .0

.3 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 79.8 3.3 1.9 1.3 3.3 3.6 3.6 3.6 3.0

3.0 3.0

1098 77.8 30.7 55.3 70.3 33.0 37.0

.3 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 79.8 3.3 1.9 1.3 3.3 3.6 3.6 3.6 3.0

3.0 3.0

1105 70.3 36.3 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 58.4 35.0 1.6

1.3 1.0 1.1 .5 .4 .6 1.3 1.3 1.9 .3

.1 .3 1105 69.0 36.4 53.1 63.3 16.0 30.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 58.4 35.0 1.6

1.3 1.0 1.1 .5 .4 .6 1.3 1.3 1.9 .3

.1 .3

1107 80.0 30.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 13.1 1.9 1.9

.9 .9 .4 .4 .4 .4 .4 .4 .4 .4 .3 .3

1108 77.8 39.9 57.0 69.7 31.0 33.0

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 .0 13.1 1.9 1.9

.9 .9 .4 .4 .4 .4 .4 .4 .4 .4 .3 .3

1113 77.7 39.6 .0 .0 .0 .0 .3 .1 .1 .1 .1 .1 .1 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 3.5 1.3 1.3

.5 .5 .5 .3 .3 .3 .3 .3 .3 .3

.3 .4 1114 79.0 39.3 59.0 70.3 31.0 38.0

.3 .1 .1 .1 .1 .1 .1 .0 .0 .0

.0 .0 .0 .0 .0 .0 .0 3.5 1.3 1.3

.5 .5 .5 .3 .3 .3 .3 .3 .3 .3

.3 .4

121
1115 54.7 38 .

B

.8 .8

.1 .1 .8 .8 .8

.8 .8 78.8 18.7 8.8

.5 .3 .3 .3 .3

.5 .9 1118 54.6 39.8 44.7 58.8

.1 .1 .8 .8 .8 .8 .8 78.8 18.7 8.8

.5 .3 .3 .3 .3 .5 .9 1184 77.8 31.0 56.4 68.6

.8 .0 .0 .0 .0 .0 .0 .0 .0 8.7

.5 .4 .6 .7 .7

.8 .5 1185 77.5 30.7 .0 .0 .0 .0 .0 .6 .0 .0 .0 .0 .6 8.7

.5 .4 .6 .7 .7

.8 .5 1130 78.0 89.7 54.8 68.9

.0 .0 .0 .0 .0 .0 .0 .0 .0 .0

.7 .7 .7 .7 .7 .8 .8 1131 78.0 30.1 .0 .0

.0 .0 .0 .0 .0

.0 .0 .0 .0 .0

.7 .7 .7 .7 .7 .8 .8 1144 80.0 31.0 58.9 78.9

.0 .0 .0 .0 .0

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.0 .0 .0 771.9 4.1

.8 .5 1145 80.5 31.5 .0 .0 .0 .0 .0 .0 .0

.0 .0 .0 .0 .0

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.8 .5 1148 78.0 30.8 .0 .0 .0 .0 .1 .0 .0

.0 .0 .0 .0 .8

.8 .0 .0 877.1 4.7

1.8 1.8

1149 83.4 30.8 65.0 73.1

.0 .0 .1 .0 .0 •0 .0 .0 .0 .0 .0 .0 .0 877.1 4.7

1.8 1.8

1178 88.5 31.6 .0 .0

•1 .1 .0 .0 .0 .0 .0 .0 .0 .0 88.7 180.0 4.9 1.8 .6

.4 7.1

1173 81.3 31.5 68.1 78.6

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88.7 188.0 4.9 1.8 .6 .4 7.1

1174 83.8 38.7 .9 .0

.1 .0 .0 .0 .0

• 0 .0 .0 .0 .0 89.1 858.8 50.6 1.4 1.3

.3 36.7

1175 88.8 38.8 64.8 73.8

.1 .0 .0 .8 .0 .0 .0 .0 .0 .0 89.1 858.8 50.6 1.4 1.3

.3 36.7

1180 79.7 30.9 .0 .0

.8 .8 .0 .0 .0

.

0

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0

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.0 .0 8.0 .7

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.0 .0 3.5 1.5

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88.0 85.0

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80.0 84.0

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122

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.8

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.7

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.8

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.0

.0

.6

.0 15.5

.3

.0

.5

.4

.0.5

.4

.0

.5

.8

.0

.5

.8

.0

.9

.8

.0.9

.8

.0

.0 1.0

.0

.0 1.0

.0.0 1.8

.0.0 1.8

.0

.0

.4

.0

.0

.4

.0.0

.8

.0

.0.8

.0 8.0

.8
1181 86.3 30.8 58.1 69.3

.8 .2 .0 .0 .0 .0 .0 .0 .0 .0

.9 .9 .5 .5 .5 .2 15.1

1188 78.7 29.5 .0 .0 1.3 .0 .0 .0 .0

.0 .0 .0 .0 .0

.4 .4 .3 .3 .3

.2 3.4

1189 76.8 29.4 58.0 71.8

1.3 .0 .0 .0 .0

.0 .0 .0 .0 .0

.4 .4 .3 .3 .3

.2 3.4

1204 * 77 . 5

* 31.5 57.1 68.9

.2 .2 .2 .2 .1 .0 .0 .0 .0 4.0

.6 .4 .4 .4 .4

.2 28.5

1205 76.7 31.0 .0 .0

.2 .2 .2 .2 .1 .0 .0 .0 .0 4.0

.6 .4 .4 .4 .4 .2 28.5

1208 * 72 . 3

* 35.5 54.5 63.9

.8 .6 .0 .0 .0

.0 .0 .0 .0 .0

.5 .4 .5 .3 .3

.3 3.9

1209 71.5 34.7 .0 .0

.3 .6 .0 .0 .0

.0 .0 .0 .0 .0

.5 .4 .5 .3 .3

.3 3.9

1212 59.0 37.9 42.2 55.8

.0 .0 .0 .0 .0

3.3 1.5 1.9 .7 .3

.1 .2 .2 .2 .2

.1 16.8

1213 S7.2 37.8 .6 .0

.0 .0 .0 .0 *0

3.3 1.5 1.9 .7 .3

.1 .2 .2 .2 .2

.1 16.8

1232 57.3 38.3 41.9 54.1

1.1 .1 .0 .0 .0

3.9 .8 .6 .5 .4 .2 .2 .2 .2 .2

.2 6.4

1233 53.7 38.0 .0 .0

1.1 .1 .0 .0 .0

3.9 .8 .6 .5 .4 .2 .2 .2 .2 .2

.2 6.4

1312 * 54 . 7

* 39.8 51.5 53.6

•1 .0 .0 .0 .0

.0 .0 .0 .0 434.7

.5 .4 .5 .5 .3

.2 .2

1313 55.8 39.7 49.7 53.2

•1 .0 .0 .0 .0

.0 .0 .0 .0 434.7

.5 .4 .5 .5 .3

.2 .2 1306 68.4 35.5 56.3 61.2

.2 .0 .0 .0 .0

.0 .0 .0 .0 5.2

.8 .6 .3 .3 .3

.2 .3 1307 68.2 35.4 55.8 61.3

.2 .0 .0 .0 .0

.0 .0 .0 .0 5.2

.8 .6 .3 .3 .3

.2 .3

21.0 24.0

.0 .0

.0 .0

.4 .4

.0 .0

.0 .0

1.2 2.2

.2 .2 22.0 27.0

.0 .0 1.2 2.2

.2 .2 20.0 26.0

.0 .0

6.8 2.0

.3 .4

.0 .0 .0 .0 6.8 2.0

.3 .4 15.0 22.0

.0 .0 7.5 26.4

.3 .3

.0 .0

.0 .0

7.5 26.4

.3 .3

11.0 15.0

.1 .1 .4 .3

.1 .2

.0 .0

.1 .1 .4 .3

.1 .2

15.0 20.0

.0 .0

.4 .3

.1 .1

.0 .0

.0 .0

.4 .3

.1 .1

.0 .0

.0 .0

15.4 2.0

.3 .2

.0 .0

.0 .0 15.4 2.0

.3 .2

.0 .6

.0 .0 24.0 2.6

.3 .3

.0 .0

.0 .0 24.0 2.6

.3 .3

.0 .0 .

0

19.0 15.5 2.0

.3 .3 .2

.0 .0 .0 1.9 1.0 .5

.3 .2 .2

.0 .0 .0 1.9 1.0 .5

•3 .2 .2

.0 .0 .

0

1.1 .7 .8

• 3 .2 .3

.0 .0 .0 1.1 .7 .8

.3 .2 .3

.0 .0 .0 1.0 .9 .7

• 3 .3 .5

.0 .0 .0

1.0 .9 .7

.3 .3 .5

6.2 104.3 79.0

.4 .4 .4

.1 .0 .0

6.2 104.3 79.0

.4 .4 .4

.1 .0 .0

.0 181.4 310.0

.4 .2 .3

.2 .2 .2

.0 181.4 310.0

.4 .2 .3

.2 .2 .2

.0 .0 .0

.0 .0 .8

.3 .2 .2

.0 .0 .0 .0 .0 .8 .3 .2 .2

.0 .0 .0 2.1 1.5 .7

.3 .3 .3

.0 .0 .0 2.1 1.5 .7 .3 .3 .3

123
1328 71.7 33.5 54.1 64.2 0 0

13.7 5.0

.0 .0

.5 .2

.5 .3

1329 73.0

13.7 5.0

.0 .0

.5 .2 .5 .3 1332 61.3

2.6 1.0

.0 .1 .4 .2 .3 .4 1333 61.8

2.6 1.9

.0 .1 .4 .2

.3 .4

1336 53.4

3.6 .0

.0 .0

.3 .3

.1 .3

1337 54.0

.1 .3

1346 61.0

.2 .0

1.6 .7

.1 .1 .2 .4 1349 63.4

.2 .0

1.6 .7

.1 .1 .2 .4

.0 .0 .0

.5 10.4 2.4

.3 .3 .3 33.6 56.0 65.0

.0 .0 .0

.5 10.4 2.4

.3 .3 .3

37.4 52.8 57.0

.0 .0 .0

21.5 11.1 1.7

.3 .3 .2 37.5 51.5 57.3

.0 .0 .0

21.5 11.1 1.7

.3 .3 .2

39.5 48.6 52.2

.0 .0 .0

242.5 249.7 1.4

.4 .2 .2

39.3 49.2 52.5

.0 .0 .0

242.5 249.7 1.4

.4 .2 .2 .0 .0 .0

.0 .0 .0

.7 .7 .8

.1 .1 .1 .0 47.3 58.6

.0 .0 .0

.7 .7 .8

.1 .1 .1

.0 .0 1.1 .9

.2 .2

.0 .0

.0 .0 1.1 .9 .2 .2 .0 .0

.0 .0

.9 .7 .3 .2

.0 .0 .0 .0 .9 .7

.3 .2

.2 .3

.0 .0

.0 .0

1.0 .8

.2 .3 .0 .0

.0 .0

.8 .6

.1 .1

.0 .0 .0 .0

.8 .6

.1 .1

.0 .0 .0

.9 .6 .4 .2 .2 .3

.0 .0 .

0

.9 .6 .4

.2 .2 .3

.0 .0 .0 .8 .5 .6 .5 .4 .3

.0 .0 .6 .8 .5 .6 .5 .4 .3

.0 .0 .0

.8 .6 .5

.3 .2 .4

.0 .0 .0

.8 .6 .5 .3 .2 .4

.0 30.6 3.0

.2 .1 .1

.1 .2 .2

.0 30.6 3.0

.2 .1 .1

.1 .2 .2

124
TABLE V

C-Curve Parameters For Aluminum Alloy 2024, Kg For All

Cases is Taken Equal to 32,000 cal/mol.

Temper Sequence Property om °0

k

2

, ns

K

3

Cal /Mol

K

4

Kelvins

Est

.

Std. Dev.

T851 A U.T.S.

ksi

71.2 38.5 14.9 125 750 2.0

A Y.S.

ksi

66.6 27.8 12.1 125 750 2.4

A HRB 82.5 22.0 23.1 125 750 2.8

A %IACS 38.7 44.5 11.3 125 750 0.55

T851 B U.T.S.

ksi

71.2 47.0 25.0 45 750 2.6

B Y.S.

ksi

66.6 34.0 17.2 45 750 2.7

B HRB 82.5 43.7 32.6 45 750 3.0

B %IACS 38.7 44.0 18.2 45 750 0.49

T351 A U.T.S.

ksi

68.5 43.0 18.1 138 760 1.6

A Y.S.

ksi

55.4 37.5 12.8 138 760 1 .8

A HRB 77.4 34.4 23.8 138 760 3.1

A %IACS 29.2 40.5 11.4 138 760 1 .1

T851 B U.T.S.

ksi

68.5 53.6 12.2 156 800 1.6

B Y.S.

ksi

55.4 43.0 21.0 156 800 3.2

B HRB 77.4 55.5 13.9 156 800 2.2

B %IACS 29.2 38.7 5.2 156 800 1 .0

125
TABLE VI

Additional Specimens Not Shown in Figure 66 Studied

(Specimen Numbers Refer to Table IV)

in the TEM

ipecimen "Pre-aging" Final

Number Treatment Temper Treatment

1001 None (direct quench) T4

1008 None (direct quench) T351

1009 None (direct quench) T851

1040 Sequence A 350°C for 20s T851

1042 Sequence A 350°C for 109s T851

126
TABLE VII

Average Hardness Values of 3 Percent Stretched

and Unstretched Specimens After Aging

Aging Temperature

(°C)

HARDNESS (HRB)

After 6 Hours,

Stretched

After 6 Hours

Unstretched

After 16 Hours

Unstretched

After 30 Hours

Unstretched

21 77 75 - -

35 78 76 - -

50 79 77 - -

150 80 74 76 79

170 - 77 79 80

180 - 80 82 77

190 84 79 - 71

127
CONDUCTIVITY,

%

I

ACS

ROCKWELL

B

WT

- %

C0PPER

Figure 1. Measured average composition (wt£Cu), hardness (HRB), and eddycurrent conductivity ( %I ACS ) distance across the thickness of a 128

15.24 cm thick plate of the 2024-T851 aluminum alloy.
CVJ

CM

O

o

ro

O

CM

o

O

6|AI % 1M

129

WT

%

Cu

Figure

2.

Calculated

solidification

"path"

(AB)

shown

on

the

liquidus

surface

of

the

aluminum

rich

corner

of

the

Al-Cu-Mg

system.

The

second

phase

to

form

is

CuMgAl^

by

way

of

the

monovariant

eutectic

L

-+

a-Al+CuMgAl^

along

BC

and

solidification

is

completed

at

the

ternary

eutectic

L

-*

a-Al+CuMgAl^+CuAl

at

point

C.
Figure

3.

Calculated

solidification

"path"

AB

shown

on

the

liquidus

surface

of

the

aluminum

rich

corner

of

the

Al-Cu-Fe

system.

The

second

phase

to

form

is

(Fe,Cu)(Al

,Cu)^

along

BC

followed

by

Cu^FeAly

along

CD.

Solidification

is

completed

at

point

D

by

the

ternary

eutectic

L

-*

a-Al+Cu

0

FeAl

7

+CuAl

0

.
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dJ

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to s_

oj •(- to -O d) T3

CL4-

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i- d»

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<U JC to P (-

CD

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C

O

O *r- r- -P

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to

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CD

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r

_ 00

O

O •(- J= CM

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S_ d> tO <8 -C -P C_3 d> CD O O -P C + -p E (0 •(“ •(- dl r- S- 3 4-> 4- 4- C 3 <C (0 C C 1— T- T— cr i 3

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oo —i •P

LD

d)S-3

CD

132
AUFe t

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CDC/1 "O ~0 > C <U = -r-

•i- o +-> -c: i—

+-> -rO P3 O

(O Ol

C ID

W CU Q.

<u cm a> =

S- lO« c r-- D1 *r o r- c ^ O.T- e£

•r- ^ +J <L> 3 «3 Ll. o -C U CM

SZ S- 4-> -r- 3

to CD 4-0

Q. O'r + E Q. +-> T3 r—

CU

O •<-

C +-> O C r— I to 0 0 8

Si

>)t-r tfl tO CU 4J f

-a *r- <u

•rC T3 -C

—1

(O-r -O -P 1 IO is CU o"* CU E r— S_ W Li- CU C e£ CU 1 s_ -r- 1 x: • 3 8 2 O CU 1

> i x: >=£ >> r"> r— +-> 1 J- r—

<C 8 <0 <

Q>

jc e

a> +j 4- t- lju > -£Z -I- o S- CM

+-> 2 CL 3 _ £= O O

4— •!—

O 4- ot/l T O S-

+-> o

a> to • »a c 4- tr >i to CJ o O > •«“ •<- c.O

r— CU *r— (0*1-

ns u_ -a

a CD

•r- •!- CU x: 4- — 4- £. a O O *r- •I- to “O <u

S_ CU *r— -C

CD 4-> r- 4-> E ro c o

3-P CU

W

C c c to *r- •r-CU<US_ E o S- cu to 3 i- Q-+-> CU r- CU <U 4- TLO

CU

S-3

CD

133

continues

from B

toward C

along

the

ternary

eutectic

line L

->

a-Al

+Cu2

FeAl?

+CuAl2.

Solidification

may

reach

point

C

where L

->

a-Al

+Cu9

FeAl7

+CuA 1?

+Si.
Figure 7. S' precipitates within borders of a-Al dendrites in the laboratory

cast ingot.

134
Figure 8. Interdendri tic structure in laboratory cast ingot.

0 -CuA1 2 and S-CuMgAl^ are identified.

The phases a-Al

,

135
Figure 9. Cu^FeAly i nterdendri tic phase identified by electron diffraction

and EDS in the laboratory cast ingot.

136
Figure 10. S' precipitates within the a-Al dendrites in the DC cast ingot.

137
Figure 11. Typical views of the interdendritic region of the DC cast ingot.

(a) Optical view. The dark gray blade is Cu^FeAl^; the large

rounded light gray phase is a-Al(Fe,M)Si and the fine eutectic

like structure consists of a-Al -CuMgAl

2 +CuAl

2

. (b) SEM view

shows the Cu^FeAly phase.

138
Figure 12. Interdendri tic eutectic structure in the DC cast ingot, are identified; a-Al , 0-CuA1 2

, S-CuMgAl

2

and Mg

2 Si.

Four phases

139
B

|0 4

Al

|0 8

Cu

X

Figure 13. (a) a-Al(Fe,M)Si phase in the indendritic region of the DC cast ingot

(b) X-ray spectrum of a-Al(Fe,M)Si polyhedrally shaped particle

in Figure 13(a)

.

140
89 cm

Figure 14. Location of various macrosegregation profiles taken from the

transverse section of the DC cast ingot.

141
11.75

C -r-

<D C I

y3ddoo %‘ s o

142
CHILL FACE 0.2 cm u 2.0 cm

5.2 cm 12.5 cm 17.0 cm 19.2 cm

Figure 16. Optical micrographs showing microstructure of the DC cast 2024

aluminum alloy ingot at various distances (in cm) from the chill

face. Keller's etch.

143
5-2 cm 12.5 cm 17,0 cm 19,2 cm

Figure 17. SEM micrographs corresponding to Figure 16.

144
11.75

11.16

10.59

lO

ro £o

•»» LUO<

OJ

OO

cr

Ll. LUO<

I—cno o3

•P

C

L. *r» O -P

.£=C

to oo

OJ I SZ T-

•P E c to

•r- <U

0) J= O -P

«3 4-

C

•r- JZ r— O

OJ

O) s_

-SC 3

-P OS O

SC

CU i—

<J CQ c

fO TS

-PC

to nS CM<c

to3

to

s-aj>pc

ai -pcoo

too

D.

t. -P

CUC

CL CU CL S- • o a> o 4- o

4- r—

<U *r- i— ora ns ns

U 3dd00 % < s

s- a> E

cu ai 3 > s- C

ns -C •T— -p E rs 3

a>

p r—

ns ns

•r— » 4-O C CXI S- o o

Q. •r~ CM

-P

C

a 4- o CU o

S- -p •i— -p

nS TS o

CJ) OS

O) CU c

s_ to *1

cj> L.

CU CU -p

to > to o to ns

s_ c O o ns

ns S- o

51 -P o

CO'

CU

s_3

as

145
DISTANCE

FROM

CHILL

FACE,

cm

(LONG

TRANSVERSE

DIRECTION)

Figure

19.

Macrosegregation

profile,

average

copper

content

versus

distance

from

the

chill

face

In

the

long

transverse

direction

(position

LT

in

Figure

14)

in

the

semi

-continuous

DC

cast

ingot

of

2024

aluminum

alloy.
2024

-AS

CAST o

cvi

(—®—) l^niS3N9VIAI % <s0

in o

CO3

CO

i-<u> T3

a> -p •r— • c 4- >>

0J •r- o

+->

a c •i— r— o r— 03 a o

CO E E 3 3 >i c

•i— •rTM CO r— E

<u 03 3 c C r—

cn o <0

fC *r— E +-> a CM

"O ai

O c s- CM ic

•I— *a 4-

s_ •i—O

a>C

Q- 3 •P Q. O O 03 CO o c C r—

<u •r* cn >>

to Ss~ r*“ o

aj r— -p > •r- 03

s:. S- a o e* JD

a) E 03 (— O r—

•r— +-> 4- +->

C

O o o

S- .a •p* a. •p

a>

a c .c CL) o -p co

•r—

+-> E co

05O co cn s- O

<u 4- s- s- o

cna

0) o -a

co c aj O 03 CJ

s- +-> 3 a CO -a

to •p“ CL) s: •a S- o

CVI

<u

5-3

cn

147
10

-I

REDUCED

CROSS

SECTION

LABORATORY

INGOT

OF

2024

AL

ALLOY o

o

in

0 3dfUVd3dlN31

148

Figure

21.

Measured

temperature-time

curves

during

solidification

of

10

to

reduced

cross-section

ingot

of

2024

aluminum

alloy

obtained

from

thermocouples

located

at

different

distances

from

the

bottom

water-cooled

chill.
0001

:

i'

III

I

ro “

suojoiiu ‘SONIOVdS 31iaaN3Q AHV0N003S

149

100

300

1000

3000

10000

MEASURED

LOCAL

SOLIDIFICATION

TIME,

seconds

Figure

22.

Measured

secondary

dendrite

arm

spacings

versus

local

solidification

time

in

unidirectionally

solidified

ingot

of

2024

aluminum

alloy.
o

CO "r- 3 +J

CO

U

S- CU

O) s- > 1-o

>)

+-> c

*f- 3 >1- +->

4-> CO

CJ CO •ocoo

O) r— JC

#0 +J o

•r- c

i- •«“ -po «

CU r— 'Z -rCVJ OCVJ 4-o

+-> o

CJ> c

>>

s-o

+-> co

S-o

-Q

tO

T3C«

CU

+-> c

CU 4-> cou .c

+-> 5-

CU E

Q.

O

CL 5- O H- 0

CU

CU o

01 c

to to

i. +J

CU 10 < -5 c0

+Ju

cuCO1 CO

ooo

S-oT3

cuu3 • -o >>

cu o

5- r—

•a to

cu

•4= i

•r- c T3 •!—

•r- E

r- 3 O i—

CO (O

OO

CVJ

cu

s-cn

•p—

Ll_

150
EoXoo

hh*O

GO

LJOz2cno C Q) O -C

•r- +-> • P +J

•rC

O

O 'r D) C C O r— *r- i—

>> i— J_ lo a o

00 4-> t— E <o O i- <V 4-> o

x: +J jo P O (0 _Q i— O

+-> ai c

-C o 4-> +-> *i- c +-> a) e o E o ai p s. w

«3 *4- I CL) CO

5- Cl) CO

+->

O

O

C S- 1— ffl u

«3 +-> o to -a

•i- T- cu c -a o

<e 3

-c u) -a 0 3 0)

a> co s- E s_ O d)P E > «/> S- rtJ

CL) -M

O

.c c

+-> CL) +-> i—

SC r—

a) o «3

+-> o c 4- o

•O S- *r- cu +-> CO Q- O

CO Q. OJ

a> o s- C O •!- •a T3 s- -a *r- (O c c

=n <o 3

«=!- CNJ

a;

s-3

CO

•r— Lu

151
TEMPERATURE

TEMPERATURE

a

SOLUTION HEAT TREATMENT

495°C (923°F ) 75 MIN

TIME

TIME

Figure 25. Schematic representation of the thermomechanical treatment sequences

given the 2024-T851 and 2024-T351 aluminum alloys.

152
|

5.08

r

-

Figure 26. Schematic of tensile test specimen. Dimensions are in mm.

153
Si

ACS

AFTER

US

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BEFORE

MACHINING

2024

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on

cu *o

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•P cum •r— TP

X o o l— C C o 3 CO i •r— CO S- TP 4- «o- JC a. C S- CM a -P CL o 3 o CO r— CO o to CM E 4- —- ©m

ID

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154
HARDNESS

AFTER

US

HARDNESS

BEFORE

MACHINING

£024

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CDo

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Q

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T851 0

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a. jo 5- cn 4-> Q.

c >>-o CO CL

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co E CD CO CO E «C E CD -Q s

z

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co E -a CO

s- s- co O

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CD SZ

S- +> SZ

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CO CD SZ •r— +-> CO sz

CD -M

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CO S- E 1—

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155
ULTIMATE

TENSILE

STRENGTH

C-CURUES

2024-T851

SEQ

A

LD

T—I—I—I——I—I—I—I

1—I—

I

-1—I— "T

" T I—I——I—I—I—

T

<S> 
 
 


ll

. 1 LD <S> LD LD 1

j -<r n m cu cu 11

QUO o

1

J

156

Figure

29.

(a)

Ultimate

tensile

strength

C-curves

for

2024-T851

sequence

A

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

tensile

strength

for

an

isothermal

sequence

A

type

"pre-aging"

heat

treatment

and

a

subsequent

aging

to

the

T851

temper.
jo o o o o

|o o o o o : io o o o o

110 o id o in

cm co co «o- .•ill

;o o o o o O lil O U1 o CM CM co co <o- j

II II II II II

I* + X

O 11

;!i j

<ro

UJ

</>

LD

00I OJ

<S> OJ

<S) Q

<E3O

<ro

LD3o

LU

CL3

cjd <rLU

cu 01 P +-> $- CO 1 (/) (— QJ a C +J T- ^ 1 •i- +-> T3 CO

cO • CO

C T- <D cd<C a r- J_ r—

<a oo CU

a

cu IO 4-> s- -a a cu t— co •r* • aj cr O CO O o

+j qj +J ro 0 -P 3

§> ••

O

O

C CO

O ID r— CU »r co -P O •O' Q. CO c o CU cu cu o o

-C r—

C

r pm -p

+-> ID •!- 1— +-> CD CO i— CO o o c I— cu +-> CU 1 T3 • ss- CU — -p o •P CM £ r— ID 1 CO O CO cu = CM CM cO > CD cu <u -a CU S= 1 tr— r- J. i— •(— CD • 1- o cu CD X C

CO 4- -C CU CO rO C -P O 1 A •r— cu cu C 01 o i-

+-> > -a 01 s- o

s- c XJ 0-0

CU 3 *0 •r—

— to r»

H +-> o 4- CM o

CD

CO 1 cu c cu o EU a o -O O O •I- o CJ +-> -P o +-> CU r— >o- r— -C 10 •P CD O 3 -P C C o o UH

-P CU *r- CM +-> CU CD-r- O S_ x: c c S- 3 1 o

<E -P -r- 3 cu -a ID 3 CO CL c CO 4- 3 4- cu CD aJ O O ID i. °r- 1 O3

-a CD 3 CO to cu cu -p CU OP c >J CO CO i- <r 3 ro T- r— S- 3 cO o r— r— t— cu cu p— 3

tO 3 4-> Q- CL cr > O CO co E to r— E ® * T3 cO CD •r- +J J= •% O Or- X o o

S- o cu c 0 3 CO CU S- CD cu O CO CU C Q- CO 3 ID

CO 3 i- CL S- O" CO

<U r— r- CO cu s: <o —- > P O > "O CO -P

•r- a CU '—* 01 i— c cu s_ o pro CO -O •r- O —' +-> co -Q -P -a co

CD

CM

cus-3

CD

•r* U_

El LU <E CD 3 CLUJ

Q ^CDkh

157
YIELD

STRENGTH

C-CURVES

a034-T8Sl

SEQ

A

CM <U J= ! to o 4-> I s_ a> c o e a; o i 4- •!- 3 +-> 1 +J O' i to 0) cn

CD r— to c > to •r- i- o r— CD 3 *r- lO to O +J E

1 -1- l- -t-> o s_ CD C

o sz a)

'

-t-> 3

4-> CD O CT

SEC

cu .c to cu 00 +-> -r- to 4- -Q

4- CU C 3

O > to to 1; •r~ +> CD S- 10 Lti c o

a) to 4- T3 ( e: o cu e

Ht

S- > -£= «0

cu i. +-> 1— CL 3 CD4->

CJ C C

_1

<r

cm a> cu

• 0) 5. E or+)+i!:

o —1— to to

cu »-4 e TJ s. 1- +-> • t— 4-» H CD 00 CU C >>•!- -P 01 CU O >> to o £- CU

+-> 1— T3 -C

c/> *3 •O <

CU

a) +J =

<oa

cnc

<o

O XJ c c

a) -r- 3

cr a» a>

to 4-> =

CD

*0

I a)

fc. j CL

o

co

cu

s-3 —r r

1

i i ri II III i i—

r

—

i—i——i—

r

1 1 ^ •rTM


 
 
 
 «s> U) 
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 , CD n 03 03 1

j

i-uje:q_lljq:<e*-:dq!Llj QlilO O S

158
li

c_> <_> o <_> o ,00000

ID00000 O ID o If)

Ocl OO CO ^ «d" OOOOO

IO IDO ID

O OsJ cm no ro «d11 11 11 11 11

j* + X o I

<roU

CD

U)

00

I—I 03

«s> 03

I-o

ui

o:h

CDQ

•J

Ui

»—

>o

UJ

I-

<EO

<EO

CDQ

UI

ck

CD

<rUi

CD < a)

<u jt o -M

to c

a> ai a> 3 3 1- r— CT «o

•d <u > to CO

<u to I— c • 3 ID *r— '—

IO 00 1— 1—

S_ I— <U

flJ I 3 > > <u a)

CsJ SZ r—

"OO to

a) cm *8 a>

4-> -T3

O

+->

u c o o o> a> r~ 4- JC "'O Cl 4-> *f- ai <+- SZ > -O e

+-> $_ c o

DI 3 10U c 0 1) I (If) i~OQ.CZ

+J o <u

to sz t— o

+-> to i_

d 01 ai r- C+l a d) Q)-r

>>

i_

C ID

M 3 CD

to

<u 4- >> sz -a o 1—

+-> f—

V

a> a> +-> l-'i- c in O >>T- E

r— *r* to <u x

<U JC IO o 3 +) S- 1— to CL

Id Ol'r Q. > C id

•r- <D > owe CU 3 •«“ T3

i- rC 3 T3 id to <u *a jQ

•O +» •!“ <U Id r— J. s:.- o ai 3 to 4-> O +> ^-'1— <U id —

JQ id -C o o I— to oCO

a>

iCT>

00


 CD


 ID CD o

03

EIUJ<CCD=3 Ck:UJO ^CDi-t

159
HARDNESS

C-CURVES

2024-T851

SEQ A ! C O CQ

<U r—

-C r— 4-> I— CD <0 s a> ^ .c • a

1/10 = >>cc ai o e • i— "O •!— s_ r- (U Oia)

io+> « a

•O I E C a a> ai

•«- s- +->

n i «n o C *r- in

a) cd co 3 a) a. 1—

CT.C >>

ai +->-*-> a> V) cnC +J f— c

in •!— CD o

00 c a +J l— -r- c i <o CD ai ^ -M 3 c

CM jQ cr r— O O CD CO CM lO CO Ss- o r— +J O <4- CO c

H- E CD 10 S- 3

lO O) CD cr

CD E sz CD > -r- 4-> co

S- +->

O -Q 3 CO 3 U «— •r— 10 1 ca o o c CO

•r- CO 10 +-> -a

10 T- s- c

<U S- o CO e o *4—

-o 4-> s- a; 10 C

»o -£= 10 CD nz CDc 5

<u -a CO > s_ <U

(O *r~ CO s_ 4->

ro

CD

S-3

CO

f—LLJEZQ_lAJQ£^t— 3 Q£UJ OUJO

O

160
MEASURED

US

CALCULATED

HARDNESS

2024-T8S1

SEQ

A

O O O O O

o o o o o D CD CD CD CD

in o in o in

ojfornrj-'t

i i i i i O O CD CD O

o m o m o

C\1 C\J CO

ll II II II II

* 4- X O

E:ui<EO)=3o:ujo ia«

161

Figure

31.

(b)

Measured

values

of

the

Rockwell

B

hardness

plotted

vs.

values

calculated

using

the

hardness

C-curve

for

2024-T851

sequence

A.

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
CONDUCTIVITY

C-CURUES

2024-T851

SEQ

A

Lf)

<S>

1 ill II T—1—1—1—T~r i'"T—i—i nr“I—1—


 <S> O

\s> 
 LO 
 LO 
 ro m CU OJ

S— UJHZQ_LijQ^<rt— ZDQiljJ QUO O

0

S

r

162

Figure

32.

(a)

Conductivity

C-curves

for

2024-T851

sequence

A

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

conductivity

for

an

isothermal

sequence

A

type

"pre-aging"

heat

treatment

and

a

subsequent

aging

to

the

T851

temper.
MEASURED

US

CALCULATED

CONDUCTIVITY

2024-T8S1

SEO A

| o o o <_> o 10 0 0 0 0

j o o o o o in o lo o m i\i ro ro ^

o o o o o

I o m o in o [ !\l !\J fl m

|

j ii n n ii ii

* + X O Q i

s:ui<ru)=3tt;ijjQ

163

CALCULATED

*IACS

Figure

32.

(b)

Measured

values

of

the

conductivity

plotted

vs.

values

calculated

using

the

2024-T851

sequence

A

conductivity

C-curve.

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
ULTIMATE

TENSILE

STRENGTH

C-CURUES

2024-T851

SEQ B

Figure

33.

(a)

Ultimate

tensile

strength

C-curves

for

2024-T851

sequence b

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

tensile

strength

for

an

isothermal

sequence B

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T851

temper.
MEASURED

US

CALCULATED

U.T.S.

2024-T851

SEQ

B

<_)<_> O O O o o o o o O O O O O in LO O ift

i i i i i o o o o o O ir> O IT)

c\j cm n n<3-

ii ii ii ii ii

* + x o 0

dUJ <£</)=) O' UJQ

165

Figure

33.

(b)

Measured

values

of

the

ultimate

tensile

strength

plotted

against

the

values

calculated

using

the

C-curve

for

2024-T851

sequence

B.

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

line

is

the

scatter

band

(approximately

95

percent

confidence

level).
YIELD

STRENGTH

C-CURVES

S024-T8S1

SEQ

B

LO

<S>


 <S> 
 
 
 <S> LO <S> LO 
 LO

''T 'T n n CU (U

{

f-UJEQ.UJQ'^t-DQ'UJ QUO O i

166

Figure

34.

(a)

Yield

strength

(0.2

percent

offset)

C-curves

for

2024-T851

sequence

B

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

yield

strength

for

an

isothermal

sequence

B

type

"pre-aging"

heat

treatment

and

a

subsequent

aging

to

the

T851

temper.
*

=

200

-

250°C

+

=

250

-

300

°C

x

=

300

-

350°C

0

=

350

-

400°C

MEASURED

US

CALCULATED

YIELD

STRENGTH

a034-T851

SEQ

B

°

=

400

-

450

°c

e:uj<lu)3q:ujo

167

Figure

34.

(b)

Measured

values

of

the

yield

strength

plotted

vs.

values

calculated

using

the

yield

strength

C-curve

for

2024-T851

sequence

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
HARDNESS

C-CURUES

2024-T851

SEQ B

168

Figure

35.

(a)

Hardness

C-curves

for

2024-T851

sequence B

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

Rockwell B

hardness

for

an

isothermal

sequence B

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T851

temper.
ooooo ooooo

o o o o o in Ln o m

t i > i i ooooo O Ifl O irt o c\j c\j ro ro <3-

It II II II II

CL)

CUIOKftSQ'lAJO Ztttt

169

Figure

35.

(b)

Measured

values

of

the

Rockwell

B

hardness

plotted

vs.

value

calculated

using

the

hardness

C-curve

for

2024-T851

sequence

B.

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
CONDUCTIVITY

C-CURVES

E024-T851

SEQ B

Figure

36.

(a)

Conductivity

C-curves

for

2024-T851

sequence B

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

conductivity

for

an

isothermal

sequence B

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T851

temper.
MEASURED

US

CALCULATED

CONDUCTIVITY

2024-T851

SEQ

B

O O O O <_> o o o o o O O O O O m o ia o m

.1111

o o o o o

o m o in

<xi <nj ro ro ^

it it n it ii

* + X O

CUi€U)DKUfi J?h<EOO)

171

Figure

36.

(b)

Measured

values

of

the

conductivity

plotted

vs.

values

calculated

using

the

2024-T851

sequence

B

conductivity

C-curve.

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
ULTIMATE

TENSILE

STRENGTH

C-CURUES

2024-T3S1

SEG A

172

i;

Figure

37.

(a)

Ultimate

tensile

strength

C-curves

for

2024-T351

sequence A

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

tensile

strength

for

an

isothermal

sequence A

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T351

temper.
MEASURED

US

CALCULATED

U.T.S.

2024-T3S1

SEQ

A

ooooo OOOOO

O O O O O in if) O to

i i i i i ooooo

o m o in o

it ii n ii ii

* + X O

173

Figure

37.

(b)

Measured

values

of

the

ultimate

tensile

strength

plotted

against

the

values

calculated

using

the

C-curve

for

2024-T351

sequence

A.

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

line

is

the

scatter

band

(approximately

95

percent

confidence

level )

.
VIELD

STRENGTH

C-CURVES

2024-T3S1

SEQ A

rI

174

Figure

38.

(a)

Yield

strength

(0.2

percent

offset)

C-curves

for

2024-T351

sequence A

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

yield

strength

for

an

isothermal

sequence A

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T351

temper.
*

=

200

-

250°C

+

=

250

-

300

°C

x

=

300

-

350°C

o

=

350

-

400°C

MEASURED

US

CALCULATED

YIELD

STRENGTH

2024-T3S1

SEO A

D

=

400

*

45

o°c

CLU<CU)=)Q^UiO

175

Figure

38.

(b)

Measured

values

of

the

yield

strength

plotted

vs.

values

calculated

using

the

yield

strength

C-curve

for

2024-T351

sequence

A.

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
HARDNESS

C-CURVES

2024-T351

SEQ A

!'1II

176

Figure

39.

(a)

Hardness

C-curves

for

2024-T351

sequence A

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

Rockwell B

hardness

for

an

isothermal

sequence A

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T351

temper.
MEASURED

US

CALCULATED

HARDNESS

2084-1351

SEQ

A

O <_)<_> o o o o o o o O O O O O in o tr> o m

cm oo co «ai i i i i o o o o o

o i.f> o in

CM CM CO 00 «3

-

II II II II II

* + X O

a*

CUICCWSQ'IaJO XCKCQ

177

Figure

39.

(b)

Measured

values

of

the

Rockwell

B

hardness

plotted

vs.

values

calculated

using

the

hardness

C-curve

for

2024-T351

sequence

A.

T

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
CONDUCTIVITY

C-CURVES

2024-T351

SEQ A

178

Figure

40.

(a)

Conductivity

C-curves

for

2024-T351

sequence A

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

conductivity

for

an

isothermal

sequence A

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T351

temper.
*=

200-

250°C +=

250-

300

°C

x=

300-

350°C o=

350-

400°C

MEASURED

US

CALCULATED

CONDUCTIUITY

2024-T351

SEQ A=

400-

450

°c

179

Figure

40.

(b)

Measured

values

of

the

conductivity

plotted

versus

values

calculated

using

the

2024-T351

sequence A

conductivity

C-curve.

The

solid

line

isa

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
ULTIMATE

TENSILE

STRENGTH

C-CURUES

2024-T3S1

SEQ B

180

Figure

41.

(a)

Ultimate

tensile

strength

C-curves

for

2024-T351

sequence B

alloys.

The

curves

give

the

critical

times

for

obtai

ning

the

indicated

tensile

strength

for

an

isothermal

sequence B

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T351

temper.
o o o o o o o o o o o o o o o lo o lo 10 i i i i i o o o o o O 1-0 o LO o c\i cm n n n*

ii ii ii X o

CCk O

LlI co

u>n

t-i rr03

03

COHo

LxJ 5~

<r3o

<co

co3Q

LJ

Q&3

CO

<E

LU

COQ

UJ

I—

<£

0-3 3O•-J <co

E:LU<CC0 3Q:LiiO ^co

—«

181

Figure

41.

(b)

Measured

values

of

the

ultimate

tensile

strength

plotted

against

the

values

calculated

using

the

C-curve

for

2024-T351

sequence

B.

The

solid

line

isa

line

of

unit

slop

and

the

dashed

line

is

the

scatter

band

(approximately

95

percent

confidence

level).
YIELD

STRENGTH

C-CURVES

2024-T351

SEQ B

182

Figure

42.

(a)

Yield

strength

(0.2

percent

offset)

C-curves

for

2024-T351

sequence B

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

yield

strength

for

an

isothermal

sequence B

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T351

temper.
*

=

200

-

250°C

+

=

250

-

300

°C

x

=

300

-

350°C

o

=

350

-

400°C

MEASURED

US

CALCULATED

YIELD

STRENGTH

2024-T351

SEQ

B

=

400

-

450

°c

SEl&J<E<A30£UIO 1

183

Figure

42.

(b)

Measured

values

of

the

yield

strength

plotted

versus

values

calculated

using

the

yield

strength

C-curve

for

2024-1351

sequence

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
HARDNESS

C-CURUES

20S4-T3S1

SEQ B

184

Figure

43.

(a)

Hardness

C-curves

for

2024-T351

sequence B

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

Rockwell B

hardness

for

an

isothermal

sequence B

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T351

temper.
*

=

200

-

250°C

+

=

250

-

300

°C

x

=

300

-

350°C

o

=

350

-

400°C

MEASURED

US

CALCULATED

HARDNESS

2024-T351

SEQ

B

°

=

400

-

450

°c

IS) r10

LD

LO

O

U£

t-

<E

O

-3

<C

o

:uj<E<n=)C£UJO xocca

185

Figure

43.

(b)

Measured

values

of

the

Rockwell

B

hardness

plotted

vs.

values

calculated

using

the

hardness

C-curve

for

2024-T351

sequence

B.

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
CONDUCTIVITY

C-CURVES

2024-T351

SEQ B

186

Figure

44.

(a)

Conductivity

C-curves

for

2024-T351

sequence B

alloys.

The

curves

give

the

critical

times

for

obtaining

the

indicated

conductivity

for

an

isothermal

sequence B

type

"pre-aging"

heat

treatment

anda

subsequent

aging

to

the

T351

temper.
*

=

200

-

250°C

+

=

250

-

300°C

x

=

300

-

350°C

o

=

350

-

400°C

MEASURED

US

CALCULATED

CONDUCTIUITY

2024-T351

SEQ

B

=

400

-

450°c

187

Figure

44.

(b)

Measured

values

of

the

conductivity

plotted

vs.

values

calculated

using

the

2024-T351

sequence

B

conductivity

C-curve.

The

solid

line

is

a

line

of

unit

slope

and

the

dashed

lines

are

the

scatter

band

(approximately

95

percent

confidence

level).
C-CURUE

COMPARISON

2024-T851

2024-T351

2219-T87*

Li)

<S> ss_ •> o O r— 4- 4-in cn oo oo c •«— *i— oo 1— t— sz oo

cu i -o +J ^ (a > >=3- c cn 4-> S-(M O C O «3 3 0 a. a> on -o U N W $- i a> o 4- s- O 1- LO 00 o -a co i—

-t-> 00 o 1— l— 1 0) 4-> a) i cn

oo c ai •r— 1—

4- cu -c >>CM CM 4- e 1— O CM O 4-> O) CM

<o jr cu

+-> a> • +-> i- -c C S. Ifl o 1—

CD +j >>4— 4- a o o

S- = r— •i— • o

O) CTii— +J CO -—

ac io

C *

Lxl cn

•ia) r-. NOTE o o co • fC 3 s_ to h- O 1 C ai -— i CL) i— CL oo

i- E SZ r— LU -C Cl 3 O O CM +->=1—

CT> K) cn c cm

cu i—

C CO (!) 3 l. t- a> * 5. CTO

s_ a) *3 4-

+-> o 00 +-> kJ (/) CH CO (J *r- <r CJ 1 <U CL) CO s V o 30) 3 S- ^ • O i— CTf— i— *r“ '—

i—

ai <u cm <0-0 0 r—

H •r— c/> CM

>> > to

(O ' H -a -a jc -a aj a:o

4- c c 4-> E c o o <a « CTO (O e c s- cu c <C « O) 4- •> s- O r— S- i— cu co a> un +-> -a un 4- 1-000 CO <1) CO <u i. ch c 1— oc

fo a> i "O •«— 1 Q. 3 «=d- 1— (O *d- e E ctcm d) +-> CM o o a) o •1- -Q O S- O 00 CVJ >> O CM 4-

Lf)

'd-CU

s_3

o>

O

O


 o

LOn ai

i- uj ea. ui oc <e i- => at uj QUO

O

188
ULTIMATE

TENSILE

STRENGTH

US

YIELD

STRENGTH

2024-T8S1

J—

I

% XhO

US

O'

E'- enq

yjH

H(n cn »—

•

189

Figure

46.

Comparison

of

ultimate

tensile

strength

vs.

yield

strength

for

2024-T851

with

the

correlations

predicted

by

the

C-curves.

The

dashed

lines

are

the

scatter

band

(^95%

confidence

level)

obtained

from

a

least

squares

quadratic

fit

to

the

data.

The

solid

and

dotted

lines

were

calculated

from

the

C-curves.
UTIIMATE

TENSILE

STRENGTH

US

HARDNESS

a0a4-T851


 O)


 00

G>r-


 COo

in


 ^r


 ro


 03 o<uc

a> •>— -e <a

S- I— +-> O -Q

+- o

to '—

<u r— > <u

i- > 3 cu 0 —1

to -o

a:x

«» cn

co

LxJ Za

01

<r

s-

+-> to

to

toO) o QJ C U a cu c

s- x: cu

*a +-> "O j= *r- >><+- • X) C

to

O > -o O

<U £P O to

CO'i— 01 CO c1 cu a) —S-0.-0c

to (O

<U

C -Q r- O

•I— •(— S- tO +J CL) C K)P

CU r- 4-> +> cu «o

s» o

S- to oO CL) E -c

i- CL) 4-> +J -C

r— 4-> CL) 3 S-

<4- 4-> O *r- tO 3 CU C C O r— •(— lO LO r—

f- CO

S_ t— -O

«3 I CU CL«d" .C E M to 0 0(0 O C\J -O

i->* «ocu

4-> «

-o

c:

(Oo

to

a>

*o p> • (O to -o cu>

CU J- J= 3

+j

ai o o

+j

<u

+-> j=

i- +j

<4- E a o

•r- S- +J <4- (O

s_ -a

-a a>

(O 4-> 3 #o cr i—3

to

a

cu i—

i- (O (O o 3

CT QJ to J_CU

4-» 3

lO

(O to

CU cu I— c

•r— CU i—

-C

4-> TJ

CU E -)-> O 4-) i- O

h- -a

cu

s-3

CO o

00


 Is-
 CO


Lf>

«S> <r
 n

=)f-cn i^cnH-i

190
ULTIMATE

TENSILE

STRENGTH

US

CONDUCTIVITY

2024-T851

inm

03

n*©

o>n

oon

ino

<rH

«* >

6—OQ

z:©o

T3

CU

+-> TJ +-> e <U O o e -a 4- <u •1— jr "3 -O flh-P c

+J JD «3 to O

T3 • •o

CO '—' »r- >> cu r— i—

-M > CU o

f— E > to > 3 at

•i“

U r— CU

4-> 1 XI

<_>

O aj t— 3 u

T3 CL> E E -E CU • O •*-> •O (O o •1- +-> >>4- «3 • -Q e -a

to o > T3 o cu

CU JZ • -E +-> +3 to -M o to cu o>t- cr> o > E T3 <? +-> E

cu ai

V 3 e e 4-> O

•*-> Q-TD •!“ 1 to E 4- c_> to fO

<U E -O

u cu r- O •1— x:

•f— •»— E +-> +-> 10 4-> cu to E «3 +-> E E

CU i— 4-> -a

O +J QJ •O IO E E O 3 4-

CU E

+i

O

to CT T3

rO cu to CU E J= cu +-> •r- a) +-> E

-M -E <0 r—

4-> CU 3 3 3 ECU x: <o to t—

4- +> *o O *r- CO 4-> O 3 cu to E E tO CU O r— •I- CU E

«/) LO 1—

J <u 1- 00 E 1

10 10 3 « 1 cu to CL«^- -E E cu E CM to o E o o (O E •I— O CVJ -O 4- r—

00

*=d"

CUe3

o>

•r—

Ll.

©©

oo r-

©

©

©

©

in in ^ n

3e-<n

191
YIELD

STRENGTH

US

HARDNESS

2024-T851


 C7>

<S> 00

r-


 CO


 Lf>

<S> ^r


 m


 nj

co

O'X

cn

co

LlI q

<L

a>

s- -C CD CD

4-> S_ 3

•r- lO 5 +-> to to tO <D r— O) tO C

LT)

C CD •»— oo •«- r— i

* h- r- 1 *o -O ^ -a CD

cvj a> E 4-> o -c= O 4-> C\J to J- O

to 4- -a

i- -a o -o -o

4- a> CD

C

.c

C lO

to 1

•r— -M tO T3

1C 4-> 1-

-a • -Q r—

to

O

O

to CD to

to > -—

a> s- i— CD e 3 CD SZ

-o CJ > t—

S- 1 CD to o r—

.c • <u ai to • -CI a +-> tO 4-> c to > cd -a >>T3

-C S3 •i- CD

+-> 4- -C CD*a C +J c ai o

a> +-> o o

s- o +-> +-> -r- tO "O LO +-> QJ CD i-

"O S- <? 4- i— CL

O) o

f- tO T3 i- >> c c +-> o to to

<4- -r- -Q S- O 4-> o ITS S- to C i— CD 3 o a) +-> c

to s- +-> r- S_ to to

s- o O CD tO o to S- Q. (O e <u CD 3 O JC sz <zr O 4-> +-> to

tT>

a>

s-3

CD

‘r—


 00


 r-

«s> co
 LD<S> <S> n


 m i

>HljJJQ COI— QiLiJZOI—b^COH-t

192

calculated

from

the

C-curves.
YIELD

STRENGTH

US

CONDUCTIUITY

E0S4-T8S1

o ® O O

00 <0 03

Lf>

’T

n

<<r

(U

cs> "T

)n

oon

O)o

<r

t-t a

3

o

Q

O

>HUJJQ 0)HKUZOHZ

193

Figure

50.

Comparison

of

yield

strength

vs.

conductivity

data

for

2024-T851

with

correlations

predicted

by

the

C-curves.

The

dashed

lines

are

the

scatter

band

(^95%

confidence

level)

obtained

from

a

least

squares

quadratic

fit

to

the

data.

The

solid

and

dotted

lines

were

calculated

from

the

C-curves.
HARDNESS

US

CONDUCTIUITY

2024-T851

Lf> n

OJ

^ro

^r

a»n

oon o

<L* >

f- ^ =>

Ij »-i o3QOO

CL) sai5

<u

S-

<o

+->

</» to a) cj id c C O) *-

co

in

oo

l— -a

<d T3

cd

CD E +-> -C O +->

NW !.

O o «o m- -a

CVJ T3

-o T3

S_ CD CD

C o .r: c *a

<+- I— f—

IC T3

(O +> *f—

+-> • -Q i

fl 1/1 o o

T3 CD od > ^ >, S- CD D 3 (Dr

*r~

O > l— > I CD

•rO n-

+-> O CD CD #0 3 r u -P op c io C CD TO O >>"0 o -Q •!- CD • -a

10 CD > 4-> a

lO *r- CO T3

CD CD C S- o Q.

s-

<a to rz co

if-O •*-> «3 C i— O CD 10 S-

•r* S-

$-

O

•o o

CLE <u o -c O +j

in c +-> oo o

+-> in 4-> CTl -r- ? «+- a

T3 *r- C +-> •a -o

i- *3

CD 3

-p cr

+j

«o CO o

CO

<D

S-«

co- co

CDS_3

cn

O

<S>


 00


 CD OJ

<EQ£ UJCDCD IKtt

194

calculated

from

the

C-curves.
ULTIMATE

TENSILE

STRENGTH

US

YIELD

STRENGTH

2034-T3S1

ID

CO

CO

ID

LD

<S> LD

LD

<S> ^r

LD

CD

<3

UJ

OC

E

CD

Q

-J

Ul

i—

•

>

CD 00 CO LD m

31- CD i^CD*-»

195

Figure

52.

Comparison

of

ultimate

tensile

strength

vs.

yield

strength

for

2024-T351

with

the

correlations

predicted

by

the

C-curves.

The

dashed

lines

are

the

scatter

band

(^95%

confidence

level)

obtained

from

a

least

squares

quadratic

fit

to

the

data.

The

solid

and

dotted

lines

were

calculated

from

the

C-curves.
UTLIHATE

TENSILE

STRENGTH

US

HARDNESS

2024-T351

00 CD LD T

<s» oi


 oo


 r-


 CD

<S>

LT>

T

<S> n

ca

o'

CD

CD

LU

Q

OL<r

3E-CD ^CD«-i

|1i1 9 196

I

Figure

53.

Comparison

of

ultimate

tensile

strength

vs.

hardness

data

for

2024-T351

with

the

correlations

predicted

by

the

C-curves.

The

dashed

lines

are

the

scatter

band

(^95%

confidence

level)

obtained

from

a

least

squares

quadratic

fit

to

the

data.

The

solid

and

dotted

lines

were

calculated

from

the

C-curves.
ULTIMATE

TENSILE

STRENGTH

US

CONDUCTIUITY

3024-T3S1

ru

^r

oon

CDnn

OJn


 n

oo

ru

u>o

<r

4k >oGZoo

00 or-


 LO

<S> ID

<S> »rQ

cn

ZDH- V)

197

Figure

54.

Comparison

of

ultimate

tensile

strength

vs.

conductivity

data

for

2024-T351

the

correlations

predicted

by

the

C-curves.

The

dashed

lines

are

the

scatter

band

(^95%

confidence

level)

obtained

from

a

least

squares

quadratic

fit

to

the

data.

The

solid

and

dotted

lines

were

calculated

from

the

C-curves.
YIELD

STRENGTH

US

HARDNESS

2024-T351


 (J)


 00

r-


 CO


 LD

«S> <r


 n

00

01

x

* CO

CO

Q

01

<r

co LO CO

I>- »—i LlI *—1 Q COJ— QILUZICSJ— X ^!CO>-»

198

Figure

55.

Comparison

of

yield

strength

vs.

hardness

data

for

2024-T351

with

the

correlations

predicted

by

the

C-curves.

The

dashed

lines

are

the

scatter

band

(%95%

confidence

level)

obtained

from

a

least

squares

quadratic

fit

to

the

data.

The

solid

and

dotted

lines

were

calculated

from

the

C-curves.
YIELD

STRENGTH

US

CONDUCTIUXTY

2024-T351

(U

00n

<U

re

re

cu

reo

re

oo

cu

if) o

<Ew*

«»

r>t-oQZoo 0) s_ i— «3 LO

ro

w to 1— CD L0 ai 1 E *0 e

*3" *r— CD ‘r— CM .— r— OCM -O «3 TD

CD ai S- -C E 4-> O CO O +-> <1-10 1. o *0 4-

*o -a P DTI a to XZ CD c UH C

•r* <o >1 #0 -a

+J . +J i—

•r- tO jQ r— > a» o o

•r- >

+-> s_ to a 3 i— CD 3 0 0) C

-o i > I— COO) O r- a a) • • -c a) *0 to • -t-> o 4-> 0) to c «3 > > >> a> -o s- -Q -a 3

XI -i- 0) o

+-> -o 4- jc i o) <d c c +-> o

+-> o

ai o o o ai

S- *r- 4-> xz

4-> -a >9 +j

tO 0) LO -t-> s- cn <- E

-o Q. <? 4-

O

i— *—S- CD to O 4— •r CT3 r—

>» o c -4-> -O

f- to 03 0) 4- +) J3 S- -4-> o *a a «3 p— 4- to r— C 0) CD 3 3 O S- +-> cr o

to S- +-> r—

•t- o «c to «o

s- o o CD

O

«3 to S- CL-C «3 a) E +-> aJ 3 S- O -r- cr cd O 2 4-> to 2

LO

LO

<u

s_3

cn

<S 
 
 <£ in

OO o ^ n cu

>MUJJQ </>t-Q£LUZ:OE-X ^V)HH

199
HARDNESS

US

CONDUCTIUITY

2024-T351

\SJ




00

n

CO

m

n

OJ

n


 n

oo

OJ

o

<E

«» >

O

Q

O

00 CO 03

x<xo'ozujcncn

200

Figure

57.

Comparison

of

hardness

vs.

conductivity

data

for

2024-T351

with

the

correlations

predicted

by

the

C-curves.

The

dashed

lines

are

the

scatter

band

(^95%

confidence

level)

obtained

from

a

least

squares

quadratic

fit

to

the

data.

The

solid

and

dotted

lines

were

calculated

from

the

C-curves.
UTLMATE

TENSILE

STRENGTH

US

HARDNESS

2024-T851

ik|\

\

V

N ^

N\

N V

\

\ \

i a

z

<cA

v T

>r

\

\

x «XN \

\ \\A

B v Yx\ \

Ul CMH

1“

oo«

Ul UJ o

M M H

\ \y

X \

\

« \

?

A

i'\

D

LINE

£0

LINE

ED

LINES

\

>

i%\s AV\ i\V\ \

*

\j\\ \

\ \V>

\ \\ V

•Ji-in oo« inoo

i i xo

—i—Wr'

\ \ \

\ \ \

\ \ y \ \

\ \

\ \

V

k \

\

\ 1

V

<S> O)

CO


 ID


 in

CO

OJ

ca

a:

«k in

in

UJ

z

Q

a:

<E

00 10


 in

®

CO

31- in b^in»-H

201

Figure

58.

Comparison

of

the

ultimate

tensile

strength

vs.

hardness

data

generated

for

2024-T851

in

this

report

with

the

data

from

soft

21

24-T851

reported

by

Petrak

and

Gunderson

(heavy

solid

and

dashed

lines,

representing

a

least

squares

mean

and

the

lower

90

percent

confidence

level,

respectively,

from

Ref.

12).
ULTIMATE

TENSILE

STRENGTH

US

CONDUCTIUITY

2024-T8S1

LD

'T

'T

-Tn

'T

CU


 T

O)

cn

oon

-ae E «3 O $- >1 S- T3 cu

4-> 4- •r* 2

•r— r— o > re o 1

• •i- +-> 00 *—

4-> «3 cu CM o -a >>-c i

3 > +J td <u *3 • c .£ CU -O 4- O 4-> sz £ O) o —-- CO cc

.£ • +-> c £ E

tO T- o CO o > 3 to CU 1-

J- E 4- .£ +-> cu H S.T3 to 0\ CD O e <u >> E CL 3 s_ cu a> CD CO %

s- s- 3 >

+J -a CT

to to £ to +J

(O o

a» .£ +-> cu r- +-> to CL

•r- CO re to

tO E s- cu CU C •!— +-> r— S- a cu

+-> r- Q_ re rt LD r—

(U CO >5 CD CU +J 1— -O £ >

cO 1 •r— cu e *0- -a +-> 1

•p- CM CU £

+-> o 4-> CU cu r- CM J- to o 3 O cu E

S- CL 1_ CU

ai o CU Cl 

a

-£ 4- s- CD T4-> 4-

-a r— £

4- <U cn

O

O 4-) 00 00a

fO 1a) c s- i c +-> o cu re- •r— £

to £ C\J r— CU

•T- CU r— o

S- CD CM T3 s- CO CU <u CL cO +J -£ Q. E -M 4- CO O CO O re o o -a •

to -a CD

CDID

CU

s-3

CD


 00
 r-


 UD


 S>

ld ^

to

DhW

202
YIELD

STRENGTH

US

HARDNESS

2024-T851


 O)


 00


 r-


 CO


 LO

<S> ^r


 n


 ru

cnX

CO

CO

Ul

Q

a

:

<x

•o

<u 4-> id JC c

+-> t/> CL) id +-> id O

T3 4- -O S_ O <1) ui uns

to c

Cl

cu E td O

c o

TD S_ -a

CT>

S_ 4- t- Sid cu

ia o 3

4-> to o

• id r—

to T3 >> > > a) cu to .c

-c <u 4-> • +->+-> JC -—

>

<LI *

-a cm

to -C C r—

*4- +J C Id H- *t— O • O S to C 4—

s_ id o

+j a» ai cc M S.TJ E

• o c E o a. 3 to o

CD 0) S- S- $_ 4- j= -a •d

+-> to c 3 •> 0>*r- Id a- >> C SI to

CU +J ss cu U «3 +-> >

4-> C S_ to T—

to •«— +-> id +>

a> cu o

T3 1— a. i— cu r— in a.

a) co >> id to

•r- |— JO cu >» 1 cn s_ -o c

HIM Cl •r- *> sz o +J P i—

+-> CM S- c <u

. o cu >

4- i- a. to cu O O CL) CU r4- S- S- c CL cu OT3.— cu o

to a» in s- c

•1- +-> CO CU i- ca 1— «-o

id S_ l to *r- CL CU «d- CU 4- E CN c c

O CU i— •«- O O CO CM i— O

olO

<U

s-

=5cn O

00 CO


 in


 
 n


 cu >HUJQ COHGCUZOHX b^COHH

203
YIELD

STRENGTH

US

CONDUCTIVITY

2024-T8S1

if)

n

cu

^r

«s>

C7> n

oon

in

o

<E

if

%

> - 2

O

Q

Z

o

o

00 CD OJ

>mujjq cnt-o'ujzof-i ^cn*-«

204

Figure

61.

Comparison

of

the

yield

strength

(0.2 %

offset)

vs.

conductivity

data

generated

with

the

data

from

soft

21

24-T851

reported

by

Petrak

and

Gunderson

(heavy

solid

and

dashed

lines,

representing

a

least

squares

mean

and

the

lower

90%

confidence

level,

respectively,

from

Ref.

12).
UTLIF1ATE

TENSILE

STRENGTH

US

HARDNESS

2024-T3S1

CT> O

00©

r-





 LD©

^r


 n

oo

O'

CO

COuQ

O'

<E

00 CO LO

IDt-CO ^COHH

205

Figure

62.

Comparison

of

the

ultimate

tensile

strength

vs.

hardness

data

generated

for

2024-T351

in

this

report

with

data

from

soft

2024-T351

as

reported

by

Petrak

and

Gunderson

(heavy

solid

and

dashed

lines,

representing a

least

squares

mean

and

the

lower

90

percent

confidence

level,

respectively,

from

Ref.

12).
ULTIMATE

TENSILE

STRENGTH

VS

CONDUCTIVITY

2024-T351

206

Figure

63.

Comparison

of

the

ultimate

tensile

strength

vs.

conductivity

data

generated

for

2024-T351

in

this

report

with

data

from

soft

2024-T351

as

reported

by

Petrak

and

Gunderson

(heavy

solid

and

dashed

lines,

representing a

least

squares

mean

and

the

lower

90

percent

confidence

level,

respectively,

from

Ref.

12).
YIELD

STRENGTH

US

HARDNESS

2024-T351


 (7>

00

rCO®

LO

<S> ^r®n

01

co

COuQ

<E

•a

tz S- »e rO <U +> 3

#a +-> -aO

a 4— •r— i

• o r—

to to o aj CM

to (/) sz r—

a> E -p c O >7 T3 l. > -a 4-

J- 4— <o c CU CO cu CO cc SZ •a JC

+J '—

c E • »o <o o

to -a c a> s- > o E 4- -c to - V -p s- to 0S +-> •r- <u cu >>

a; s -a i.

to c <o 'aj 4- •p 3 3 >

4- s- CD a~i- o o to +-> Q.T3 o

a)

C -p cu CM s_ to to CL • «o to O to cu a) —

•I— <0

r- i- JZ s- £ -P 4-> CO A

*-> (U p—

cn c Q_ cn cu e •r— c >

cu >1 •r- a>

s_ r* -Q 4-> r—

+J to SZ

to CO T3 cu cu 1

cu to

a

T3 1 +-> cu c

r— *0" s- s- 0J

CU CM o cl-a

•I—

O CL CU •f >>CM cu S- 4-

s- c

4- s- 0% o o o to CO o

4- •e cu c c p O -O r— •1— c

to CU LO r« cu

•r

-t-> CO CJ

S- tO l— -a ss- i cu OJ CL CU «=s- -C Cl E c CM to o cu O *o o o o>cm -a ao

UDa

s-3

cn

®®

®

co lo

®

®

co

>»-«LtJ*JO (DHOIUIZOI-Z ^CO»-i

207
YIELD

STRENGTH

US

CONDUCTIVITY

2024-T351

ru

^r




00

n

CD

cn

n

OJ

n


 n

oo

03

CD

o

<r

* >

o

3

Q

z:o

o

CD LO CO 03

LOHQ'UJZOI-X ^CD*-»

208

Figure

65.

Comparison

of

yield

strength

(0.2%

offset)

vs.

conductivity

data

generated

for

2024-T351

in

this

report

with

data

from

soft

2024-T351

as

reported

by

Petrak

and

Gunderson

(heavy

solid

and

dashed

lines,

representing

a

least

squares

mean

and

the

lower

90

percent

confidence

level,

respectively,

from

Ref.

12).
o

_o

o

to o

c

o

o

cu CO

OO

m

O'

c

O'

<

ICD

CD CD CD -• -• - • < < <

O

OCD CD CD

< < <

o_

c\J

o o O O o o o

°

o m O in o in o

rO rO CM CM

( 0o ) duuaj_ 6ui6\/- 9Jd

209

Figure

66.

0.635

cm

thick

plate

"pre-aged"

specimens

examined

by

means

of

TEM.

A

and B

refer

to

the

"pre-aging"

sequence

treatment.
Al 2 4 6 8 10 12 14 16 18 20

Weight Percentage Magnesium

Figure 67. Phase diagram of Al-Cu-Mg system showing the aluminum rich solvus

surface (40).

210
211

(b)

Electron

diffraction

pattern

from

(a).

[112]

zone.
CO

CD . -t-> CO E 4-> O •f—CL

Z>

C4— •r

•rO

U

-a CD e >> S- O 4-> U -r- CL e 1—

*0" •!— 1 t— 0 CO ' *1— E > CO e

CD E >. E •«- S-

•r— CD O CO

CD CD >

CL r— CD

CO

O +-> •r— O

-a +-> E

CD s- JC cO

O

O o. co E r—

CD TD 3 -r- cC

cr

CO -Q , CD

4-> CL O CO • ai •1— CO • s- O CD CO

•r— r- CD -O O *r- O T- S- E 4-> co

•1“ CD S-O

CD CO E

co S_ CL 3

CD M—O ,

+-> -Q

a CO E

•r“ E CD E

+-> 0 3-1-

i- •1— +J rO

cO CD-I- S- CL CD +-> OI

5- CO -O E 4-> •r— CD O CO O +-> O

CO

O

S- 2: s_ CD CD o. cn

CO -—' S-

•r— cO CO Q '— 1—

cr> co

CDs-3

CD

•r—

Li212

present

at

grain

boundaries.
ure 70. (a) Cu^Mn^Al^Q dispersoid particle.

(b) Microdiffraction pattern from (a). Dispersoid particle zone

is [020]; a-Al matrix zone is [110]. Matrix spots identified by arrows

213
Figure 71. Helical dislocations in directly quenched specimen (T4 condition).

All dislocations visible have same Burgers vector, 1 /2 [1 1 0]

Projection of Burgers vector is parallel to [200]. Weak beam

dark field image, g/3g, g = 200.

i

214
Figure 72. Structure after processing to T351 condition. Direct quench.

215
o

o O

® ® ®

o ® o

Figure

73.

(a)

Structure

after

processing

to

T851

condition.

Direct

quench.

(b)

Electron

diffraction

pattern

[001

]

a

zone.

(c)

Indexed

pattern

corresponding

to

(b).

Reflections

from

eight

S

phase

variants

are

present.

Spots

not

shown

arise

from

double

di

ffraction.
A

AOCTC -

500s

217

Figure

74.

(a-f)

Microstructures

of

specimens

given

sequence A

"pre-aging"

treatments

(low

magnification).

Region

of

light

contrast

around

precipitate

helps

identify

it

as

the S

phase.

Light

contrast

is

result

of

preferential

thinning

during

foil

preparation.
218

(Figure

74.

Continued)
V>V

V,

:

<D

(Figure

74.

Continued)
A

AOO°C

-

500s

220

Figure

75.

(a-f)

Same

as

Figure

74

but

at

higher

magnification.
221

(Figure

75.

Continued)
222

(Figure

75.

Continued)
3 1—

C\J o • 1 1 1 1 S- o aj — o +-> r— 4- CO 1—

tO e i—l o o -E E •!- r— +J

ai •+> o o E *o 1—

-O

•i- • +j

<_> CO E QJ qj o qj a S- 0.0 *i— i—

tO CO LO S- i

o o co e s- QJ

•rO 1—1 E

4-

O

>)

N 1.0 (O -

tO

O CO

X

"O

O I—

i— E o E <— S- OO +J o s- o to -Q -*-> 4- i—

E

to E E CO -a •r4J 1. >, E

lOC 0 11 to

S_ <D -M -Q

CD E +->

O QJ

4-> to +J

+J tO Q.

C tO to QJ *r- +J

$- e to •r— QJ 4-> O i~ CL

4-> T- CO •t— tO = -4-> o

+1 0)0

D Q) • 1- c to c S- Q-’t— S_ O a. •1-014- o to't-r

QJ 1 -t- +-> S- OJ T3 •!— *3 CL S- 2

CL E

qj = o cl

CO S— •r— o

to

C +-> -e 1 1 .e o co — a ai aj c — O r— o -

c/1

C LxJ •!— 00

QJ +J 1 1 3 tO r— ^

cr-~*i— o

tO QJ -O QJ o '— CO -— S_ 1—

co

P",

aj

s-3

01 r—

U_

223
IMi

to

cu

s-o

CD

<3

to c i—

H3 o r—

•r- o c +-> o ai <c i—

E +-> •r- C a a)

<u -r- s

Q_ S- I 1 to O i

C\J a; o

fO OO ra

to • • E '

' '

1 COO o •i— C i—

(O <+-

O rs_ I I I—

CD

C

s_ O) CO) s s-

•r— 4-> 1 1 ta .c +-> o +-> (O o to

•1- Q.I— aj 3 >—

c c o

a> o —

N

+-> -r- -

(O +-> toX

+->

O I—

•i

r mo S- Q- i_ O +J

•r- 4- r- ro O <4- 1 1 E

aj •!- S— "O to -o

Q. -f- c c <o

a> o cl

to S_

T a>

fO +-> -C +-> .c a to «3 atu c +-> r— O *r— (O LJ •!- Q.

+J •r— fC O / ./ N (— 0J IBJD 0) S- -—'— s- Q_

r->.

aj

sZ5CD

•I— U224
225

Figure

78.

(a)

G

phase

precipitate.

Sequence

A

"pre-aged"

400°C

for

500s.

(b)

Electron

diffraction

pattern

from

(a).

Matrix

zone

is

[211]

and

G

zone

is

[001].
Figure 79. Q-phase precipitate with dispersoid particle at center. Precipitate

lies on [111] plane which is parallel to plane of figure. Sequence A

"pre-aged" at 350°C for 20s.

226
-o -o O

cd CO

+-> s- oo o CD XJ Q- o •r— co c 5- r—

<o 4-5 o co © •i— cn

-a c

cn •i

c >5-0 • 1— i

c 2 co o o

_c o S- CO O) s_ c C5 CO CD co

-SC cn

Q. o CO

«3 E c

s- o o

cn .sc • 1— o cn

S- -a CD o c s_

•f— cd E 4-5 CO (O

--

—

Cl) x:

JO ,

4->

—

O•«— CD T3 +-> +-> r— S- O

CD fO

•r— Q. 4-

“O . .X v— X

S- o •r— m co S- 1 Q S- 4-> CD fO -a Q. E c CO

CO •r— e

-a *i

- ca +j CO — *t3 CD

4-> CO CO r— CD 4-> <u 4-> •r— •f— "3 CL 4- 4-5 •r—

•r—

O > 4-> Q. CD sz •r— S- cn CJ CL

•i— CD S- S- 9

CQ Q- CO O

co

a>

S-05

•f— U227

particle-precipitate

clusters

are

free

of

homogeneously

distributed

S'.

Sequence A

"pre-aged"

250°C

for

400s.
500r 0

at

disp <

a>ac

ajZJcr

<uto

s_o

4-

00

a)>

s-3uco

»o

aj

•i— • +-> +-> <a e

+-> aj £= E

O) +-> to ra

cu cu

S- 5- Q.4J

a;s_ =cd o c

4-> CD

<c

m E i aj aj -c s_ o a. go =

00

cus-

=3CD iZ

(Do) duel 228
229

treatments

(low

magnification)
B

350°C

230

(Figure

82.

Continued)
5

<

•*,

J 1

V 9

v;^.

Wi, T-'* * ••'••r-.kj.]

f . V/•i. /

* mi. ^

4 / «

231

(Figure

82.

Continued)
232

Figure

83.

(a-f)

Same

as

Figure

82

but

at

higher

magnification.
233

(Figure

83.

Continued)
»-5

. § fH

- ~ ^

f ' s ^ 'j />y ,"^*7>

s. £> <&idL

££ ’^

x234

(Figure

83.

Continued)
Figure 84. Dark field micrograph showing large heterogeneously nucleated S'

precipites at dispersoid particles and at dislocations (bands of

adjacent precipitates). Homogeneously distributed precipitates

are present in matrix except in vicinity of large heterogeneously

nucleated precipitates. Sequence B "pre-aged" at 300°C for 20s.

235
Figure 85. Optical micrograph

7200 minutes . (b)

of (a) sequence A "pre-aged" at

sequence B "pre-aged" at 400°C

400°C for

for 7200 minutes

i 236
1 i 1 1 o o o O o

o o o o o

in ro CM *—

(Do) dwoi

CO

237

Figure

86.

Schematic

representation

of

nucleation

curves

for

sequence

"pre-aging"

treatment.
238

Figure

87.

Specimens

after

sequence

A

"pre-aging"

treatment

at

350°

for

(a)

20s

and

(b)

109s.

Specimens

were

then

processed

according

to

T851

specifications.

Small

needle

shaped

precipitates

are

S'

formed

during

final

aging

treatment.

Many

S

phase

precipitates

are

surrounded

by

light

(thin)

regions.
«* b

Figure 88. Age-hardening response at room temperature (24-25°C) of "pre-aged"

2024 aluminum alloy specimens, a) Sequence A "pre-aging" treatment,

(b) Sequence B "pre-aging" treatment.

239
8UH ‘ssaupjDH

o

o

om O

o

o

o

o

o

o

c

£

a>E

I240

Figure

89.

Room

temperature

(24-25°C)

age-hardening

behavior

of

direct

quench

specimen.
b

Figure 90. Affect of sequence A "pre-aging" treatment on hardness, (a) Final

hardness after room temperature aging (T4). (b) Incubation hardness,

(c) Incremental change in hardness due to room temperature aging.

Data taken from Figures 88 and 89.

241
CVJ £

O'+3ococ

Q>Ocoo

242

Figure

91.

Incremental

change

in

hardness

asa

function

of

(Cu+Mg)

concentration

due

to

room

temperature

aging.

Data

taken

from

Beton

and

Rollason

(38)

Figure

2.

Solvus

temperatures

for

the

ot/(a+S)

boundary

are

indicated.
Figure 92. As-quenched and final hardness curves for Al-2.2:1 (Cu+Mg) alloys

from Hardy (35) Figure 7.

243
Figure 93. Effect of sequence B "pre-aging" treatments on hardness.

(a) Final hardness after room temperature aging (T4).

(b) Incubation hardness, (c) Incremental change in hardness

due to room temperature aging. Data taken from Figures 88 and 89.

244
2024-

T

ro ro ro O0S/ UJ ‘ xi 1 0013 A ONflOS

c

£

o

C5

<

UJ

<r

CL

245

Figure

94.

Variation

of

the

sound-wave

velocity

in

2024-T4

aluminum

alloy

as

a

function

of

"pre-aging"

time

at

different

"pre-aging"

temperatures

and

heat

treatment

sequences.
o

o

c

e

LlI 5

ho

z

o

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246

Figure

95.

Variation

of

hardness

in

2024-T4

aluminum

alloy

for

different

"pre-aging"

temperatures

and

heat

treatment

sequences.
2024-T4

>>0 to r- O) I— s- •O 3

4-> 1 2

c <u

•r- Q.

3 <U r— +> « _ ^ co I— c

I *r“CO

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S- C o.

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•r- O) +J S-

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to -i— 1C 4-> s- +J = r— CO 3 C

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ra

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cu

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o o o

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in

• W0/8P ‘N0llVnN311V OINOSVdinn

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in

ro

247

and

heat

treatment

sequences.
1

00 “

T“ "T~

CM T” O

”T“

00 T“

(0

-T“

<fr T”CVi i o

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C\J

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ro

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249

Figure

98.

Correlation

between

hardness

and

ultrasonic

attenuation

for

2024-T4

aluminum

alloy.

Dependence

on

"pre-aging"

treatment

sequence.
o

00

oo r^- to & in

aaH 4 ss3NaavH

250

ULTRASONIC

ATTENUATION,

dB/cm

Figure

99.

Correlation

between

ultrasonic

attenuation

and

hardness

for

2024-T351

aluminum

alloy

(sequence

A).
2024-

T4

C\J

00ro

CD

GO

N-»o

CD

CD

”

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CM

CD

CD

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GO

in

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to

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ro

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o

m

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251

V

Figure

100.

Correlation

curve

between

sound-wave

velocity

and

hardness

for

2024-T4

aluminum

alloy.
00

aaH ‘ssaNaavH

252

SOUND

VELOCITY,

m/sec

Figure

101.

Correlation

curve

between

sound-wave

velocity

hardness

for

2024-T351

aluminum

alloy.
2024

-T85I

00

ro

VO

L

o

VO

o o

in

8dH ‘SS3NaaVH

vom

ro

vo

Ld

>

a

o

ro

vo

253

Figure

102.

Correlation

curve

between

sound-wave

velocity

and

hardness

for

2024-T851

aluminum

alloy.
20V I %‘AllAllOnaNOO

254

Figure

103.

Eddy-current

conductivity

as

a

function

of

aging

time,

at

different

aging

temperatures,

of

2024

aluminum

alloy

(unstretched).
mOinOmO^JO co cooc

'£

LlI 2

I-o

z

QdH ‘SS3N0dVH

255

Figure

104.

Hardness

as

a

function

of

aging

time,

at

different

aging

temperatures,

of

2024

aluminum

alloy

(unstretched).
9UH ' SS3NQ9VH

incocoo N-r-uSio in o in o in o m in

ce

LlT 1

I-o

<

oo

256

Figure

105.

Eddy-current

conductivity

and

hardness

asa

function

of

aging

at

190

of

2024

aluminum

alloy

(stretched

and

unstretched).
3

5.0

m

co o

oo m

r8dH 4 SS3N0dVH

£ IS o

to m

in 8

o

o

c

E

lJ

2

h-CDz

o

<

257

Figure

106.

Eddy-current

conductivity

and

hardness

as

a

function

of

aging

time

during

aging

at

150

°C,

of

2024

aluminum

alloy

(stretched

and

unstretched).
L , INCHES

I 2 3 4 5 6

Figure 107. Predicted minimum yield strength in different thickness plates

under the "worst case" heat flow conditions for both T351 and

T851 using sequence A or B data.
MINIMUM

ULTIMATE

TENSILE

STRENGTH,

ksi

L, INCHES

I 2 3 4 5 6

Figure 108. Predicted minimum ultimate tensile strength in different

thickness plates under the "worst case" heat flow

conditions for both T351 and T851 using sequence A or B data.

259

MPa
L, INCHES

2 3 4 5 6

PLATE THICKNESS, L cm

Figure 109. Predicted minimum hardness in different thickness plates

under the "worst case" heat flow conditions for both T351

and T851 using sequence A or B data.

260
L , INCHES

5 10 15

PLATE THICKNESS , L cm

Figure 110. Predicted maximum conductivity in different thickness plates

under the "worst case" heat flow conditions for both T351

and T851 using sequence A or B data.

261
U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

1. PUBLICATION OR

REPORT NO.

2. Performing Organ. Report No. 3. Publication Date

SHEET (See instructions) 83-2669 April 1983

4. TITLE AND SUBTITLE

NbS: Processing/Microstructure/Property Relationships in 2024 Aluminum Alloy Plates

5. author(S) L. Ives, L. Swartzendruber, W. Boettinger, M. Rosen, 5. Ridder,

F. Biancaniello, R. Reno, D. Ballard and R. Mehrabian

6. PERFORMING ORGANIZATION (If joint or other than NBS, see Instructions)

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

7. Contract/Grant No.

S. Type of Report & Period Covered

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City, State. ZIP)

NASA

Washington, DC

10. SUPPLEMENTARY NOTES

J Document describes a computer program; SF-185, FIPS Software Summary, is attached.

plates

and the

11. ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bibliography or literature survey, mention it here) .... . , , nondestructive evaluation (NDE) using eddy-current conductivity and hardness

measurements form an essential part of the quality control of aluminum alloy

used for aerospace vehicles. The relationships between the NDE measurements

important mechanical properties are affected by a large number of variables

including: chemical composition, cast structure, ingot scalping, solution heat

treatment and quenching, mechanical working, and aging treatment. At the request of

the National Aeronautics and Space Administration, a number of these relationships

nas been explored for 2024 aluminum alloy. This work is a continuation of our previous efforts on 2219 aluminum alloy and is motivated by a concern that a nuniDer of improperly treated plates with "soft spots" may have been incorporated

into aerospace structures. A major result of our research has been a delineation of

wriicn alloy tempers and plate thicknesses are most likely to contain "soft spots

due to specific processing errors.

12. KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolons)

Aging; aluminum alloy; eddy-current conductivity; hardness; heat flow; microstructure

nondestructive evaluation; precipitation; processing; segregation; solution heat

treatment; ultrasonics.

'

13. AVAILABILITY

[X] Unlimited

[~~| For Official Distribution. Do Not Release to NTIS

~\ Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C.

20402.

[X] Order From National Technical Information Service (NTIS), Springfield, VA. 22161

14. NO. OF

PRINTED PAGES

268

15. Price

$ 22.00

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