~ 

NASA Contractor Report 178227 

8lETHODS FCR EOWGEE HETALZURGY B1-CU-f3g-X-X PUOYS Final Report (BcDconell-Douglas Unclas 

CONSOLIDATION PROCESSING 

PARAMETERS AND ALTERNATIVE 

PROCESSING METHODS FOR POWDER 

METALLURGY AI = CU = Mg = X = X ALLOYS 

I 

I K. K. SANKARAN 

MCDONNELL DOUGLAS ASTRONAUTICS COMPANY - ST. LOUIS DIVISION 

MCDONNELL DOUGLAS CORPORATION 

ST. LOUIS, MISSOURI 63166 

Contract NASI - 16967 

February 1987 

NASA 

National Aeronautics and 

Space Administration 

Langiey Research Center 

Hampton, Virginia 23665-5225 
NASA Contractor Report 178227 

CONSOLIDATION PROCESSING 

PARAMETERS AND ALTERNATIVE 

PROCESSING METHODS FOR POWDER 

METALLURGY AI = CU - Mg - X - X ALLOYS 

K. K. SANKARAN 

I MCDONNELL DO1 G AS ASTRO JA JTICS COMPAP 

I MCDONNELL DOUGLAS CORPORATION L ST. LOUIS, MISSOURI 63166 

I 

I 

I 

Contract NASI - 16967 

February 1987 

NASA 

National Aeronautics and 

Space Ad mi nist ration 

Langley Research Center 

Hampton, Virginia 23665-5225 

Y - ST. LOUIS DIVISION 
PREFACE 

This report presents the results of research performed from 27 September 

1982 through 29 August 1986 by the McDonnell Douglas Astronautics Company - 

St. Louis Division under the National Aeronautics and Space Administration 

Contract NAS1-16967 entitled "Consolidation Processing Parameters and Alternative Processing Methods for Powder Metallurgy A1-Cu-Mg-X-X Alloys." The 

intent of the study was to systematically investigate the relationships between the powder consol idation processing parameters and the microstructure and 

properties of the fully processed alloys, and to evaluate alternative alloying 

and consolidation methods which have the potential to reduce material and 

processing cost and enhance product quality. 

The research was performed in the Product and Materials Engineering Department, with Dr. K. K. Sankaran as the Principal Investigator. Novamet, a 

unit of INCO Alloy Products Company, was the principal sub-contractor, with 

Dr. S. J. Donachie and Dr. P. S. Gilman supervising all of the powder processing. Mr. D. L. Dicus of the NASA-Langley Research Center was the Technical Manager for the contract. 

i 
CONTENTS 

Page 

1 

2 

1. INTRODUCTION 

2. BACKGROUND 

3. RESEARCH OBJECTIVES AND APPROACH 4 

3.1 Objectives 4 

3.2 Approach 4 

3.2.1 Selection of Alloy Compositions 

3.2.2 Selection of Consolidation Processing Parameters 

5 

5 

4. EXPERIMENTAL PROCEDURES 7 

4.1 Powder Production 7 

4.1.1 Li-Containing Alloys 

4.1.2 Fe- and Ce-Containing Alloys 

7 

7 

4.2 Billet Consolidation Processing 7 

4.2.1 Conventional Consolidation of Li-Containing 

Alloys (Tasks I and 11) 

4.2.2 Conventional Consolidation of Fe- and 

Ce-Containing Alloys (Tasks I and 11) 

4.2.3 Containerless Vacuum Hot Pressing of LiContaining and Fe- and ce-Containing Alloys (Task 111) 

7 

8 

8 

4.3 Extrusion 10 

4.4 Post-Extrus i on Process i ng 11 

13 

14 

14 

18 

4.5 Characterization of the Microstructure and Properties 

5. RESULTS AND DISCUSSION 

5.1 Powder Characteristics 

5.2 Billet and Extrusion Characteristics 

5.2.1 

5.2.2 Chemical Composition of the Extrusions 

5.2.3 Density of the Extrusions 

5.2.4 Appearance, Microstructure, and Deformation Sub5.2.5 Microstructure and Sub-structure of the Solution 

Hydrogen Contents of the Hot-Pressed Billets and 

Extrusions 

Structure of the Extrusions 

Heat Treated Alloys 

18 

20 

21 

22 

25 

iii 
CONTENTS (continued) 

5.3 Influence of Post-Extrusion Processing on the Properties 

5.3.1 Age-Hardening Behavior 

5.3.2 Effect of Cold-Work on Tensile Properties 

5.3.3 Effect of Solution Heat Treatment Temperature on 

Tensile Properties 

5.3.4 Effect of Aging Conditions on Tensile Properties 

5.4 Influence of Vacuum Degassing Parameters on the Properties 

5.4.1 Task 1: Alloys Prepared by the Conventional 

Consolidation of Prealloyed Powder 

5.4.2 Task 11: Alloys Prepared by the Conventional 

Consolidation of Mechanically Alloyed Powder 

5.4.3 Task 111: Alloys Prepared by Containerless Vacuum 

Hot Pressing 

5.5 Influence of Alloying Approach on the Properties 

5.6 Influence of Consolidation Processing Method on the 

Properties 

5.7 Influence of Orientation on the Properties 

6. CONCLUSIONS 

7. REFERENCES 

8. APPENDIX 

iv 

27 I 

28 

30 I I 

31 

32 

32 

I 

32 I 

39 I 

45 i 

52 

56 

59 

62 

63 

65 
LIST OF FIGURES 

Figure 

1. 

2. 

3. 

4. 

5. 

6. 

7. 

8. 

9. 

10. 

11. 

12. 

13. 

14. 

15. 

16. 

17. 

18. 

Particle Size Distribution of Atomized A 

Powder 

Appearance of the Preal loyed Powder 

As-Solidif ied Microstructure of the Prea 

Microstructure of the As-Extruded Alloys 

-4Cu-lMg-1.5Li-0.2Zr 

loyed Powder 

(111) Pole Figure of the As-Extruded Alloys 

Microstructure of the Solution Heat Treated A1 loys 

(111) Pole Figure of the Solution Heat Treated Alloys 

Age-Hardening Curves for Alloys 3, 9, and 13 

Tensile Fracture Surface of Alloy 3 in the Artificially 

Aged Condition 

Tensile Fracture Surface of Alloy 6 in the Artificially 

Aged Condition 

Tensile Fracture Surface of Alloy 9 in the Artificially 

Aged Condition 

Tensile Fracture Surface of Alloy 12 in the Artificially 

Aged Condition 

Tensile Fracture Surface of Alloy 15 in the Artificially 

Aged Condition 

Tensile Fracture Surface of Alloy 17 in the Artificially 

Aged Condition 

Tensile Fracture Surface of (a) alloy 19 and (b) alloy 21 

in the Artificially Aged Condition 

Influence of Alloying Approach on the Properties of LiContaining Alloys Prepared by conventional consolidation 

Influence of Alloying Approach on the Properties of LiContaining Alloys Prepared by Containerless Vacuum Hot 

Pressing 

Influence of Alloying Approach on the Properties of Fe- and 

Ce-Containing Alloys Prepared by Conventional Consolidation 

Pase 

15 

16 

17 

23 

24 

26 

27 

28 

35 

38 

41 

44 

46 

49 

51 

53 

54 

55 

V 
LIST OF FIGURES (contl nued) 

Figure 

19. 

20. 

21. 

22. 

23. 

24. 

A1 . 

A2. 

A3. 

A4. 

A5. 

A6. 

A7. 

A8. 

Influence of Alloying Approach on the Properties of Fe- and 

Ce-Containing Alloys Prepared by Containerless Vacuum Hot 

Pressing 

Influence of Consolidation Processing Method on the 

Properties of Li-Containing Alloys Prepared from Preal loyed 

Powder 

Influence of Consolidation Processing Method on the 

Properties of Li-Containing Alloys Prepared from Mechanically 

A1 loyed Powder 

Influence of Consolidation Processing Method on the 

Properties of Fe- and Ce-Contai ni ng A1 1 oys Prepared from 

Mechanically Alloyed Powder 

Influence of Orientation on the Properties of the Alloys 

Influence of Orientation on the Notch Tension Strength 

Thermogravimetric Analysis Scans From Prealloyed 

A1-4Cu-lMg-1.5Li-0.2Zr Powder 

Thermogravimetric Analysis Scan From Prealloyed 

A1-4Cu-lMg-l.5Fe-0.75Ce Powder 

Thermogravimetric Analysis Scan From Prealloyed 

Al-l.6Fe-0.8Ce Powder 

Transmission Electron Micrograph of Alloy 3 in the 

Artificially Aged Condition 

Transmission Electron Micrograph of Alloy 6 in the 

Extruded Condition 

Transmission Electron Micrograph of Alloy 9 in the 

Solution Heat Treated Condition 

Transmission Electron Micrograph of Al.loy 15 Following 

Containerless Vacuum Hot Pressing (a) Bright Field and 

(b) Dark Field 

Transmission Electron Micrograph of Alloy 15 in the 

Extruded Condition 

Paqe 

56 

57 

58 

59 

60 

61 

66 

67 

67 

68 

69 

69 

70 

71 

vi 
LIST OF TABLES 

Page 

6 

10 

11 

12 

14 

18 

20 

21 

Table 

1. 

2. 

3. 

4. 

5. 

6. 

7. 

8. 

9. 

Consolidation Processing Parameters Initially Selected for Study 

Billet Consolidation Processing Parameters 

Extrusion Dimensions 

Strengthening Mechanisms for the Alloys 

Chemical Composition of the Prealloyed Powder 

Hydrogen Content of the Alloys 

Chemical Composition of the Extrusions 

Density of the Extrusions 

Influence of Cold Working by Stretching Prior to Aging 

on the Longitudinal Tensile Properties 30 

10. Influence of Solution Heat Treatment (SHT) Temperature 

on the Longitudinal Tensile Properties 31 

11. Influence of Aging Time at 463K (374OF) on the 

Long1 tudinal Tensile Properties 32 

Longjtudinal Tensile Properties of the A1-4Cu-lMg-1.5Li-0.2Zr 

Alloy Prepared by the Conventional Consolidation of 

Prealloyed Powder 

12. 

33 

13. Notch Tension Properties of the A1-4Cu-lMg-1.5Li-0.2Zr 

Alloy Prepared by the Conventional Consolidation of 

Prealloyed Powder 36 

Longitudinal Tensile Properties of the A1-4Cu-1Mg-1 .5Fe-0.75Ce 

Alloy Prepared by the Conventional Consolidation of 

Prealloyed Powder 

14. 

37 

15. Longitudinal Notch Tension Properties of the 

A1-4Cu-lMg-l.SFe-0.75Ce Alloy Prepared by the Conventional 

Consolidation of Prealloyed Powder 39 

16. Longitudinal Tensile Properties of the A1-4Cu-lMg-1.5Li 

Alloy Prepared by the Conventional 

Consolidation of Mechanically Alloyed Powder 40 

vi i 
LIST OF TABLES (continued) 

Tab1 e 

17. 

18. 

19. 

20. 

21. 

22. 

23. 

24. 

25. 

Notch Tension Properties of the A1 -4Cu-1Mg-1.5Li A1 loy 

Prepared by the Conventional Consolidation of Mechanically 

A1 loyed Powder 

Longitudinal Tensile Properties of the Al-4Cu-lMg-l.5Fe-0.75Ce 

Alloy Prepared by the Conventional Consolidation of 

Mechanically Alloyed Powder 

Notch Tension Properties of the A1-4Cu-lMg-l.SFe-0.75Ce Alloy 

Prepared by the Conventional Consolidation of 

Mechanically A1 loyed Powder 

Longitudinal Tensile Properties of the A1-4Cu-lMg-1.5Li-0.2Zr 

Alloy Prepared by the Containerless Vacuum Hot Pressing of 

Preal loyed Powder 

Notch Tension Properties of the A1-4Cu-lMg-1.5Li-0.2Zr 

Prepared by the Containerless Vacuum Hot Pressing of 

Preal loyed Powder 

Longitudinal Tensile Properties of the A1-4Cu-lMg-1.5L 

A1 loy 

Alloy Prepared by the Containerless Vacuum Hot Pressing of 

Mechanically Alloyed Powder 

Notch Tension Properties of the A1-4Cu-lMg-1.5Li Alloy 

Prepared by the Containerless Vacuum Hot Pressing of 

Mechanically A1 loyed Powder 

Longitudinal Tensile Properties of the A1-4Cu-lMg-l.5Fe-0.75Ce 

Alloy Prepared by the Containerless Vacuum Hot Pressing of 

Mechanically Alloyed Powder 

Notch Tension Properties of the A1-4Cu-lMg-l.5Fe-0.75Ce Alloy 

Prepared by the Containerless Vacuum Hot Pressing of 

Mechanically Alloyed Powder 

Page 

42 

42 I 

43 

45 

47 

48 

48 

50 

52 

viii 
1. INTROOUCTION 

Vacuum degassing is an important step in the consolidation processing of 

aluminum alloy powder into billets for hot working. Insufficient degassing 

may cause defects such as hydrogen-induced porosity and blistering which degrade the properties of the wrought alloys. While innovative processing techniques are being developed to improve the quality and reduce the cost of 

powder-processed aluminum a1 loys (1, 2), the successful near-term use of 

these alloys will depend upon improving the efficiency of vacuum degassing 

through a metallurgical understanding of the changes occurring during consolidation. To this end, the present effort is aimed at determining the influence 

of varying the vacuum degassing parameters on the microstructure and properties of powder-processed aluminum alloys. In addition, using alloys prepared 

by consolidating prealloyed rapidly solidified powder in sealed aluminum containers as the baseline, the study also attempts to determine the effectiveness of alternative processing techniques such as mechanical alloying for 

powder production, and containerless vacuum hot pressing for powder consolidation. The intent is to systematically investigate the relationships between 

the powder consol idation processing parameters and the microstructure and properties of the fully processed alloys, and to evaluate alternative alloying 

and consolidation methods which have the potential to reduce material cost and 

enhance product quality. 

1 
2. BACKGROUND 

Wrought aluminum alloys prepared from rapidly solidified powder have demonstrated a potential for developing superior combinations of properties compared with ingot metallurgy (IM) alloys (3). Such improvements are due to 

the microstructural refinements and the increased solid solubilities resulting 

from the rapid solidification. Examples of property enhancement include the 

better strength and stress corrosion resistance of alloys 7090 and 7091 compared with that of 1M 7XXX-series alloys (4), the improved strength and 

toughness of powder metallurgy (PM) 2XXX-series alloys over that of equivalent 

IM alloys (5), and the elevated-temperature stability of A1-Fe-X (X=Ce, Cr, 

Ho, or< V+Si) alloys not achievable in IM alloys (6-9). However, the use of 

such alloys has generally been limited by a lack of complete understanding of 

the relationships between processing and properties. Additionally, the manufacturing technology for the rolling of plate and sheet products from large 

billets is still under development, and the alloys are available principally 

only as forgings and extrusions. 

In the conventional consolidation processing of powder-processed aluminum 

alloys, the alloy powder is pressed into a porous compact, vacuum degassed, 

consolidated into a fully dense billet, and formed into a wrought product 

(10). The purpose of vacuum degassing is to dehydrate the hydrated oxide 

film on the powder surface and to liberate the moisture prior to full densification. During the diffusion of the liberated water through the porous cornpact, additional reaction with the powder surface may result in thicker oxide 

fllms and the evolution of hydrogen. Since the minimum temperature at which 

the alloy must be heated for complete dehydration is not known, degassing is 

commonly performed above the highest temperature to which the alloy will be 

subsequently exposed (the solution heat treatment temperature for precipitation-hardened alloys) (10). This approach eliminates potential problems 

such as porosity and surface blistering by avoiding the evolution of hydrogen 

subsequent to full densification. However, high degassing temperatures irreversibly coarsen the refined microstructures and degrade the properties, particularly of the dispersion-strengthened alloys. 

A recent study (11) conducted on alloy 7091 and an Al-Fe-Ce alloy observed the break-up and distribution of the oxide film during the hot pressing 

and subsequent hot working of the degassed billet. This study reported that 

the amorphous oxide film remaining on the surface of 7091 alloy powder was 

fully dehydrated following vacuum degassing. While these and other studies 

have recognized the importance of the consolidation processing steps in the 

total PM processing of aluminum alloys, information on the effects of varying 

the vacuum degassing parameters on the microstructure and properties is not 

generally available. Also, important systems such as the aluminum-lithium 

alloys which possess more complex and reactive surface oxides have not received attention. 

Recently, approaches such as mechanical alloying (12) and containerless 

vacuum hot pressing (13) have been developed as alternatives to powder preparation by the rapid solidification of a prealloyed melt and conventional 

2 
powder consolldation using sealed aluminum containers, respectively. The advantages of mechanical a1 loying include the incorporation of strengthening by 

oxides, carbides, and a high density of sub-structural defects. Containerless 

as leakage into the container and the need to cool and reheat the degassed 

billet, and has the potential to reduce production costs and enhance material 

quality. However, the influence of vacuum degassing parameters on the stability of the sub-structures and the resultant properties of the mechanically 

alloyed materials, and on the microstructures and properties of alloys processed by containerless vacuum hot pressing has not been studied. The potential effectiveness of employing these alternative approaches to improve the 

properties of PM aluminum alloys has also not been determined systematically. 

I vacuum hot pressing avoids the disadvantages of conventional processing such 

I 

I 

3 
3. RESEARCH OBJECTIVES AND APPROACH 

This study is a systematic investigation of the relationships between the 

consolidation processing parameters and the microstructure and properties of 

powder metallurgy (PM) A1-Cu-Mg-X-X alloys. The study also investigates the 

consolidation processing of a1 loys prepared by alternative powder preparation 

and consolidation techniques, and compares their characteristics with those 

prepared from atomized, rapidly solidified powder and consolidated by conventional means using sealed aluminum containers. The overall goal is to develop 

an understanding of the sensitivity of PM alloys to processing variations, and 

to explore simplified processing methods which can reduce production costs and 

enhance material quality. 

3.1 Objectives 

The objectives of this investigation are to (1) identify the critical 

parameters associated with the conventional consolidation processing of prealloyed A1-Cu-Mg-X-X alloy powder, and evaluate the effects of varying these 

parameters on the microstructure and properties of the fully processed alloys 

(Conventional consolidation processing is defined as the cold compaction of 

prealloyed powder followed by vacuum degassing and hot pressing in a sealed 

aluminum canister.), (2) investigate alternative means that do not employ 

fully prealloyed powder for producing A1-Cu-Mg-X-X alloy compositions, and 

evaluate the effects of these alternative means on the consolidation processing, and the homogeneity, microstructure, and properties of the fully 

processed alloys, and (3) investigate alternative consolidation processing 

methods for A1-Cu-Mg-X-X alloys, identify the critical parameters associated 

with these methods, and evaluate the effects of varying these parameters on 

the microstructure and properties of the fully processed alloys. 

3.2 Approach 

Alloys prepared by the conventional consolidation of atomized, rapidly 

solidified powder were used as the baseline for this study. Mechanical 

alloying was chosen as the alternative powder production method, and containerless vacuum hot pressing was chosen as the alternative consolidation 

method. Because complete dehydration of the surface oxides and the removal of 

moisture are critical to achieving the full potential of PM alloys, vacuum 

degassing parameters were chosen as being critical for both conventional and 

containerless consolidation. Consolidated billets were hot extruded and heat 

treated for evaluation. To achieve the program objectives, the effort was 

divided into the following three tasks: 

Task I: Conventional Consolidation of Prealloyed (PA) Powder 

Task 11: Conventional Consolidation of Mechanically Alloyed (MA) Powder 

Task 111: Consolidation by Containerless Vacuum Hot Pressing (CVHP) of PA 

and MA powder. 

4 
3.2.1 Selection of Alloy Compositions - Two alloys with nominal compositions 

of Al-4Cu-lMs-l.5Li-0.2Zr and A1-4Cu-lHg-l.5Fe-0.75Ce, respectively, were 

selected for this study. The program approach was to study these alloys prepared from both PA and MA powder under Tasks I and 11, and the PA Li-containing alloy and the MA Fe- and Ce-containing alloy under Task 111. 

The A1-4Cu-lMg-1.5Li-0.2Zr composition was selected on the basis of the 

beneficial effects of Li additions and the current emphasis on A1-Li alloys 

(14). In addition, Li-containing alloys are expected to possess complex and 

more reactive surface oxides compared with the other aluminum alloys. Among 

the various A1-Li alloys being developed, A1-Cu-Mg-Li alloys possess superior 

combinations of strength and toughness. The amounts of Cu, Hg, and Li in the 

alloy are such that their full dissolution is expected at the selected 

solution heat treatment temperature of 773K (932OF). The addition of Zr is 

based on the beneficial effects of Al3Zr in controlling recrystallization 

and enhancing properties. Because of the availability of oxides and carbides 

for sub-structural refinement, Zr was not included in the Li-containing alloy 

prepared by mechanical alloying. 

The A1-4Cu-lHg-l.5Fe-0.75Ce composition was selected on the basis of the 

known beneficial effects of adding Fe and Ce on the elevated temperature properties. Similar to the amounts of Fe and Ni in the 2618 aluminum alloy, the 

amounts of Fe and Ce were selected to provide strength enhancement at moderately elevated temperatures near 422K (3OOOF). The ratio of the amounts of Fe 

and Ce was based on the superior elevated temperature properties reported for 

an A1-8Fe-4Ce alloy (6). Elemental additions of Fe and Ce may not be feasible by mechanical alloying, since the aluminum may coat these elements 

during processing and prevent their homogeneous distribution. To achieve 

homogeneous alloying, the Fe- and Ce-containing alloys were to be mechanically 

alloyed using elemental Al, Cu and, Hg powder, and prealloyed, rapidly solidified A1-1.6Fe-0.8Ce powder prepared by atomization. 

3.2.2 Selection of Consolidation Processing Parameters - For each of the 

alloys under the three tasks, the vacuum degassing temperatures were chosen 

both above and below the solution heat treatment temperature of 773K (932OF). 

The intent was to determine the need for vacuum degassing above the highest 

post-consolidation processing temperature. For conventional consolidation 

(Tasks I and 11), the approach was to cold isostatically press the powder, encapsulate and seal in an aluminum alloy container, vacuum degas, hot press at 

675K (756OF) to full density, and remove the container prior to extrusion. 

For CVHP (Task 111), the approach was to degas the loose powder in a vacuum 

chamber which is itself part of the pressing die, and hot press at the same 

temperature to full density. A nominal billet diameter of 150 nun (6 in.) was 

chosen for both conventional and CVHP consolidation. A reduction ratio of 

16:l was chosen for extruding the billets at 600K (621°F) into bars of rectangular cross-section with an aspect ratio of 5:l. The 5:l aspect ratio is 

expected to prevent strengthening from axisymmetric deformation texture. 

Additionally, due to the influence of dispersoids in stabilizing the hotworked sub-structure, an extrusion ratio of 8:l with a 5:l aspect ratio was 

also planned for the Fe- and Ce-containing alloy under Tasks I and 11. The 

consolidation processing parameters targeted for this study are presented in 

Table 1. 

5 
TABLE 1. CONSOLIDATION PROCESSING PARAMETERS INITIALLY SELECTED FOR STUDY 

- TASK 

I 

111 

ALLOY 

PA A1 -4Cu-lMg-1.5Li -0.2Zr 

PA Al-4Cu-lMg-l.SFe-0.75Ce 

MA A1 -4Cu-1Mg-1.5Li 

MA A1-4Cu-lMg-1.5Fe-0.75Ce 

PA A1 -4th-1Mg-1.5Li -0.2Zr 

MA A1-4Cu-lMg-1.5Fe-0.75Ce 

VACUUM DEGASSING EXTRUSION 

TEMPERATURE TIME RATIO 

K OF 

750 891 

750 891 

790 963 

750 891 

750 891 

790 963 

750 891 

750 891 

790 963 

750 891 

750 891 

790 963 

750 891 

790 963 

750 891 

790 963 

h 

6 

12 

6 

12 

12 

6 

6 

12 

6 

12 

12 

6 

6 

6 

6 

6 

16 

16 

16 

8 

16 

16 

16 

16 

16 

8 

16 

16 

16 

16 

16 

16 
4. EXPERIMENTAL PROCEDURES 

4.1 Powder Production 

4.1.1 Li-Containinq Alloys - PA A1-4Cu-lMg-1.5Li-0.2Zr powder was produced by 

helium gas atomization at Valimet, Inc., Stockton, CA. During the collection 

of the atomized powder, a flash occurred in the system which filters the very 

fine fractions for disposal. The recovered powder, which was not affected by 

the flash, was screened through an 80 mesh (170 pm) sieve for further processing. 

MA A1-4Cu-lHg-1.5Li powder was produced by Novamet (a unit of INCO Alloy 

Products Company), Wyckoff, NJ at its Haskell, NY facilities. The powder was 

produced by a process which is proprietary to Novamet. After attrition, the 

MA powder was heated at 700K (8OOOF) under a vacuum to remove the residual 

stearic acid. The stearic acid is used as a process control agent to balance 

the welding and fracturing of the powder particles during attrition. This 

organic agent also supplies the carbon necessary for the formation of aluminum 

carbides in the alloy. 

4.1.2 Fe- and Ce-Containing Alloys - PA A1-4Cu-lMg-l.5Fe-0.75Ce and Al-l.6Fe0.8Ce Dowder were Droduced by flue gas atomization at the Alcoa Technical 

Center,' Alcoa Center, PA. The atomized powder was screened through an 100 

mesh (149 pm) sieve for further processing. 

MA A1-4Cu-lMg-l.5Fe-0.75Ce powder was produced by Novamet using PA Al1.6Fe-0.8Ce powder and elemental Al, Cu, and Mg powder. As discussed in the 

following paragraphs, problems during powder consolidation also necessitated 

the mechanical alloying of this powder from the respective elemental powders. 

4.2 Billet Consolidation Processing 

4.2.1 Conventional Consolidation of Li-Containing Alloys (Tasks I and 11) - 

Cold isostatic pressing (CIP) of PA A1-4Cu-lMg-1.5Li-0.2Zr powder was believed 

to present potential safety hazards, because in addition to the flash that had 

occurred during powder collection, it was also reported by Valimet that 52.5 

pct. of the atomized powder was smaller than 44 pm (-325 mesh size). Should 

leaks develop in the bag during the CIP of such powder, reaction of the powder 

with the pressurizing fluid may cause a hazardous situation. Therefore, it 

was decided to cold compact the PA powder and, for consistency, the MA powder 

uniaxially using mechanical means. However,, the first attempt to uniaxially 

press the PA powder in an extrusion press at Reactive Metals Inc., Ashtabula, 

OH, resulted in an explosion and a fire. Although the extrusion press liner 

was maintained at 644K (7OO0F), uniaxial pressing is accomplished in one 

minute or less and the powder itself is not heated while pressing. It is 

believed that under the pressure of the cold compaction, the reaction between 

the reactive powder surface and the entrapped air built up a large internal 

pressure that eventually blew off the welded container and exposed the 

powder. Possible ignition sources include the residual lubricant in the 

extrusion chamber, or the flame heating the extrusion press liner. 

7 
The problems discussed in the preceding paragraph necessitated the 

pressing of the Li-containing alloy powder under controlled laboratory conditions so as to prevent further hazardous incidents. Since laboratory facilities for the cold pressing of 150 mm (6 in.) diameter compacts were not 

readily available, it was decided to cold press both the PA and MA Licontaining powder into 114 mm (4.5 in.) diameter compacts at the Novamet faci1 i ties. 

Three batches each of appropriate amounts of the PA and MA powder were 

uniaxially cold compacted in 114 mm (4.5 in.) diameter mild steel containers 

to about 75 pct. of the theoretical density. Following this step, the mild 

steel containers were machined off and the compacts were loaded into 114 mn 

(4.5 in.) diameter 6061-A1 alloy containers fitted with evacuation tubes. The 

compacts were heated in a fluidized bed furnace and vacuum degassed using the 

parameters shown in Table 1. Following vacuum degassing, the billets were 

cooled, the evacuation tubes were crimped and cut, and the compacts were 

sealed by welding. The degassed and sealed compacts were then heated at 675K 

(756°F) for three hours and hot upset to full density at the same temperature 

in an extrusion press. The 6061-A1 containers were then machined off and the 

billets were ready for extrusion. 

4.2.2 Conventional Consolidation of Fe- and Ce-Containing Alloys (Tasks I and 

II) - To expedite powder processing, it was initially decided to uniaxially 

cold press both the PA and MA A1-4Cu-lMg-l.5Fe-0.75Ce powder using procedures 

similar to that attempted at Reactive Metals, Inc. for the Li-containing 

alloys. While three batches each of the PA and MA Fe- and Ce-containing 

powder were pressed without incident, the densities of these cold-pressed compacts were about 90 pct. of the theoretical rather than the intended 70-80 

pct. The pores are not inter-connected at 90 pct. density, and the vacuum 

degassing of these compacts was not considered feasible. Since the Fe- and 

Ce-containing powder is safer to handle than the Li-containing powder, it was 

again decided to CIP this powder into 150 mm (6 in.) diameter compacts for 

Tasks I and 11. A new batch of powder was prepared by mechanically alloying 

elemental Al, Cu, and Mg powder and PA Al-1.6Fe-0.8Ce powder. However, during 

the heating of this MA powder for removing the process control agent, failure 

of the vacuum system resulted in its contamination and caused its rejection. 

Since all of the PA A1-1.6Fe-0.8Ce powder had been consumed, the alloy was 

made by mechanically alloying elemental Al, Cu, Mg, Fe, and Ce powder for 

further processing. 

Three batches each of PA and MA powder (prepared from elemental powder) 

were cold isostatically pressed under a pressure of 207 MPa (30 ksi) at 

Dynamet Technology, Burlington, MA. The compacts, which were 140 m (5.5 in.) 

in diameter and about 80 pct. dense, were loaded into 6061-A1 alloy containers 

fitted with evacuation tubes. Using procedures similar to that used for the 

Li -containing a1 loys, the compacts were vacuum degassed in accordance with the 

parameters shown in Table 1, hot upset, and the 6061-A1 containers were 

machined off . 

4.2.3 Containerless Vacuum Hot Pressing of Li-Containing and Fe- and CeContaining Allovs (Task 111) - During the period in which problems were being 

experienced in the consolidation processing of alloys under Tasks I and 11, 

8 
two A1-4Cu-lMg-1.5Li-0.2Zr billets (using prealloyed powder) and two Al-4CulMg-1.5Fe-0.75Ce billets (using mechanically alloyed powder) were prepared 

under Task 111 by containerless vacuum hot pressing. The two A1-4Cu-1Mg1.5Fe-0.75Ce billets were consolidated from powder mechanically alloyed using 

elemental Al, Cu, and Mg powder and prealloyed A1-7.6Fe-0.8Ce powder. 

The only die available at Novamet for containerless vacuum hot pressing 

was designed for pressing 280 mm (11 in.) diameter billets. To press the four 

150 mm (6 in.) diameter billets, an aluminum tube was placed inside and concentric with the 280 mm (11 in.) die. The tube, whose inner diameter was such 

as to yield a billet of the desired diameter, was filled with the alloy powder 

to be consolidated. The remainder of the die was filled with commercial grade 

aluminum powder with appropriate demarcation using copper powder. The tube 

was withdrawn, and the entire assembly was heated in the CVHP chamber. After 

the interface between the commercial aluminum powder and the alloy powder 

reached the vacuum degassing temperature (determined' by prior calibration), 

the powder was degassed for the desired time and pressed. The 280 mm (11 in.) 

composite billets were then machined to produce the alloy billets of the 

appropriate diameter. The vacuum degassing parameters for these four billets 

were in accordance with that shown in Table 1. 

In addition to the above four billets, five other billets were also processed by CVHP. These five billets included one 114 mm (4.5 in.) diameter 

billet made from the PA A1-4Cu-lMg-1.5Li-0.2Zr powder, one 114 mm (4.5 in.) 

and one 150 mm (6 in.) diameter billet each made from the MA A1-4CU-lHg-1.5Li 

powder, and two 150 mm (6 in.) diameter billets made from the MA A1-4Cu-1Mg1.5Fe-0.75Ce powder prepared from fully elemental powder that had been used 

for billet pressing under Task 11. The first three of these five billets were 

added to provide a comparison between billets of different diameters, and between the PA and MA Li-containing materials consolidated using both conventional and containerless processing. The last two billets were added to provide a comparison with the alloys prepared earlier under this task using elemental Al, Cu, and Mg powder and prealloyed Al-l.6Fe-0.8Ce powder. These five 

billets were again prepared as composite billets using the 280 mm (11 in.) die 

and an inner aluminum tube of the appropriate diameter, followed by machining 

the 280 mm (11 in.) composite billets to produce the alloy billets of the 

appropriate diameter. The vacuum degassing parameters for the five additional 

billets were selected to provide relevant comparisons with the sixteen billets 

originally planned for the program. 

The consolidation processing parameters for the billets of the twenty-one 

alloys, numbered 1 through 21, are presented in Table 2. 

9 
- TASK 

TABLE 2. BILLET CONSOLIDATION PROCESSING PARAMETERS 

ALLOY ALLOY VACUUM DEGASSING HOT PRESSING BILLET 

Nm TEMPERATURE TIME TEMPERATURE DIAMETER 

K OF h K "F mm in. , 

I PA A1-4Cu-lMg-1.5Li-0.2Zr 1 750 891 6 675 756 114 4.5 

2 750 891 12 675 756 114 4.5 

3 790 963 6 675 756 114 4.5 

PA A1-4Cu-lMg-l.5Fe-0.75Ce 4 750 891 12 675 756 150 6.0 

5 750 891 12 675 756 150 6.0 

6 790 963 6 675 756 150 6.0 

I1 MA A1-4Cu-lMg-l.SLi 7 750 891 6 675 756 114 4.5 

8 750 891 12 675 756 114 4.5 

9 790 963 6 675 756 114 4.5 

MA A1-4Cu-lMg-l.5Fe-0.75Ce 10 750 891 12 675 756 150 6.0 

11 750 891 12 675 756 150 6.0 

12 790 963 6 675 756 150 6.0 

111 PA A1-4Cu-lMg-1.5Li-0.2Zr 13 750 891 6 750 891 114 4.5 

14 750 891 6 750 891 150 6.0 

15 790 963 6 790 963 150 6.0 

MA A1-4Cu-lMg-1.5Li 16 750 891 6 750 891 114 4.5 

17 750 891 6 750 891 150 6.0 

MA A1-4Cu-lMg-l.5Fe-0.75Ce 18 750 891 6 750 891 150 6.0 

(Using PA A1-1.6Fe-0.8Ce) 19 790 963 6 790 963 150 6.0 

MA A1-4Cu-lMg-1.5Fe-0.75Ce 20 750 891 6 750 891 150 6.0 

(Using Elemental Powder) 21 790 963 6 790 963 150 6.0 

4.3 Extrusion 

Based upon Novamet's experience with similar alloys, the extrusion temperature was increased from the planned 600K (621°F) to 644K (700°F). The 

billets were coated with Fel-Pro C 300 lubricant and preheated for three hours 

at 644K (700°F). The 114 mm (4.5 in.) diameter billets were extruded through 

a rectangular die with a nominal cross section of 51 mm (2 in.) by 13 mm (0.5 

in.) at a speed of 91 m/s (3.6 in./s) and a pressure of 483 MPa (70 ksi).The 

150 mm (6 in.) diameter billets, except those of alloys 4 and 10, were extruded through a rectangular die with a nominal cross section of 63 mm (2.5 in.) 

by 15 mm (0.6 in.). Billets of alloys 4 and 10 were extruded through a rectangular die with a nominal cross section of 102 mm (4 in.) by 25 mm (1 in.). 

10 
The extrusion speed and pressure were not recorded for the 150 mn (6 in.) diameter billets. The dimensions of the twenty-one extrusions are presented in 

Table 3. 

TABLE 3. EXTRUSION DIMENSIONS 

- TASK - ALLOY 

I PA A1-4Cu-lMg-1.5Li-0.2Zr 

PA A1-4Cu-lMg-l.5Fe-0.75Ce 

I1 MA A1-4Cu-lMg-1.5Li 

MA A1-4Cu-lMg-l.5Fe-0.75Ce 

I11 PA A1-4Cu-lMg-1.5Li-0.2Zr 

MA A1 -4Cu-1Mg-I. 5Li 

MA A1-4Cu-lMg-l.SFe-0.75Ce 

(Using PA Al-l.6Fe-0.8Ce) 

MA A1-4Cu-lMg-l.5Fe-0.75Ce 

(Using Elemental Powder) 

ALLOY 

NUMBER 

1, 2 and 3 

4 

5 and 6 

7, 8 and 9 

70 

11 and 12 

13 

14 and 15 

16 

17 

18 and 19 

20 and 21 

REDUCTION ASPECT 

RATIO RATIO 

mn (in.) 

16 4 

7 4 

19 4 

16 4 

7 4 

19 4 

16 4 

19 4 

16 4 

19 4 

19 4 

19 4 

51 x 13 

(2.0 x 0.5) 

102 x 25 

(4.0 x 1.0) 

63 x 15 

(2.5 x 0.6) 

51 x 13 

(2.0 x 0.5) 

102 x 25 

(4.0 x 1.0) 

63 x 15 

(2.5 x 0.6) 

51 x 13 

(2.0 x 0.5) 

63 x 15 

(2.5 x 0.6) 

51 x 13 

(2.0 x 0.5) 

63 x 15 

(2.5 x 0.6) 

63 x 15 

(2.5 x 0.6) 

63 x 15 

(2.5 x 0.6) 

4.4 Post-Extrusion Processing 

All of the alloys were solution heat treated at 773K (932°F) for one 

hour, quenched in water at room temperature, and age-hardened. To determine 

11 
the aging conditions for the alloys for evaluating their mechanical properties, attempts were made to determine the age-hardening behavior of one alloy 

each of the PA and MA Li-containing alloys processed by both conventional consolidation and containerless vacuum hot pressing (alloys 3, 9, 13, and 16) at 

room temperature, 433K (320"F), 463K (374"F), and 493K (428°F). Because of a 

defect in the extrusion of alloy 16, the age-hardening study of this alloy was 

not completed. Based on the observed age-hardening behavior of alloys 3, 9, 

and 13, all of the Li-containing alloys were aged both at room temperature for 

a minimum of seven days to stable hardness, and at 463K (374°F) for 16 hours 

to peak hardness. Also, since these aging conditions are similar to those 

used for the conventional 2XXX-series alloys, and the aging temperatures are 

not expected to influence the precipitation/coarsening of the Fe- and Cecontaining dispersoids, the same aging conditions were chosen for all of the 

PA and MA A1-4Cu-lHg-l.5Fe-0.75Ce alloys as chosen for the Li-containing 

a1 loys. 

To enhance the kinetics of precipitation and to increase the strength of 

the alloys, precipitation-hardened 2XXX-series aluminum alloys are commonly 

cold worked following solution heat treatment and prior to aging. Additionally, powder-processed aluminum alloy extrusions are also cold worked by 

stretching prior to aging so that the stresses induced during quenching from 

the solution heat treatment temperature can be relieved. The influence of 

cold working prior to aging on the age-hardening response of these alloys is 

determined by the respective precipitation mechanisms responsible for strengthening. Based upon the alloy composition and the alloying method, the alloys 

can be grouped, as shown in Table 4, into four different classes with each 

class representing a certain combination of strengthening mechanisms. To determine the need for cold working the alloys prior to aging, attempts were 

made to stretch one representative alloy from each of the four different 

classes. While alloys 3 and 6 were successfully stretched, the high strength 

of the MA alloys precluded the stretching of alloys 9 and 12. Based upon 

these results, except for stretching a few selected alloys to determine the 

influence on tensile properties, the mechanical properties of all of the 

alloys were determined without stretching. Additionally, to determine the 

influence of the solution heat treatment temperature, alloys 5 and 6 were 

solution heat treated at 753K (896OF) and 788K (959OF) and naturally aged. 

The aging time at 463K (374°F) was also varied for alloys 2, 5, 8, and 11. 

TABLE 4. STRENGTHENING MECHANISMS FOR THE ALLOYS 

ALLOYS - STRENGTHENING MECHANISMS 

1, 2, 3, 13, 14, and 15 Precipitation Hardening (PH) 

7, 8, 9, 16, and 17 

4, 5, and 6 

PH + OxideKarbide Dispersion 

Strengthening (O/C-DS) 

PH + Intermetallic Dispersion 

Strengthening (IM-OS) 

10, 11, 12, 18, 19, 20, and 21 PH + O/C-DS + IM-DS 
1 

4.5 Characterization of the Microstructure and Properties 

The chemical composition of the PA A1-4Cu-lMg-1.5Li-O.ZZr and A1-4Cu-1Hg1.5Fe-0.75Ce powder, and all of the twenty-one extrusions were determined from 

triplicate samples at the McDonnell Douglas Corporation. The chemical composition of the PA A1-4Cu-lMg-l.5Li-O.ZZr powder was also determined by Valimet 

and from duplicate samples by Homogeneous Metals, Inc., Clayville, NY, and 

that of the PA A1-4Cu-lMg-l.5Fe-0.75Ce powder was also determined by Alcoa. 

The chemical composition was determined by atomic absorption spectrometric 

analysis of solutions of the alloys in a mixture of nitric and hydrochloric 

acids. The hydrogen contents of the PA powder and all of the hot-pressed 

billets and extrusions were determined from duplicate samples by a fusion 

method at Leco, Inc., St. Joseph, MI. Since all of the MA powder had been 

consumed in billet processing, it was not available for characterization. The 

PA A1-4Cu-lMg-1.5Li-0.2Zr and A1-4Cu-lMg-l.5Fe-0.75Ce powder were also examined by optical and scanning electron microscopy. 

The densities of all of the extrusions were determined, and samples of 

selected alloys were examined by optical microscopy. The sub-structures of 

selected alloys were characterized at Lambda Research, Cincinnati, OH by determining the (111) pole figures by the back-reflection technique using the 

Schultz apparatus, and the back-reflection pin-hole patterns. 

The tensile properties at room temperature of all of the alloys in the 

longitudinal orientation and of selected alloys in the transverse orientation 

were determined in accordance with the ASTM Standard E8. Unless indicated 

otherwise, the tensile properties are based on the results of duplicate 

tests. Round sub-size specimens with threaded ends and a nominal reducedsection of 4 mn (0.160 in.) were used to determine the tensile properties in 

the longitudinal direction of all of the alloys, except alloys 4 and 10, 

following solution heat treatment at 773K (932°F) for one hour and natural 

aging to stable hardness or artificial aging at 463K (374°F) for 16 hours to 

peak hardness. Such specimens were also used to determine the tensile properties of alloys 3 and 6 which were stretched prior to aging, and the transverse 

tensile properties of alloys 3, 7, 14, 17, 18, and 19 following natural 

aging. However, the longitudinal properties of alloys 7, 14, 15, 16, and 17, 

and the transverse properties of alloys 3 and 7 could not be determined owing 

to failure in the threaded ends of the specimens. Round specimens with unthreaded ends and a nominal reduced-section diameter of 6 mn (0.25 in.) were 

then used to determine the longitudinal tensile properties of (i) alloys 4, 7, 

10, 14, 15, 16, and 17 following solution heat treatment at 773K (932°F) for 

one hour and either natural aging or artificial aging at 463K (374°F) for 16 

hours, (ii) alloys 2, 5, 8, and 11 following a similar solution heat treatment and aging at 463K (374°F) for six and 33 hours, (iii) alloy 2 following a 

similar solution heat treatment and aging at 463K (374°F) for 12 hours, and 

(iv) alloys 5 and 6 following solution heat treatment at either 753K (896°F) 

or 788K (959°F) for one hour and natural aging. 

To obtain a measure of the plane strain fracture toughness of the alloys, 

notched tension tests were conducted on selected alloys in both the longitudinal and transverse orientations. The tests were conducted in accordance 

with the ASTM Standard E602 using 12.7 mm (0.5 in.) diameter specimens. 

13 
5. RESULTS AND DISCUSSION 

As described in the preceding sections, a total of twenty-one extrusions 

representing varying combinations of alloy compositions, alloying approach, 

and consolidation parameters were prepared under the three tasks of the program. In this section, the composition, microstructure, and the age-hardening 

behavior of the alloys are discussed first. Following this, the effects of 

varying the vacuum degassing parameters on the properties of the alloys are 

discussed separately for each of the three tasks and are also compared with 

each other. 

5.1 Powder Characteristics 

The chemical compositions of the PA powder alloys are presented in Table 

5. The copper content measured by all of the sources of analysis is lower 

than the nominal in the PA powder samples of both the A1-4Cu-lMg-l.SLi-0.2Zr 

and A1-4Cu-lMg-1.5Fe-0.75Ce alloys, although, as reported later, it is closer 

to the nominal in all of the extrusions prepared from the PA powder. The 

reason for this observation is not clear. The amount of Li is higher than the 

nominal, but the amounts of the other elements are close to the nominal corn 

sition. As expected, the hydrogen concentration of the Li-containing a1 

powder is also significantly higher than that of the Fe- and Ce- contain 

a 1 1 oy powder. 

TABLE 5. CHEMICAL COMPOSITION OF THE PREALLOYED POWDER 

ALLOY SOURCE OF CHEMICAL COMPOSITION (WT.%) 

ANALYSIS 

0- 

OY 

n9 

Cu Mg Li Zr Fe Ce 

A1-4Cu-lMg-1.5Li-0.2Zr Valimet 3.87 1.08 1.85 0.22 -- -- 

Homogene- 3.64 1.00 1.74 0.25 .- -- 

~US Metals 

McDonnell 3.76 0.86 1.73 0.19 0.02 -- 

Doug 1 as 

A1-4Cu-lMg-l.SFe-0.75Ce Alcoa 3.8 1.0 - -- 1.3 0.7 

McDonnell 3.59 0.91 .- -- 1.51 0.74 

Doug 1 as 
The particle size distribution+at tk as-atomized A1-4Cu-lMg-l.SLi-0.2Zr 

powder, which was determined by Valimet, is plotted in Figure 1. The size 

distribution follows log-normal behavior, as is expected for metal powders 

gas-atomized under constant pressure (15). From the linear, least-squares 

fit, a geometric mean particle diameter of 40 prn was determined. As 

reported by Alcoa, about 85 pct. of the PA Al-4Cu-lMg-l.SFe-0.75Ce and 

Al-1.6Fe-0.8Ce powder particles were finer than 325 mesh (44 pm) size. 

M11-6453 

1 

90 

ao 

70 

60 

50 

40 40 50 70 100 150 200 

PARTICLE SIZE ( p m) 

FIGURE 1. PARTICLE SIZE DISTRIBUTION OF ATOMIZED 

AI - 4Cu - 1 Mg - 1.5 Li - 0.2 Zr POWDER 

Scanning electron photomicrographs of the PA Al-QCu-lMg-l.5Li -0.22r and 

A1-4Cu-lMg-l.SFe-0.75Ce powder samples are presented in Figure 2. Powder particles are predominantly spherical, which is typical of inert gas atomized 

powder. The larger size powder particles of the A1-4Cu-1Mg-1 .5Fe-0.75Ce alloy 

are somewhat elongated and irregular in shape. The surface of the powder particles of the Li-containing alloy is generally also rougher than that of the 

surface of the A1-4Cu-lMg-l.5Fe-0.75Ce alloy. The surface characteristics and 

the morphology of the powder particles of these alloys are determined by the 

atomization parameters which are proprietary to the powder producers. It is 

also observed that the size of a significant fraction of the powder particles 

of both the alloys is less than about 20 pm. 

Optical micrographs of samples of the PA A1-4Cu-lMg-l.SLi-0.2Zr and Al4Cu-lMg-l.SFe-0.75Ce powder particles are presented in Figure 3. lhe alloys 

exhibit a cellular-dendritic solidification microstructure with dendrite arm 

spacings on the order of 1 pm. Based on the correlation developed by Matyja 

et al. (16), a solidification rate of about lO5K/s is indicated for these 

a1 loys. 

Results of analytical characterization of powder by therniogravimetric 

analysis, mass spectrometry, and x-ray photoelectron spectroscopy suggested 

the absence of chemically bound water on the surface of PA A1-4Cu-lMg-1.5Li0.22r powder sample. 7hese studies, which were conducted independently, are 

summarized in the appendix. 

15 
ORIGINAL PAGE 1s 

OF POOR QUALITY 

11 -6454 

AI-4 Cu-1 Mg-1.5 Li-0.2 Zr 

AI4 Cu-1 Mg-1.5 Fe-0.75 Ce U 

10prn 

FIGURE 2. APPEARANCE OF THE PREALLOYED POWDER 

16 
11 -6455 

AI-4 Cu-1 Mg-1.5 Li-0.2 Zr 

AI-4 Cu-1 Mg-1.5 Fe-0.75 Ce U 

lop m 

FIGURE 3. AS-SOLIDIFIED MICROSTRUCTURE OF THE PREALLOYED POWDER 

17 
- 

5.2 'dllet" and kxtrhsion Characteristics 

5.2.1 Hydrogen Contents of the Hot-Pressed Billets and Extrusions - The 

hydrogen contents of the billets and extrusions are presented in Table 6. The 

hydrogen contents of the hot-pressed billets of the A1-4Cu-lMg-l.5Li-0.2Zr 

TABLE 6. HYDROGEN CONTENT OF THE ALLOYS 

- TASK ALLOY ALLOY VACUUM DEGASSING HOT PRESSING HYDROGEN 

NUMBER TEMPERATURE TIME TEMPERATURE CONTENT 

I PA A1 -4Cu-lMg-1.5Li -0.22r 1 

2 

3 

PA A1-4Cu-lMg-l.5Fe-0.75Ce 4 

5 

6 

11 MA A1-4Cu-1Mg-1.5Li 7 

8 

9 

MA A1-4Cu-lMg-l.5Fe-0.75Ce 10 

11 

12 

111 PA A1 -4Cu-lMg-1.5Li-0.2Zr 13 

14 

15 

.5Li 16 

17 

.5Fe-0.75Ce 18 

6Fe-0.8Ce) 19 

MA A1 -4Cu-lHg-1.5Fe-0.75Ce 20 

(Using Elemental Powder) 21 

MA A1 -4Cu-lMqMA A1 -4Cu-1 Mg - 

(Using PA Al-1 

K OF 

750 891 

750 891 

790 963 

750 891 

750 891 

790 963 

750 891 

750 891 

790 963 

750 891 

750 891 

790 963 

750 891 

750 891 

790 963 

750 891 

750 891 

750 891 

790 963 

750 891 

790 963 

h 

6 

12 

6 

12 

12 

6 

6 

12 

6 

12 

12 

6 

6 

6 

6 

6 

6 

6 

6 

6 

6 

K OF P PM 

BILLET EXTRUSION 

675 756 38.40 34.41 

675 156 29.56 25.69 

675 756 20.65 34.19 

675 756 0.96 2.44 

675 756 0.97 1.05 

675 756 1.92 3.18 

675 756 29.42 33.92 

675 756 26.83 31.48 

675 756 23.80 24.11 

675 756 5.65 6.15 

675 756 1.88 2.52 

675 756 2.60 3.30 

750 891 176.71 72.74 

750 891 91.93 29.75 

790 963 45.57 49.81 

750 891 62.29 140.00 

750 891 349.23 193.32 

750 891 183.17 58.81 

790 963 22.50 22.52 

750 891 7.29 11.96 

790 963 18.55 8.59 

(alloys 1, 2, and 3) and A1-4Cu-lMg-1.5Fe-0.75Ce (alloys 4, 5, and 6) alloys 

prepared by the conventional consolidation of PA powder are about an order of 

magnitude lower than that of the powder used in preparing these alloys. The 

hydrogen contents of the MA Li -containing alloys prepared by conventional consolidation (alloys 7, 8, and 9) are similar to those of the corresponding PA 

alloys (alloys 1, 2, and 3). Although mechanical alloying can be expected to 

18 
increase the hydrogen content of aluminum alloys, the effect is not observed 

in the Li-containing alloys due to their inherently higher hydrogen content. 

However, such an effect is observed in the Fe- and Ce-containing alloys 

(compare alloys 10, 11, and 12 with alloys 4, 5, and 6). 

For both the PA and MA Li-containing alloys prepared by conventional consolidation, the hydrogen content is lowered by increasing either the degassing 

time or the temperature (compare alloys 1 and 2, 1 and 3, 7 and 8, and 7 and 

9). It has been reported in a previous investigation (17) that the hydrogen 

concentration of vacuum degassed Li-containing a1 loy powder passes through a 

minimum with increasing degassing temperature. Above the degassing temperature corresponding to the minimum in the hydrogen concentration, more of the 

liberated hydrogen-bearing species is chemically fixed by Li than is physically removed. Such an observation was made in an A1-3Li-1 .5Cu-lMg-0.5Co0.2Zr alloy for which a degassing temperature of 723K (842°F) resulted in the 

lowest level of residual hydrogen. However, the lower concentration of Li in 

the present alloys is probably not sufficient for the chemical fixing of 

hydrogen, and the efficiency of hydrogen removal increases, although not significantly, with increasing degassing time or temperature. Such an effect is 

also observed in the PA Li-containing alloys prepared by containerless vacuum 

hot pressing (compare alloys 13 and 14 with 15). The reason for the increasing hydrogen content with increasing diameter of the MA A1-4Cu-lMg-1.5Li 

alloy billet (alloys 16 and 17) is not clear. 

The hydrogen contents of the Fe- and Ce-containing alloys are generally 

lower than those of the Li-containing alloys, and a systematic variation of 

the hydrogen content with varying degassing parameters is not observed for 

these alloys prepared by conventional consolidation (alloys 4, 5, 6, 10, 11, 

and 12). The MA Fe- and Ce-containing alloy powder prepared from partially 

prealloyed pobder appears to have a large initial hydrogen, and a significantly higher amount is removed by degassing at the higher temperature (compare alloys 18 and 19). 

Except for the Li-containing alloy billets prepared by the conventional 

consolidation of the PA or MA powder (alloys 1, 2, 3, 7, 8, and 9). the lack 

of a systematic variation of the hydrogen contents of the billets with varying 

degassing parameters is probably because these parameters were not varied over 

a wide enough range of values. While the two degassing temperatures were chosen to lie above and below the subsequent solution heat: treatment temperature, 

respectively, their proximity precludes wide variations in the degassing efficiency. Nevertheless, for those alloys for which a systematic variation is 

observed, the degassing efficiency, as measured by the hydrogen content, increases with increasing degassing temperature or time. Additionally, the 

higher hydrogen content of the billets of both Li-containing and Fe- and Cecontaining alloys processed by CVHP compared with that of the corresponding 

billets processed by conventional consolidation could be due to the composite 

billet approach used in containerless consolidation. 

The hydrogen contents of the extrusions of the alloys prepared by conventional consolidation (alloys 1 through 12) are generally similar to those 

of the corresponding hot-pressed billets. For the extrusions of the alloys 

prepared by containerless consolidation, the hydrogen content, except for 

alloy 16, is either similar to or lower than that of the corresponding billet. 

19 
5.2.2 Chemical Composition of the Extrusions - The chemical compositions of 

the twenty-one extrusions are presented in Table 7. The overall compositions 

of the alloys are close to the target. Except for the Cu content of alloy 8 

and the Zr content of alloy 1, specific deviations from the nominal composition are systematic. The deviations are the higher Li content (0.25 - 0.35 

pct.) of all of the alloys made from the PA powder (alloys 1, 2, 3, 13, 14, 

and 15), the higher Fe content (0.20 - 0.35 pct.) of all of the alloys made 

from the MA powder (alloys 10, 11, 12, 18, 19, 20, and 21), and the higher Ce 

content (0.13 - 0.14 pct.) of the alloys made from the MA powder that used the 

PA A1-1.6Fe-0.8Ce powder for alloying (alloys 18 and 19). 

TABLE 7. CHEMICAL COMPOSITION OF THE EXTRUSIONS 

ALLOY ALLOY CHEMICAL COMPOSITION (WT.%) 

NUMBER 

PA A1 -4Cu-1Mg-l.5Li -0.2Zr 1 

2 

3 

PA A1-4Cu-lMg-1.5Fe-0.75Ce 4 

5 

6 

MA A1-4Cu-lMg-1.5Li 7 

8 

9 

MA A1-4Cu-lMg-1.5Fe-0.75Ce 10 

11 

12 

PA A1 -4Cu-1Mg-1.5Li -0.2Zr 13 

14 

15 

MA A1 -4Cu-1Mg-1.5Li 16 

17 

MA A1 -4Cu-lMg-1.5Fe-0.75Ce 18 

(Using PA Al-l.6Fe-0.8Ce) 19 

MA A1-4Cu-lMg-1.5Fe-0.75Ce 20 

(Using Elemental Powder) 21 

cu Hg Li Zr 

3.90 1.02 1.85 0.40 

4.00 1.03 1.83 0.19 

3.92 1.01 1.79 0.19 

3.94 0.94 -- -- 

3.90 0.93 -- 

3.82 0.92 -- 

-- 

-- 

4.08 1.02 1.49 -- 

4.56 1.15 1.62 -- 

4.10 1.04 1.44 -- 

4.03 1.02 -- -- 4.13 0.99 -- 

3.97 0.99 -- 

-- 

-- 

3.92 1.02 1.77 0.20 

3.94 1.03 1.78 0.19 

4.03 1.06 1.83 0.20 

3.98 1.01 1.40 -- 4.08 1.03 1.45 -- 

4.02 0.97 -- -- 

4.06 0.97 -- -- 

-- 3.93 0.99 -- 

4.11 1.04 -- -- 

Fe Ce 

0.028 -- 

0.034 -- 

0.060 -- 

1.61 0.78 

1.61 0.80 

1.59 0.78 

0.049 -- 

0.022 -- 

0.022 -- 

1.81 0.77 

1.78 0.76 

1.71 0.74 

0.052 -- 

0.020 -- 

0.017 -- 

0.027 -- 

0.020 -- 

1.70 0.88 

1.72 0.89 

1.73 0.73 

1.83 0.78 

20 

L 
5.2.3 Density of the Extrusions - The measured density of the alloys are 

presented in Table 8. The density values of the alloys are consistent with 

the corresponding alloy compositions. The density of the Li-containing alloys 

is about 3.5 - 4.5 pct. lower than that of 2024-Al , while that of the Fe- and 

Ce-containing alloys is about 1.5 - 3.0 pct. higher. Except for the density 

of alloy 13, the measured values are also consistent with the deviations in 

the chemical compositions of the respective alloys. 

TABLE 8. DENSITY OF THE EXTRUSIONS 

- ALLOY 

g/cm3 

DENSITY 

1 b/in. 

PA A1-4Cu-lMg-1.5Li-0.2Zr 1 2.644 0.0955 

2 2.651 0.0957 

3 2.653 0.0958 

PA A1-4Cu-lMg-1.5Fe-0.75Ce 4 2.818 0.1018 

5 2.824 0.1020 

6 2.812 0.1016 

MA A1 -4Cu-lMg-1.5Li 7 2.673 0.0966 

8 2.667 0.0964 

9 2.676 0.0967 

MA A1 -4Cu-lMg-1.5Fe-0.75Ce 10 2.847 0.1029 

11 2.838 0.1025 

12 2.845 0.1028 

PA A1 -4Cu-1Mg-1.5Li -0.2Zr 13 2.670 0.0965 

14 2.646 0.0956 

15 2.640 0.0954 

HA A1 -4Cu-1Mg-1.5Li 16 2.668 0.0964 

17 2.669 0.0964 

MA A1-4Cu-lMg-l.5Fe-0.75Ce 18 2.844 0.1027 

(Using PA Al-l.6Fe-0.8Ce) 19 2.853 0.1031 

MA A1-4Cu-lMg-1.5Fe-0.75Ce 20 2.835 0.1024 

(Using Elemental Powder) 21 2.845 0.1028 

- 4.55 

- 4.30 

- 4.22 

+ 1.73 

+ 1.95 

+ 1.52 

- 3.50 

- 3.72 

- 3.39 

+ 2.78 

+ 2.45 

+ 2.71 

- 3.61 

- 4.48 

- 4.69 

- 3.68 

- 3.65 

+ 2.67 

+ 3.00 

+ 2.35 

+ 2.71 

21 
5.2.4 Appearance, Microstructure, and Deformation Sub-structure of the 

Extrusions - All of the extrusions exhibited hot tearing on the surface with 

the degree of hot tearing varying from slight to severe. Residual container 

feed was also observed in some of the alloys prepared by conventional consolidation. Blisters were observed in only one of the alloys (alloy 15). These 

were present only in the region near the end of push. Extrusions of the 

alloys 18, 19, 20, and 21 had a rough surface finish. A crack was also present over the entire length of the extrusion of alloy 16. 

Optical micrographs of selected alloys in the as-extruded condition are 

presented in Figure 4. The microstructures are typical of that of powderprocessed aluminum alloys in the hot-worked and highly overaged condition. 

Considerable microstructural refinement is observed. In all of the alloys, 

the size of the second-phase particles is no larger than about 10 pm. The 

microstructures of the Li-containing alloys (alloys 3, 7, and 15) exhibit a 

high volume fraction of 3 - 10 pm particles. The particles in the alloy 

made from the PA powder (alloys 3 and 15)) are etched to varying degrees by 

the Keller's etch. The smaller (<2 pm) particles that are not strongly 

etched are believed to be AlzCu, while the larger particles that are 

attacked strongly are probably ternary compounds containing Li (18). In the 

alloy made from the HA powder (alloy 7), all of the second-phase particles are 

strongly attacked suggesting that these are Li-containing phases. 

The microstructures of the extrusions of the Al-4Cu-lMg-l.SFe-0.75Ce 

alloy made from both the PA powder (alloy 6), and the MA powder that was made 

using the PA A1-1.6Fe-0.8Ce powder (alloy 19) show extremely fine and uniformly distributed second-phase particles. The volume fraction of these particles is higher in the latter alloy. In contrast, the presence of coarse 

particles in the alloy made by the mechanical alloying of elemental powder 

(alloy 12) suggests that complete and homogeneous alloying of the elemental Fe 

and Ce particles has not occurred. 

The (111) pole figures of the extrusions of alloys 2, 5, 8, and 11 showed 

similar deformation texture. The difference among these alloys was in the 

degree of preferred orientation, with alloys 2 and 5 (made from the PA powder) 

showing a higher degree of texture than alloys 8 and 11 (made from the MA 

powder). The pole figures of alloys 2 and 5, shown in Figure 5, are representative of hot-worked and fully recovered sub-structures. The maximum x-ray 

intensities for alloys 8 and 11 were also considerably lower than that for 

alloys 2 and 5, suggesting that, as expected, the carbides and oxides result 

in a finer sub-structure in the mechanically alloyed materials. 

Attempts to determine the degree of recrystallization of the extrusions 

of alloys 2, 5, 8, and 11 from pin-hole, back-reflection patterns were rendered difficult by the fluorescence of the alloys, and the relatively low 

diffracted x-ray intensity. Patterns of moderate quality that were obtained 

by using special x-ray films and extended exposure times revealed that these 

alloys possessed unrecrystallized microstructures. These results are also 

consistent with the microstructural observations and texture measurements. 
mIGINAL PAGE 19 

W WlOR OUALm 

11 -6456 

ALLOY 6 t ALLOY 3 

ALLOY 7 U ALLOY 12 

10 pm 

ALLOY 15 ALLOY 19 

FIGURE 4. MICROSTRUCTURE OF THE AS-EXTRUDED ALLOYS 

23 
FIGURE 5. 

24 

Ml l-6457 

(111) POLE FIGURE OF THE AS-EXTRUDED ALLOYS: 

(a) ALLOY 2 AND (b) ALLOY 5. CONTOUR VALUES 

ARE RELATIVE TO THE RANDOM INTENSITY. 
5.2.5 Microstructure and' Sub-Strk?ur$ of the Solution Heat Treated Alloys - 

Optical micrographs of the extrusions of the alloys which were solution heat 

treated at 773K -(932"F) for 1 hour are presented :n Figure 6. Solution heat 

treatment did not result in observable recrystallization of these alloys. 

Several observations can be made by comparing these microstructures with those 

of the same alloys in the as-extruded condition (Figure 4). 

In the PA A1-4Cu-lMg-l.SLi-O.ZZr alloy (alloys 3 and 15), full dissolution of the excess solute-rich phases is not achieved at 773K (932OF). 

The Li-content of all of the PA A1-4Cu-lMg-l.5Li-O.ZZr alloys is higher than 

the nominal which probably results in a total solute content exceeding the 

limit of solubility at the solution heat treatment temperature. Except for 

alloy 8, all of the MA Li-containing alloys have compositions that are close 

to the nominal, and, as observed in Figure 6, complete dissolution of the 

coarser, solute-rich phases is achieved in the MA A1-4Cu-lMg-1.5Li alloy 

(alloy 7). 

The microstructures of alloys 6 and 19 in the solution heat treated condition suggest that full dissolution of the second phases is not achieved at 

the selected solution heat treatment temperature. In the alloys made from the 

PA powder, the undissolved second phase particles in the Fe- and Ce-containing 

alloy (alloy 6) are finer than those in the Li-containing alloys (alloys 3 and 

15) and are also uniformly distributed. The size and distribution of the 

coarse, unalloyed elemental Fe and Ce particles in the Al-4Cu-lMg-l.SFe-0.75Ce 

alloy made by the mechanical alloying of elemental powder (alloy 12) are also 

unaffected by the solution heat treatment. 

The (111) pole figures of the extrusions of alloys 2, 5, 8, and 11 in the 

solution heat treated condition were similar to those in the as-extruded condition. The difference among the alloys was again in the degree of preferred 

orientation, with the alloys made from the PA powder exhibiting a higher 

degree of texture. The pole figures of alloys 2 and 5 in the solution heat 

treated condition, shown in Figure 7, are representative of hot-worked and recovered alloys. As with the alloys in the as-extruded condition, the maximum 

x-ray intensities from the alloys made from the MA powder (alloys 8 and 11) 

were also considerably lower than for the alloys made from the PA powder 

(alloys 2 and 5). Consistent with the microstructural observations and 

texture measurements, the pin-hole, back-reflection patterns of alloys 2, 5, 

8, and 11 in the solution heat treated condition revealed unrecrystallized 

microstructures. However, recrystallization of the alloys made from the MA 

powder cannot be ruled out, since the extremely fine grains typical of these 

alloys will not be identified by these techniques. 

25 
ORIGINAL PAGE 1s 

OF PooR QUALm 

11 -6458 

CT 

ALLOY 3 ALLOY 6 31 

L 

ALLOY 7 U ALLOY 12 

10 pm 

ALLOY 15 ALLOY 19 

FIGURE 6. MICROSTRUCTURE OF THE SOLUTION HEAT TREATED ALLOYS 
ED (4 MI 1-6459 

FIGURE 7. (111) POLE FIGURE OF THE SOLUTION HEAT TREATED ALLOYS: 

(a) ALLOY 2 AND (b) ALLOY 5. CONTOUR VALUES 

ARE RELATIVE TO THE RANDOM INTENSITY. 

5.3 Influence of Post-Extrusion Processing on the Properties 

The primary purpose of this study is to investigate the influence of 

varying the critical consolidation processing parameters on the properties. 

However, selecting the proper post-consolidation processing parameters is also 

27 
essential to obtaining meaningful mechanical properties. Thus, prior to 

investigating the influence of varying the vacuum degassing parameters on the 

mechanical properties of the a1 loys, their age-hardening characteristics were 

determined for selecting the appropriate aging conditions. Additionally, the 

influence of cold working by stretching prior to aging, and of varying the 

solution heat treatment temperature and the aging times at the selected peak 

aging temperature was also determined for a limited number of representative 

a1 loys. 

60 

5.3.1 Age-Hardening Behavior - The age-hardening curves of alloys 3, 9, and 

13 at room temperature, 433K (320°F), 463K (374OF), and 493K (428OF) following 

solution heat treatment at 773K (932°F) for 1 hour and water quenching are 

presented in Figure 8. 

- ROOM TEMPERATURE 

I I I I I Ill I I I I1111 I I I I Ill 

Ml l-6460 

ALLOY 0 

+9 

-8- 13 

70 

65 

a 

I 

85 - 

80 ALLOY - 

*3 

+9 

+ 13 

433K (32OOF) 

55 

100 101 1 02 103 ioo 101 1 02 103 

AGING TIME (h) 

100 

95 I 463K (374OF) - 

- 9ok 

m 

ALLOY n 

r *3 

+9 

-m- 13 65 

60 - I I I I I Ill I I I I IIIL 

10" 10' 1 

AGING TIME (h) 

AGING TIME (h) 

493K (428°F) 

90 

85 

ALLOY 

+-3 

t9 

+ 13 70 - 

65 - 

! 

60 

100 

* 

101 102 

AGING TIME (h) 

FIGURE 8. AGE-HARDENING CURVES FOR ALLOYS 3,9, AND 13 
I 

The hardening behavior of these alloys indicates that the Li-containing 

alloys made from the PA powder (alloys 3 and 13) exhibit aging kinetics similar to those of conventional alloys such as 2024-A1. Maximum hardness at 463K 

(374°F) is achieved after 16 hours, and stable hardness at room temperature is 

achieved after seven days. Also, longer aging times are required at 433K 

(320°F) than at 463K (374°F) to achieve peak hardness, although the value of 

the maximum hardness itself is higher at the lower aging temperature. The 

hardness of the Li-containing alloy made from the MA powder (alloy 9) is 

rather insensitive to aging. 

Extensive analytical transmission electron microscopic (TEM) studies are 

required to identify the precipitating phases that are responsible for the 

observed age-hardening in these alloys. The TEM studies (summarized in the 

appendix) that were independently conducted were 1 imited to observing the 

sub-structural features such as the sub-grain and grain structure and the dispersoids. Nevertheless, despite the uncertainties arising from the lack of 

TEM data, the known precipi tation-hardening behavior of the component binary 

and ternary systems was used to explain the age-hardening behavior of the more 

complex alloys of this investigation. 

It is observed that both of the Li-containing alloys made from the PA 

powder (alloys 3 and 13) exhibit similar behavior. The dual kinetics exhibited by these alloys, especially at the lower temperatures, may be attributed 

to a combination of the solid-solution hardening by Li, and the co-precipitation of GP zones/elt/e' and 6'(A13Li) which occur independently of each 

other. The solubility of Li decreases with increasing Mg, and an initial 

increase in the rate of A13Li precipitation may result from the Mg in the 

solid solution. With the onset of the precipitation of A12CuMg, the supersaturation of Li and hence the driving force for the precipitation of A13Li 

decreases. The initial drop in the hardness of alloy 13 at 433K (32OOF) and 

463K (374°F) results from reversion which is a partial breakdown of the GP 

zones that may have formed upon exposure to room temperature during sample 

preparation following solution treatment and quenching. Possible precipitation reactions in the Li-containing alloys made from the PA powder are: 

GP I+ elt+ e' + AI~CU 

Solid Solution- &I+ A12MgLi (A12CuLi) 

GP 1- S"-S' + A12CuMg 

AlzCu, AlzCuLi, and A12CuMg are non-coherent precipitates, and their 

formation on prolonged exposures at elevated temperatures causes overaging and 

decreases the hardness. 

In the A1-4Cu-lMg-1.5Li alloy made from the MA powder (alloy 9), dispersion strengthening by the oxides and carbides is expected to superimpose on to 

the mechanisms discussed for the all~ys made from the PA powder. The agehardening kinetics of alloy 9 indicates that the oxides and carbides impart 

considerable strengthening. The stability of these phases also results in a 

lower sensitivity of the hardness to the aging conditions. The dispersoids in 

the alloys made from MA powder are initially amorphous. At temperatures approaching 673K (752*F), the oxides and carbides crystallize to provide yA1203 and Al4C3, respectively (19). The solution heat treatment 

29 
temperature is sufficiently high to crystallize these phases, and the relatively low aging temperatures have little influence on dispersion strengthening. The rise and fall of the aging curve for alloy 9 is a result of the 

interaction between hardening by the oxides, carbides, and the solute-rich 

precipitates. While the variations are repeated in the same sequence at each 

aging temperature, the exact nature of the superposition of the strengthening 

mechanisms is uncertain. 

The strengthening mechanisms of the A1-4Cu-lMg-l.SFe-0.75Ce alloys made 

from the PA powder are expected to be similar to those of the equivalent 

2XXX-series alloys. The intermetallic Fe- and Ce-containing dispersoids are 

not expected to coarsen at the low aging temperatures. In these alloys made 

are also expected to be similar to those in the Li-containing alloys. 

I 

I 

from the MA powder, the dispersion strengthening by the oxides and carbides I 

I 

MPa ks i MPa ks i x 

Based upon the above observations, unless otherwise indicated, the mechanical properties of all of the alloys were determined after solution heat 

treatment at 773K (932°F) for 1 hour followed by water quenching and aging 

either at room temperature for a minimum of 7 days (natural aging, NA) or at 

463K (374°F) for 16 hours (artificial aging, AA). While higher peak hardness 

values were achieved for alloys 3 and 13 after aging for over 100 hours at 

433K (320"F), based on the commercial practice for 2024-A1 and the practicality to achieve maximum hardness within a reasonably short time, aging for 

16 hours at 463K (374°F) was chosen as representative of peak aging. 

I S - Stretched. Properties of alloy 3(S) in the AA condition are from one test. 

5.3.2 Effect of Cold-Work on Tensile Properties - As reported in Section 4.4, 

only alloys made from the PA powder could be stretched successfully. The influence of cold working by a nominal stretch of 1.5 pct., following solution 

heat treatment at 773K (932°F) for one hour and water quenching and prior to 

aging, on the mechanical properties of a representative Li-containing a1 loy 

(alloy 3, stretched to a permanent set of 1.5 pct.) and an Fe- and Ce-containing alloy (alloy 6, stretched to a permanent set of 1.75 pct.) is presented 

in Table 9. During the interval between quenching and stretching, the alloys 

were held at a temperature lower than 248K (-13°F). 

TABLE 9. INFLUENCE OF COLD WORKING BY STRETCHING PRIOR TO AGING 

ON THE LONGITUDINAL TENSILE PROPERTIES 

I 30 

ALLOY AGING 0.2% OFFSET ULTIMATE TENSILE ELONGATREATMENT YIELD STRENGTH STRENGTH - TION 

3 NA 441 64 54 5 79 6 

3(S) NA 448 65 51 0 74 5 

3 AA 393 57 462 67 4 

3(S) AA 469 68 538 78 5 

6 NA 193 28 352 51 15 

6 AA 179 26 290 42 16 

6(S) AA 2 69 39 331 48 13 

6(S) NA 248 36 352 51 10 
The influence of plastic deformation prior to age-hardening on the mechanical properties of aluminum alloys depends on the effects of the defects on 

precipitate nucleation. In A1-Cu-Mg alloys, the precipitation kinetics of GP 

zones is significantly enhanced by the defects. However, in Al-Cu-Li alloys, 

the response to cold work is dependent on the relative levels of Cu and Li 

which determine the type of the strengthening precipitates (20). The addition of Li to A1-Cu alloys is also known to modify the GP zones present in the 

binary alloy (21). Thus, in contrast with conventional alloys such as 

2024-A1, the lack of response to stretching in alloy 3 in the NA condition is 

perhaps due to the presence of Li-modified GP zones whose kinetics of precipitation is not enhanced by defects. However, in the AA condition, precipitation of the A12Cu(Mg, Li) phase is enhanced by cold working with a resultant 

increase of strength. For alloy 6 in both the NA and AA conditions, the influence of stretch is similar to that observed in alloys such as 2024-A1. The 

increase in the yield strength is more pronounced than the increase in the 

tensile strength. The strength enhancement is also accompanied by a decrease 

in the ductility. 

5.3.3 Effect of Solution Heat Treatment Temperature on Tensile Properties - 

As described earlier, all of the alloys were vacuum degassed both above and 

below the selected solution heat treatment temperature (773K, 932°F). This 

temperature was chosen so as to result in the full dissolution of the solute 

elements in the alloys of the nominal compositions. However, the very low 

strength of the Fe- and Ce-containing alloys made from the PA A1-4Cu-1Mg1.5Fe-0.75Ce powder indicated that the solute elements did not fully dissolve 

at 773K (932°F). To investigate the influence of the solution heat treatment 

temperature, the tensile properties of alloys 5 and 6 were determined in the 

longitudinal direction following solution heat treatment at both 753K (896°F) 

and 788K (959OF), and natural aging. The results of these tests, which are 

presented in Table 10, clearly show increasing strength with increasing 

solution heat treatment temperature. Increasing the solution heat treatment 

temperature increases the solubility of the solute elements thereby increasing 

the amount of solute available for precipitation hardening. Despite these 

results, all of the alloys were still solution heat treated at 773K (932°F) 

since this temperature lies between the two vacuum degassing temperatures and 

its choice was dictated once the billets were consolidated. 

TABLE 10. INFLUENCE OF SOLUTION HEAT TREATMENT (SHT) TEMPERATURE 

ON THE LONGITUDINAL TENSILE PROPERTIES 

ALLOY SHT 0.2% OFFSET ULTIMATE TENSILE ELONGAANDAGING TEMPERATURE YIELD STRENGTH STRENGTH TION 

CONDITION 

K OF MPa ksi MPa ksi % 

5 NA 753 896 172 25 331 48 16 

5 NA 773 932 207 30 365 53 16 

5 NA 788 959 234 34 400 58 15 

6 NA 753 896 165 24 324 47 16 

6 NA 773 932 193 28 352 51 15 

6 NA 788 959 2 34 34 393 57 13 

31 
5.3.4 Effect of Aging Conditions on Tensile Properties - To determine the 

influence of aging conditions on the mechanical properties, representative 

alloys (alloys 2, 5, 8, and 11) were aged for varying times at 463K (374°F) 

following solution heat treatment at 773K (932°F) and water quenching. The 

longitudinal tensile properties of these alloys, which are presented in Table 

11, indicate that the peak strength is achieved after aging for a period of 

time that is lower than that at which peak hardness is achieved. The effect 

is particularly pronounced in alloy 2, since the PA Li-containing alloys are 

strengthened solely by precipitation hardening. In the other alloys, which 

are additionally strengthened by the dispersoids, the effect is less pronounced due to the relative insensitivity of the dispersion strengthening to exposure at the aging temperature. Overaging by aging for 33 hours does not significantly affect the strength. 

I 

I 

TABLE 11. INFLUENCE OF AGING TIME AT 463K (374OF) 

ON THE LONGITUDINAL TENSILE PROPERTIES 

ALLOY AGING 0.2% OFFSET ULTIMATE TENSILE ELONGATIME YIELD STRENGTH STRENGTH - TION 

h HPa ksi MPa ksi x 

2 6 455 66 524 76 4 

2 12 462 67 524 76 3 

2 16 407 59 469 68 3 

2 33 421 61 483 70 3 

5 6 200 29 324 47 14 

5 16 186 27 296 43 15 

5 33 193 28 290 42 15 

8 6 531 77 600 87 2 

8 16 496 72 593 86 3 

8 33 496 72 593 86 2 

11 6 379 55 434 63 

11 16 352 51 421 61 

11 33 372 54 421 61 

6 

6 

6 

5.4 Influence of Vacuum Degassing Parameters on the Properties 

5.4.1 Task I: Alloys Prepared by the Conventional Consolidation of 

Prealloyed Powder - 

Al-4Cu-lMg-1.5Li-0.2Zr: The tensile properties of the A1-4Cu-1Mg1.5Li-0.2Zr alloy made by the conventional consolidation of PA powder are 

presented in Table 12. Increasing the vacuum degassing temperature 

results in a moderate increase in the strength of the naturally aged (NA) 

alloys, as observed by comparing the properties of alloys 1 and 3. However, a 

a1 loys. 

very sma 

32 

similar effect is not- observed- in the artificially aged (AA) 

Additionally, increasing the vacuum degassing time results in a 

1 increase in the yield stress of both the NA and AA alloys. 
ALLOY 

AND AGING 

CONDITION 

1 NA 

2 NA 

3 NA 

1 AA 

2 AA 

3 AA 

TABLE 12. LONGITUDINAL TENSILE PROPERTIES OF THE 

A1-4Cu-lHg-1.5Li-0.2Zr ALLOY PREPARED BY THE 

CONVENTIONAL CONSOLIDATION OF PREALLOYED POWDER 

VACUUM DEGASSING 0.2% OFFSET ULTIMATE ELONGATEMPERATURE TIHE YIELD TENS1 LE TION 

STRENGTH STRENGTH 

K OF h MPa ksi HPa ksi x 

750 891 6 400 58 503 73 5 

750 891 12 414 60 510 74 3 

790 963 6 441 64 545 79 6 

750 891 6 393 57 469 68 3 

750 891 12 407 59 469 68 3 

790 963 6 393 57 462 67 4 

Also, the ductility of neither the NA nor the AA alloy changes systematically 

with the variations in the vacuum degassing conditions. These observations 

suggest that the possible influence of varying the vacuum degassing parameters 

on the extent of powder degassing is not responsible for the observed changes 

in the mechanical properties. Additionally, although increasing the vacuum 

degassing temperature and/or the time decreased the hydrogen content of the 

billets of this alloy, the relatively low magnitude of this reduction indicates that, as mentioned earlier, these parameters were not varied over a wide 

enough range of values. The measured values of the tensile properties also 

suggest that the range over which the parameters were varied is above the 

minimum necessary for achieving reasonable properties. The increase in the 

strength of the NA alloy with increasing degassing temperature must then be 

due to a change in the sub-structure caused by the variation in the extent of 

elevated temperature exposure during vacuum degassing. 

Except for the changes in the vacuum degassing temperature, the processing sequence is identical for alloys 1 and 3. The differences between 

these alloys could then arise from the changes in the sub-structure due to the 

varying volume fraction of the Al32r dispersoids. The Zr remains in the 

supersaturated solid solution in the as-atomized powder. A certain amount of 

Zr precipitates as A13Zr during vacuum degassing, while the remainder precipitates during the thermal exposure prior to hot pressing the billet. Because 

of the increasing solid solubility of Zr in A1 with increasing temperature, 

the alloy degassed at the higher degassing temperature (alloy 3, degassed at 

790K, 963°F) can be expected to have more of this element available in solid 

solution for the precipitation of Al3Zr at the hot pressing temperature. 

Thus, increasing the vacuum degassing temperature will increase the volume 

fraction of Al3Zr that is formed during hot pressing at 675K (756°F) with a 

corresponding reduction of the sub-grain size following extrusion and solution 

heat treatment. It is also interesting to note that superplastic .Al-Cu-Zr 

alloys are aged at 643K (698OF) for the precipitation of Al3Zr that restricts the recrystallization and grain growth during elevated temperature 

forming (22). 

33 
The effect of decreasing the sub-grain size on the strength of these 

alloys can be illustrated by the following approximate and simple calculation. The diameter of the Al3Zr particles is typically 10 nm and their 

volume fraction is about 0.005 (22). Using the Zener relation, R = k.r/f, 

where R is the limiting radius of the growing grain, f is the volume fraction 

of the particle, and k is close to 1, increasing the f from 0.005 to 0.01 

would decrease the sub-grain diameter from 2 pm to 1 pm. By using the 

Hall-Petch relation, i 

U K . d-l I2 

where, for aluminum, the frictional stress (do) is 15.6 MPa (2.3 ksi) and 

the slope (K) is 0.068 MPa(m)l/2 (0.062 ksi(in.)1I2) (23), it is found 

that a decrease in the sub-grain diameter (d) from 2 pm to 1 pm increases 

the Hall-Petch strengthening (a) from 64 MPa (9.2 ksi) to 84 MPa (12.2 

ksi). While the Zener relationship has been derived for calculating the size 

of the recrystallized grains which grow in the presence of non-shearable particles, the above calculation does serve to illustrate the contribution to the 

flow stress by the sub-structural refinement due to the A13Zr dispersoids. 

The volume fraction of Al3Zr and the size of the sub-grains were not 

determined in this study. However, the microstructure of alloy 3 (Figure A4 

in the appendix) does suggest that the size of the sub-grains in these alloys 

can be expected to vary over the range that results in the observed variations 

in the strength of these alloys. 

In the alloys strengthened predominantly by shearable precipitates such 

as GP zones and/or dl in which deformation tends to concentrate on long and 

narrow planar slip bands, sub-grain boundaries are effective obstacles to further slip. In contrast, in the alloys strengthened predominantly by precipitates such as TI, SI, or e', sub-grain boundaries are not as effective. 

Thus, superposition of the strengthening by the sub-grains on to that from 

precipitation hardening is effective in the alloys hardened by shearable precipitates such as GP zones and 6', and not in the alloys hardened by the 

precipitation of phases such as TI, SI. or 8'. The changes in the tensile 

properties of the NA alloys (strengthened by shearable precipitates) with 

varying vacuum degassing temperature, and the lack of similar changes in the 

properties of the AA alloys (strengthened by phases such as TI, SI. or e') 

can thus be explained on the basis of the variations in the strengthening by 

the substructure. The relative insensitivity of the properties of the AA 

alloys is also consistent with the fact that these alloys are highly overaged 

(aged past the peak yield strength) and the dislocations can by pass the precipitates by the Orowan mechanism. Thus, slip would be more homogeneous and 

less planar thereby rendering the sub-grains rather ineffective. 

The ductility values of the alloys are somewhat lower than to be expected. As discussed earlier, the Li content of these alloys is higher than intended and results in an excess of the solute-rich phases with an accompanying 

loss of the ductility. The tensile fracture surface of alloy 3, shown in 

Figure 9, reveals a ductile fracture mode. The 3 - 5 pm undissolved particles are also similar to that revealed by optical metallography. Energy 

dispersive x-ray analysis identified these particles to be rich in Cu. 
ORtGlNAL PAGE 6 

OF FQOR QU11 -6461 

5 Pm 

FIGURE 9. TENSILE FRACTURE SURFACE OF ALLOY 3 IN THE 

ARTIFICIALLY AGED CONDITION 

35 
*=+ 

Attempts to determine the tensile properties of these alloys in the 

transverse direction were not successful due to the low ductility values and 

the resultant failure in the threaded ends of the specimens. To obtain a 

measure of the fracture toughness of these alloys, tension tests were conducted on notched specimens of a few selected alloys. The results of these 

tests are presented in Table 13. Included in the results are the values of 

the notch tension strength (NTS, the ratio of the maximum load to the initial 

cross-sectional area in the plane of the notch). For those alloys for which 

the smooth bar tensile properties were available, the notch yield ratio (NYR, 

the ratio of the NIS to the smooth bar yield strength), the notch strength 

ratio (NSR, the ratio of the NIS to the smooth bar tensile strength), and the 

calculated values of the plane strain fracture toughness (KIc) are also 

shown. 

TABLE 13. NOTCH TENSION PROPERTIES OF THE A1-4Cu-lMg-1.5Li-0.2Zr ALLOY 

PREPARED BY THE CONVENTIONAL CONSOLIDATION OF PREALLOYED POWDER 

ALLOY ORIENTA- NTS -- NY R -- NSR !iI - c 

_AND AGING - TION -- 

3 REATMENT 

MPa ksi MPa( m)’ j2 ks i ( in . )l2 

1 NA L 365 53 0.91 0.73 18.9 17.2 

3 NA L 352 51 0.80 0.65 18.5 16.8 

1 AA L 269 39 0.68 0.57 13.9 12.6 

3 A/\ L 262 38 0.67 0.57 13.5 12.3 

.- 3 NA T 165 24 

3 AA T 200 29 

- - 

- - -- - 

L - Longitudinal and T - Transverse 

K1C values were calculated from the relation: 

NY R - (4X) (1 - (X2/2))2 

where, X = (KIc/YS) (n/D)-ll2, YS is the srriooth bar yield strength, and 

D is the major diameter of the specimen (24). For materials with NYR 5 1.1, 

the calculated values of KIC agree to within 10 pct. bf the values obtained 

f rom more direct measurements. However, because of the approximate nature of 

the plasticity corrections, agreement is not good for sufficiently ductile 

a1 1 oys . 

The notch tension properties and the calculated values of KIC are not 

affected by the changes in the sub-structure of the alloys caused by the 

variations in the vacuum degassing parameters. In agreement with the tensile 

properties, the notch tension properties of the AA alloys are inferior to that 

of the NA alloys. The poor properties of the alloys in the transverse direction also results in a very low value of the NTS. 

36 
nA1:4Cu-lMq-1.5Fe -0.7!$&: The tensile properties of the A1 -4Cu-1Mg1 .5Fer0.75Ce alloy made by the conventional consolidation of the PA powder are 

presented in lable 14. As discursed in Section 5.3.3, increasing the solution 

heal treatment temperature increased the strength of this alloy thereby indicating a total solute content in excess of the maximum solid solubility at 

713K (937"l). It is possible that the total solute content exceeded the maximum 5olid solubility even at the higher degassing temperature (790K, 963OF). 

lhe excess second-phase particles present in these alloys is nominally composed of three parts - that which would dissolve at 713K (932°F) during solution 

heat treatment, that which would not dissolve at 773K (932°F) but would 

dissolve at 790K (963"F), and that which would not dissolve even at 790K 

(963°F). lhe extent of coarsening of that part of the second-phase particles 

attributable to the solute content in excess of the maximum solid solubility 

at 790K (963"F), is higher in the alloy degassed at this temperature (alloy 6) 

than in the alloy degassed at the lower temperature (alloy 5). This would 

explain the small reduction in the strength of alloy 6 compared with that of 

alloy 5, in either the NA or the AA condition. Additionally, the size of the 

dispersoids is possibly also not sufficiently small for stabilizing the substructure during hot working, and consequently changing the extrusion ratio 

does not change the tensile properties. However, while the extent of the dispersion and the sub-structure strengthening may be limited, it is superimposed 

upon strengthening by shearable precipitates. Consequently, the strength of 

the NA alloys is higher than that of the corresponding AA alloys. 

A I.. LOX 

- AND AGING 

CONDITION 

4 NA 

5 NA 

6 NA 

4 AA 

5 AA 

6 AA 

TABLE 14. LONGITUDINAL TENSILE PROPERTIES OF THE 

A1-4Cu-lMg-l.5Fe-0.75Ce ALLOY 

CONVENTIONAL CONSOLIDATION OF 

PREPARED BY THE 

PREALLOYED POWDER 

K "F h MPa ksi 

7 750 891 12 207 30 

19 750 s'9l I2 207 30 

19 790 963, 6 193 28 

7 750 891 12 193 28 

19 750 891 12 186 27 

19 790 963 6 179 26 

IJ LTIMATE -____ ELONGATENSILE -- T ION 

STRENGTH 

MPa ksi % 

3 12 5'1 14 

365 53 16 

352 51 15 

296 43 13 

296 43 15 

290 42 16 

The tensile fracture surface of alloy 6, shown in Figure 10, reveals a 

ductile failure mode. The absence of the coarser particles similar to that in 

the A1 -4Cu-lMg-1.5Li-0.22r alloys made from the PA powder results in =good 

values of the ductility. The notch tension properties, presented in Table 15, 

indicate a high notch tension strength for all of the alloys. Because of the 

very high plasticity, the KIC vdues were not calculated. 

37 
11 -6462 

U 

. 

FIGURE 10. 

U 

5 Clm . TENSILE FRACTURE SURFACE OF ALLOY 6 IN THE 

ART1 Fl Cl ALLY AGED CONDITION r, 
I. 

TABLE 15. LONGITUDINAL NOTCH TENSION PROPERTIES OF THE 

A1-4Cu-lMg-l.5Fe-0.75Ce ALLOY PREPARED BY THE 

CONVENTIONAL CONSOLIDATION OF PREALLOYED POWDER 

ALLOY - NTS 

AND AGING 

-- TREATMENT 

- NY R - NSR 

M Pa ks i 

6 NA 434 63 2.21 1.22 

6 AA 386 56 2.15 1.33 

From the above discussion, it is clear that the only effect of varying 

the vacuum degassing temperature and/or the time on the properties of A1-4CulMg-l.SLi-0.2Zr and A1-4Cu-lMg-1.5Fe-0.75Ce alloys made by the conventional 

consolidation processing of PA powder appears to be due to the changes in the 

sub-structure and dispersion strengthening. Together with the success of such 

techniques as depuretive degassing in achieving good combinations of properties in rapidly solidified 7XXX-series alloys (25), these observations 

suggest that vacuum degassing above the solution heat treatment temperature 

may not be necessary, and that once sufficient degassing is achieved, further 

improvements in the properties of powder-processed aluminum alloys are possible through a proper control of post-consolidation processing. 

5.4.2 Task 11: Alloys Prepared by the Conventional Consolidation of 

Mechanically Alloyed Powder - 

Al-4Cu-1Ms-1.5Li: The tensile properties of the A1-4Cu-lMg-l.SLi alloy 

made by the conventional consolidation of MA powder are presented in Table 16. 

In contrast with the behavior of the alloys made from PA powder, increasing 

the vacuum degassing temperature and/or the time decreases the strength of the 

alloy in both the NA and AA conditions. As with the Al-4Cu-lMg-l.SLi-0.2Zr 

alloy made from the PA powder, such behavior can be explained on the basis of 

the operating strengthening mechanisms. 

A significant part of the flow stress of mechanically alloyed materials 

is due to lhe dispersion strengthening by the oxides and carbides, and the 

Hall-Petch strengthening by the very small recrystallized grains. The oxides 

and carbides are produced during the manufacture of the powder, and any coarsening of these particles during processing will result in a reduction of the 

strengthening by these particles. The Orowan strengthening by the dispersoids, although not thermally activated, is strongly influenced by the size 

and spacing of the strengthening particles. The flow stress due to the dispersion strengthening varies inversely with the inter-particle spacing, and an 

increase of this spacing by the coarsening of the particles will cause a 

corresponding reduction in the strength. Thus, irrespective of the aging 

treatment, the strength of the A1-4Cu-lMg-l.SLi alloy made from the MA powder 

decreases with increasing vacuum degassing temperature and/or time. 

39 
TABLE 16. LONGITUDINAL TENSILE PROPERTIES OF THE 

Al-4Cu-lMg-l.SLi ALLOY PREPARED BY THE CONVENTIONAL 

CONSOLIDATION OF MECHANICALLY ALLOYED POWDER 

ALLU VACUUM DEGASSING 0.2% OFFSET ULTIMATE ELONGAAND AGING TEMPERATURE YIELD TENSILE TION 

- CONDITION -.STRENGTH STRENGTH 

K OF h MPa ksi MPa ksi % 

7 NA 750 891 6 655 95 676 98 3 

8 NA 750 891 I2 579 84 64 1 93 3 

9 NA 790 963 h 572 83 614 89 3 

7 AA 750 891 6 524 76 621 90 3 

8 AA 750 891 12 496 72 593 86 3 

9 AA 790 963 6 483 70 565 82 3 

Properties of alloy 8 AA are from one test 

The strength of the alloys in the AA condition is lower than those of the 

corresponding alloys in the NA condition. Generally, the strength of unrecrystallized, precipitation-hardened aluminum alloys is higher than that of 

recrystallized alloys, although grain size itself does not influence the 

strength. However, in the NA alloys, which are strengthened by shearable precipitates, the small sub-grains (in an unrecrystallized microstructure) such 

as in the A1-4Cu-lMg-1.5Li-0.2Zr alloys made from the PA powder (alloys 1, 2, 

and 3), and the small recrystallized grains such as in the A1-4Cu-lMg-1.5Li 

alloys made from the MA powder (alloys 7, 8, and 9) significantly add to the 

strength of the a1 loy. 

The tensile fracture surface of alloy 9 in the AA condition is shown in 

Figure 11. While the overall appearance of the fracture surface is oxidized 

and somewhat unusual, small (0.2 - 0.3 pm) dimples, possibly associated with 

the oxide and carbide dispersoid particles are observed at the higher magnification. While the failure mode is ductile (fracture results from ductile void 

nucleation and coalescence), the very high strength of these alloys results in 

low values of ductility. The notch tension properties and the calculated 

values of the KIC, shown in Table 17, are lower than the typical values for 

high-strength aluminum alloys. The properties of the AA alloys are also inferior to those of the NA alloys. 

40 
ORIGINAL PAGE iS 

OF POOR QUALITY 

U 

1Oprn 

FIGURE 11. TENSILE FRACTURE SURFACE OF ALLOY 9 IN THE 

ARTIFICIALLY AGED CONDITION 

41 
TABLE 17. NOTCH TENSION PROPERTIES OF THE A1-4Cu-lMg-l.SLi ALLOY 

PREPARED BY THE CONVENTIONAL CONSOLIDATION OF MECHANICALLY 

ALLOYED POWDER 

ALLOY ORIENTA- - NTS - NY R NSR 

AND AGING 

TRFATMFNT . . . -. . . . . - . . . 

MPa ksi 

KI - c 

7 NA L 345 50 0.53 0.51 17.1 15.5 

9 NA L 448 65 0.78 0.73 23.1 21 .o 

7 AA L 262 38 0.49 0.41 13.0 11.8 

9 AA L 338 49 0.70 0.60 17.5 15.9 

7 NA T 359 52 

7 AA T 248 36 

.- -- - -- -- -- .- -- 

L - Longitudinal and T - Transverse 

Al-4Cu-lMs-1.5Fe-0.75Ce: The tensile properties of the A1-4Cu-lMg-l.SFe0.75Ce alloy made by the conventional consolidation of MA powder are presented 

In Table 18. In contrast with the behavior of the Li-containing alloys, the 

strength of these alloys extruded using a reduction ratio of 19:l increases 

with increasing vacuum degassing temperature, as observed by comparing the 

properties of alloys 11 and 12. Additionally, increasing the extrusion ratio 

also reduces the strength significantly, as observed by comparing the properties of alloys 10 and 11. As expected, the strength of the alloy in the AA 

condition is also lower than that of the corresponding alloy in the NA condition. 

TABLE 18. LONGITUDINAL TENSILE PROPERTIES OF THE A1-4Cu-lMg-1.5Fe-0.75Ce 

ALLOY PREPARED BY THE CONVENTIONAL CONSOLIDATION OF MECHANICALLY 

ALLOYED POWDER 

ALLOY EXTRUSION 

AND AGING RATIO 

CONDITION 

10 NA 7 

11 NA 19 

12 NA 19 

10 AA 7 

11 AA 19 

12 AA 19 

Properties of alloy 

VACUUM DEGASSING 

TEMPERATURE TIME 

K 

7 50 

7 50 

790 

7 50 

7 50 

790 

"F h 

891 12 

891 12 

963 6 

891 12 

891 12 

963 6 

are , ,-om one test. 

0.2% OFFSET 

YIELD 

STRENGTH 

MPa ksi 

455 66 

400 58 

455 66 

407 59 

352 51 

414 60 

ULTIMATE 

TENSILE 

STRENGTH 

MPa ksi 

496 72 

469 68 

503 73 

455 66 

421 61 

462 67 

E LONGA- 

- TION 

x 

5 

9 

6 

4 

6 

5 

42 
Although the reasons for tbe,t$se&\Cesr, qriation in the mechanical properties of these alloys are not cler, an explanation can be sought in the effect 

of processing on the homogeneity and the distribution of the alloying elements. As discussed earlier, these alloys, which are made by the mechanical 

alloying of elemental powder, contain coarse undissolved particles both before 

and after solution heat treatment (refer to the microstructure of alloy 12 in 

Figures 4 and 6). Compared with alloy 11 which is degassed at 750K (891"F), 

degassing of alloy 12 at the higher degassing temperature (790K, 963°F) may 

result in a better homogenization of the elements and cause additional dispersion strengthening by the precipitation of the intermetallic dispersoids 

during processing subsequent to vacuum degassing. Thus, the strength of alloy 

12 is higher than that of alloy 11 in both the NA and AA conditions. However, 

the observed effect of the extrusion ratio cannot be explained. These alloys 

are strengthened by a combination of the fine grain structure, precipitates, 

and the oxide, carbide, and intermetallic dispersoids. Ihe complexity of the 

influence of processing on these strengthening features and the interaction 

among them, renders a prediction and explanation of the deformation and the 

mechanical properties of these alloys very difficult. 

Consistent with the good combination of strength and ductility, the 

alloys possess reasonable notch tension properties, as shown in Table 19. The 

tensile fracture surface of alloy 12 in the AA condition is also shown in 

Figure 12. The fracture surface reveals a significant amount of undissolved 

particles which together with the very fine oxide and carbide dispersoids 

results in a bimodal distribution of the dimples. 

TABLE 19. NOTCH TENSION PROPERTIES OF THE A1-4Cu-lMg-l.5Fe-0.75Ce ALLOY 

PREPARED BY THE CONVENTIONAL CONSOLIDATION OF MECHANICALLY 

ALLOYED POWDER 

ALLOY ORIENTA- NTS - NY R - NSR 

AND AGING TION 

TREATMENT 

MPa ksi 

12 NA L 524 76 1.15 1.04 

12 AA L 510 74 1.23 1.10 

.- -- 12 NA T 414 60 

12 AA T 414 60 -- -- 

MPa (m) l2 ksi ( i n . )1/2 

29.2 26.5 

28.9 26.3 

L - Longitudinal and T - Transverse 

From the above discussion, it is clear that, as with the alloys made from 

PA powder, the effect of varying the vacuum degassing parameters on the pro43 
11 -6464 

FIGURE 12. TENSILE FRACTURE SURFACE OF ALLOY 12 IN THE 

ARTIFICIALLY AGED CONDITION 

44 

10 pn 
i 

I 

I 

I 

I 

1 

i 

I 

I 

i 

I 

I ' 

i 

I 

~ 

I 

1 

I 

perties of the A1-4Cu-lMg-l.SLi and A1-4Cu-lMg-l.SFe-0.75Ce alloys made by the 

conventional consolidation of MA powder appears to be due to the changes in 

the sub-structure and dispersion strengthening. Also, mechanical alloying 

significantly increases the strength over that of .similar alloys made from PA 

powder. The latter observation is discussed in more detail in a subsequent 

section. 

5.4.3 Task 111: Alloys Prepared by Containerless Vacuum Hot Pressing - 

Al-4Cu-lMs-l.5Li-0.2Zr Alloy Made from PA Powder - The tensile properties 

of the A1-4Cu-lMg-l.5Li-0.2Zr alloy prepared by the CVHP of PA powder are presented in Table 20. In contrast with the alloys prepared by conventional consolidation, increasing the vacuum degassing temperature decreases the strength 

of these alloys, as seen by comparing the tensile properties of alloys 14 and 

15. In contrast with alloys 1, 2, and 3 which are hot pressed at 675K (756°F) 

following vacuum degassing, alloys 13, 14, and 15 are hot pressed at the same 

temperature at which they are vacuum degassed. The only thermal exposure 

following the CVHP of the billets of alloys 13, 14, and 15 is at 644K (7OOOF) 

for three hours followed by the extrusion. Thus, while AlgZr precipitation 

occurs in alloys 1, 2, and 3 during hot pressing, it does not occur similarly 

in alloys 13, 14, and 15. Consequently, in contrast with the alloys prepared 

by conventional consolidation, the strength of the alloys prepared by CVHP 

decreases with increasing vacuum degassing temperature. However, the effect 

of billet diameter on the tensile properties is not clear. Except for the 

yield strength of alloy 13, the strength of the alloy in the AA condition is 

also lower than that of the corresponding alloy in the NA condition. 

TABLE 20. LONGITUDINAL TENSILE PROPERTIES OF THE Al-4Cu-lMg-l.SLi-0.2Zr 

ALLOY PREPARED BY THE CONTAINERLESS VACUUM HOT PRESSING OF 

PREALLOYED POWDER 

mm in. K OF h MPa ksi MPa ksi % 

13 NA 114 4.5 750 891 6 434 63 531 77 3 

14 NA 150 6.0 750 891 6 455 66 531 77 5 

15 NA 150 6.0 790 963 6 434 63 496 72 5 

13 AA 114 4.5 750 891 6 448 65 503 73 3 

14 AA 150 6.0 750 891 6 427 62 490 71 4 

15 AA 150 6.0 790 963 6 393 51 441 64 4 

14 NA-T 150 6.0 750 891 6 365 53 414 60 2 

T - Transverse 

45 
ORlGtNAL PAGE I$ 

OF #KIR QUM 

The tensile fracture surface of alloy 15 in the AA condition (Figure 13) 

shows a ductile fracture mode, but as with alloys 1, 2, and 3 the excess 

solute-rich phases result in low values of the ductility in these alloys. The 

notch tension properties of these alloys, presented in lable 21, are in 

agreement with the corresponding tension properties. 

lop rn 

FIGURE 13. TENSILE FRACTURE SURFACE OF ALLOY 15 IN THE 

ARTIFICIALLY AGED CONDITION 
I 

i 

I '

I 

1 

1 

I 

I 

I 

I 

i I

I 

TABLE 21. NOTCH TENSION PROPERTIES OF THE A1-4Cu-lMg-1.5Li-0.2Zr ALLOY 

PREPARED BY THE CONTAINERLESS VACUUM HOT PRESSING OF 

PREALLOYED POWDER 

ALLOY ORIENTA- - NTS - NY R - NSR 

AND AGING - TION 

TREATMENT 

MPa ksi MPa (m) 1 12 ks i ( i n . ) 1 12 

14 NA L 324 47 0.71 0.61 16.6 15.1 

15 NA L 434 63 1.00 0.88 23.1 21 .o 

14 AA L 207 30 0.48 0.42 10.6 9.6 

15 AA L 283 41 0.72 0.64 14.6 13.3 

9.3 14 NA T 200 29 0.55 0.48 -- 

14 AA T 138 20 -- -- -- -- 

L - Longitudinal and T - Transverse 

Al-4Cu-lHs-l.SLi Alloy made from MA Powder - The tensile properties of the 

A1-4Cu-lMg-l.SLi alloy prepared by the CVHP of MA powder are presented in 

Table 22. While the yield strength of the alloys pressed from billets of 

varying diameter are nearly identical, the tensile strength of the alloy extruded from the larger diameter billet is significantly lower. The reason for 

this reduction in the tensile strength is not clear. The strength of the 

alloy in the AA condition is also lower than that of the corresponding alloy 

in the NA condition. The very high strength and the low ductility result in 

poor notch tension properties (Table 23), although the tensile fracture 

surface of alloy 17 reveals very fine dimples and a ductile failure mode 

(Figure 14). 

47 
TABLE 22. LONGITUDINAL TENSILE PROPERTIES OF THE A1-4Cu-lMg-l.SLi 

ALLOY PREPARED BY THE CONTAINERLESS VACUUM 

HOT PRESSING OF MECHANICALLY ALLOYED POWDER 

ALLOY BILLET VACUUM DEGASSING 0.2% OFFSET ULTIMATE ELONGAANDNG DIAMETER TEMPERATURE TIME YIELD TENSILE - TION 

- COND IT ION STRENGTH STRENGTH 

m in. K OF h MPa ksi MPa ksi % 

16 NA 114 4.5 750 891 6 634 92 710 103 3 

17 NA 150 6.0 750 891 6 641 93 648 94 1 

16 AA 114 4.5 750 891 6 524 76 641 93 2 

17 AA 150 6.0 750 891 6 510 74 586 85 2 

17 NA-T 150 6.0 750 891 6 565 82 579 84 1 

T - Transverse 

TABLE 23. NOTCH TENSION PROPERTIES OF THE A1-4CU-lMg-l.SLi ALLOY 

PREPARED BY THE CONTAINERLESS VACUUM HOT PRESSING OF 

MECHANICALLY ALLOYED POWDER 

ALLOY ORIENTA- - NTS - NY R - NSR !!E 

AND AGING - TION 

TREATMENT 

MPa ksi MPa( rn) 12 ks i ( i n . )1/2 

17 NA L 269 39 0.42 0.42 13.6 12.4 

17 AA L 228 33 0.45 0.39 11.6 10.5 

17 NA T 159 23 0.28 0.27 8.0 7.3 

17 AA T 172 25 -- -- . .- -- 

L - Longitudinal and T - Transverse 
ORffiMAL PAGE 13 

ff POOR QUALm 

11 -6466 

10pm 

FIGURE 14. TENSILE FRACTURE SURFACE OF ALLOY 17 

IN THE ARTIFICIALLY AGED CONDITION 

49 
Al-4Cu-lMs-l.5Fe-0.75Ce Alloy made from MA Powder - The tensile properties of 

the A1-4Cu-lMg-l.5Fe-0.75Ce prepared by the CVHP of MA powder are presented in 

Table 24. Unlike alloys 10, 11, and 12, alloys 18, 19, 20, and 21 are not hot 

pressed at 675K (756°F). Thus, additional dispersoid precipitation does not 

occur in the latter alloys subsequent to vacuum degassing. Consequently, the 

coarsening of the dispersoids with increasing vacuum degassing temperature 

causes a corresponding decrease of strength. Thus, the strength of alloys 18 

and 20 is higher than that of alloys 19 and 21, respectively. Additionally, 

the alloys prepared by using PA Al-l.6Fe-0.8Ce powder are more homogeneous and 

contain a higher amount of the strengthening dispersoids. Thus, the strength 

of alloys 18 and 19 is higher than that of alloys 20 and 21, respectively. 

The strength of all of these alloys is also higher in the NA condition than in 

the AA condition. The absence of the coarse unalloyed particles in the alloys 

prepared from PA A1-1.6Fe-0.8Ce powder is evident from a comparison of the 

tensile fracture surfaces of alloys 19 and 21 (Figure 15). Alloys 18 and 19 

also possess reasonable notch tension properties (Table 25). 

TABLE 24. LONGITUDINAL TENSILE PROPERTIES OF THE A1-4Cu-lMg-l.5Fe-0.75Ce 

ALLOY PREPARED BY THE CONTAINERLESS VACUUM HOT PRESSING 

OF MECHANICALLY ALLOYED POWDER 

ALLOY VACUUM DEGASSING 0.2% OFFSET ULTIMATE ELONGAANDNG TEMPERATURE TIME YIELD TENSILE TION 

CONDITION STRENGTH STRENGTH 

18 NA 

19 NA 

20 NA 

21 NA 

18 AA 

19 AA 

20 AA 

21 AA 

18 NA-T 

19 NA-T 

T - Transverse 

K OF 

750 891 

790 963 

750 891 

790 963 

750 891 

790 963 

750 891 

790 963 

750 891 

790 963 

h M Pa 

6 483 

6 448 

6 441 

6 407 

6 455 

6 421 

6 400 

6 37 2 

6 476 

6 455 

ksi M Pa 

70 545 

65 524 

64 51 0 

59 496 

66 503 

61 490 

58 476 

54 455 

69 51 7 

66 51 0 

ksi % 

79 5 

76 5 

74 5 

72 5 

73 4 

71 5 

69 5 

66 3 

75 2 

74 2 
ORIGINAL: PAGE 1s 

Ic]IE POOR QUALnV 

11 -6467 

FIGURE 15. TENSILE FRACTURE SURFACE OF (a) ALLOY 19 AND 

(b) ALLOY 21 IN THE ARTIFICIALLY AGED CONDITION 

51 
TABLE 25. NOTCH TENS" PROPERTIES OF THE Al-4Cu-lMg-l.SFe-0.75Ce ALLOY 

PREPARED BY THE CONTAINERLESS VACUUM HOT PRESSING OF 

MECHANICALLY ALLOYED POWDER 

ALLOY 

AND AGING 

TREATMENT 

ORIENTATION 

MPa ksi 

- NSR 

18 NA L 483 70 1 .OO 0.89 

19 NA L 545 79 1.22 1.04 

18 AA L 407 59 0.89 0.81 

19 AA L 531 77 1.26 1.08 

19 NA T 448 65 0.98 0.88 

19 AA T 469 68 -- -- 

L - Longitudinal and T - Transverse 

- 

-I K c 

MPa (m) l2 ks i ( in . )I 

25.7 23.4 

30.4 27.6 

21.6 19.6 

30.4 27.6 

24.2 22.0 -- -- 

5.5 Influence of Alloying Approach on the Properties 

The influence of the alloying approach (prealloying, mechanical alloying 

using prealloyed and elemental powder, and mechanical alloying using fully 

elemental powder) on the properties of the alloys is shown for various compositions and consolidation methods in Figures 16, 17, 18, and 19. In all 

cases, the dispersion strengthening due to carbides and oxides increases the 

strength of the alloy prepared from MA powder over that achieved in the 

corresponding alloy made from PA powder. However, the magnitude of this increase is influenced by the alloy composition, consolidation method, and heat 

treatment. 

In the Li-containing alloys prepared by conventional consolidation processing, the increase in the strength due to mechanical alloying is offset by 

the coarsening of the dispersoids at the higher vacuum degassing temperature. 

Additionally, as discussed earlier, the strength of the alloys made from the 

PA powder increases with increasing vacuum degassing temperature due to the 

decreasing sub-grain size. Thus, as shown in Figure 16, the incremental 

strengthening due to mechanical a1 loying decreases with increasing vacuum degassing temperature and/or time. The strengthening due to mechanical alloying 

is not accompanied by a significant loss of toughness, and for the alloy 

vacuum degassed at 790K (963OF) for 6 hours results in an increased value of 

KIC. Since the Li-containing alloy prepared by the CVHP of MA powder was 

not vacuum degassed at 790K (963"F), a similar comparison as made for the 

alloys prepared by conventional consolidation processing is not possible. 

In the Fe- and Ce-containing alloys prepared by conventional consolidation processing, the increase in the strength due to mechanical alloying increases with increasing vacuum degassing temperature (compare alloys 5 and 6 

52 
700 

600 

500 

h 

m 

n 

3 400 

F 

P 

6 

300 

200 

100 

0- 

500 I 8ot 

- 100 

90 

- 80 

E?! 

- -2 70 

- 

.- 

Y .- 

In 

60 

.e 50 

? 40 

- 0 

Y 

In 

- 

4 

- E 30 

20 

10 

0 

v) 

- 

400 

h 

E 

I 6 300 

z 

w 

v) 

F 

200 

100 

0- 

1 2 3 7 8 9 

ALLOY 

0.2% YIELD STRENGTH TENSILE STRENGTH 

ELONGATION Ea KIC 

Q7 

h 

C 

- :S$ 

z. 

Y 

0 

- .$ 

Y, 

I 

I- 

- g 

LT 

kJ 

- 

1 2 3 7 8 9 

ALLOY 

-4 

A 

- 3E 

- 2% 

z 

Q 

0 - 1-I 

w 

-0 

FIGURE 16. INFLUENCE OF ALLOYING APPROACH ON THE PROPERTIES OF Li - 

CONTAINING ALLOYS PREPARED BY CONVENTIONAL CONSOLIDATION 

53 
Ml l-6469 

120 

(a) NATURALLY AGED 

600 - 

500 

m 

0- 

E 400- I 

I- : 300 - 

v) 

- h 

E 

200 - 

100 - 

0- 

13 14 16 17 

(b) ARTIFICIALLY AGED 

N 

ALLOY 

0.2% YIELD STRENGTH TENSILE STRENGTH 

ELONGATION rn KIC 

-5 

-4 

h E 

-2 g 

3 

z 

-3 g 

-1 Lu 

-0 

13 14 16 17 

FIGURE 17. INFLUENCE OF ALLOYING APPROACH ON THE PROPERTIES OF 

Li - CONTAINING ALLOYS PREPARED BY CONTAINERLESS 

VACUUM HOT PRESSING 

54 
MI 1-6470 

ALLOY 

ou 

70 500 - 

- 

2 60 

400 - 7.- 

5 50 s 

5 300 - y 

c; 40 

.- c 

v 

m 0 v 

.- 

F 2. 4 

P 

Lu 30 LT 200 

Icn w 

- 

E 20 

100 - 

10 

0- 0 

4 5 - 6 10 11 12 

ALLOY 

0.2% YIELD STRENGTH TENSILE STRENGTH 

ELONGATION KIC 

FIGURE 18. INFLUENCE OF ALLOYING APPROACH ON THE PROPERTIES 

OF Fe - AND Ce - CONTAINING ALLOYS PREPARED BY 

CONVENTIONAL CONSOLIDATION 

55 
500 

400 

h 

m 

300 w 

I 

z 

6 

200 

v) 

100 

C " 

18 19 20 21 

M 

18 19 20 21 

1 1-6471 

-5 

-4 h 

$ 

-32 0 

w 

2 

-22 

0 -I 

LU 

-1 

-0 

ALLOY 

0.2% YIELD STRENGTH TENSILE STRENGTH 

ELONGATION N KIC 

FIGURE 19. INFLUENCE OF ALLOYING APPROACH ON THE PROPERTIES 

OF Fe - AND Ce - CONTAINING ALLOYS PREPARED BY 

CONTAINERLESS VACUUM HOT PRESSING 

with alloys 11 and 12). as shown in Figure 18. The reason for this increase 

can again be traced to the influence of the vacuum degassing temperature on 

the properties of the alloys made from the PA or the MA powder. For the 

Fe- and Ce-containing alloys prepared by the CVHP of MA powder, use of PA 

powder for the alloying of Fe and Ce results in a more homogeneous alloy with 

a consequent increase of strength over that of the alloy made from fully elemental powder (Figure 19). 

5.6 Influence of Consolidation Processing Method on the Properties 

The influence of the consolidation processing method on the properties of 

the alloys is shown in Figures 20, 21, and 22. The data, available for alloys 

with similar compositions, alloying approach, and vacuum degassing parameters 

indicates that consolidation by CVHP is beneficial for the Li-containing 

alloys prepared from the PA powder (Figure 20), increases the tensile strength 

of the Li-containing alloy prepared from the MA powder (Figure 21), and 

reduces the yield strength of Fe- and Ce-containing alloys prepared from the 

MA powder (Figure 22). However, these data are very limited, and preclude a 

general assessment of the influence of varying the consolidation processing 

method on the mechanical properties. 

56 
1 

500 

400 

- 

m 2 300 Y 

F 

9 2 200 

G 

100 

a 

80 

70 

- 

60 ,- c5 

.- c 

Y 

.- 2 50 

0 

.- A 40 

Y 

I 

. 6 30 

z 

UI 

cc 

u) 

$ 20 

10 

0 

(a) NATURALLY AGED 

Ml l-6472 

(b) ARTIFICIALLY AGED 

-3 

-2 = 0 

(!I -1 z 

2 

3 

Lu 

-0 

I 13 .- 1 13 

ALLOY 

0.2% YIELD STRENGTH TENSILE STRENGTH 

ELONGATION rn KIC 

FIGURE 20. INFLUENCE OF CONSOLIDATION PROCESSING METHOD ON 

THE PROPERTIES OF Li - CONTAINING ALLOYS PREPARED 

FROM PREALLOYED POWDER 

57 
Ml l -6473 

700 

600 

- 2 500 

I 6 400 

z 

W 

U 6 300 

200 

E 

100 

120 

(a) NATURALLY AGED 8oor 

- 

- 

- 

- 

- 

- 

- 

n- - 7 16 

(b) ARTIFICIALLY AGED 

- 

h 

s 

-3 - z 

-2 5 

c.!J 

-1 5 

-0 

7 16 

ALLOY 

0.2% YIELD STRENGTH TENSILE STRENGTH 

ELONGATION Ea KIC 

FIGURE 21. INFLUENCE OF CONSOLIDATION PROCESSING METHOD ON THE 

PROPERTIES OF Li -CONTAINING ALLOYS PREPARED FROM 

MECHANICALLY ALLOYED POWDER 

58 
500 

400 

m 2 300 

I 

IY 

s 

2 20c 

6 

1 oc 

C 

80 

70 

- 

'- 2 60 h 

C .- 

v 

.- 2 50 

.o_ Y 

.- A 40 

Y, 

I 

Y 

fn 

. 30 s 

$ 20 

u 

a: 

10 

0 

(a) NATURALLY AGED (b) ARTIFICIALLY AGED 

12 21 

ALLOY 

0.2% YIELD STRENGTH 

ELONGATION 

Ml l-6474 

12 21 

TENSILE STRENGTH 

FIGURE 22. INFLUENCE OF CONSOLIDATION PROCESSING METHOD ON THE 

PROPERTIES OF Fe - AND Ce - CONTAINING ALLOYS PREPARED FROM 

MECHANICALLY ALLOYED POWDER 

5.7 Influence of Orientation on the Properties 

The influence of orientation on the tensile properties of the alloys is 

shown in Figure 23. The transverse tensile properties of the Li-containing 

alloys are significantly lower than the properties in the longitudinal orientation. The properties of the Fe- and Ce-containing alloys are only slightly 

affected. As observed in a recent study (17), these results indicate deficiencies in the consolidation processing of the Li-containing alloys. The 

transverse properties are strongly dependent on the distribution of priorparticle surface oxide. Insufficient densification during the hot pressing of 

the degassed billet could result in reduced disintegration and a less uniform 

distribution of the oxide particles during extrusion with deleterious effects 

on the properties in the transverse direction. The oxide layer in Li-containing alloys can be expected to be thicker and more adherent than in the 

Fe- and Ce-containing alloys. Thus, the effect is more pronounced in the 

Li-containing alloys, thereby explaining the observed difference between the 

two types of alloys in the dependence of the tensile properties on the 

orientation. 

59 
Mi 1-6478 

700 

600 

500 

h 

m 8 400 

I 

IP F 300 

v) 

200 

100 

.- 

- 

- 

- 

- 

- 

- 

0- 

1001 I 

" 

14-L 14-T 17-L 17-T 18-L 18-T 19-L 19-T 

ALLOY 

0.2% YIELD STRENGTH TENSILE STRENGTH 

ELONGATION m KIC 

FIGURE 23. INFLUENCE OF ORIENTATION ON THE PROPERTIES OF THE ALLOYS. 

ALL ALLOYS NATURALLY AGED (L: LONGITUDINAL, T: TRANSVERSE) 

The influence of orientation on the NTS is shown in Figure 24. Except 

for alloy 7 in the NA condition, the NTS of all of the alloys in the transverse orientation is lower, by varying amounts, than in the longitudinal 

orientation. 

60 
I 

I 

I 

i 

1 

I 

I 

I 

I 

I 

I 

I 

I 

I 

i I 

I 

500 

400 

h 

m 

300 v 

I 

z 

6 

2 200 

v) 

100 

0 

500 

400 

h 

m 

300 v 

I 

!- 

Lu 

cc 200 Iv) 

s 

100 

0 

80 

70 

.- 3 60 

'I 

ICY 5 50 

E 

-v) 

Z 40 

v) 

z 

I 

0 

F 

Q 

. 30 

g 20 

10 

I 0 

80 

70 

.- 

v) 

Y, 60 

CY 5 50 

Z 40 

v) 

z 

I 

-F 

E 

-v) 

Q 

. 2 30 

2 

0 20 z 

10 

0 

M11-6479 

(b) TRANSVERSE 

3 7 12 14 17 19 3 7 12 14 17 19 

ALLOY 

(NATURALLY AGED) 

(a) LONGITUDINAL m (b) TRANSVERSE 

3 7 12 14 17 19 3 7 12 14 17 19 

ALLOY 

(ARTIFICIALLY AGED) 

FIGURE 24. INFLUENCE OF ORIENTATION ON THE NOTCH TENSION STRENGTH 

61 
6. CONCLUSIONS 

1. 

2. 

3. 

4. 

5. 

6. 

7. 

8. 

9. 

I The influence of vacuum degassing parameters on the properties of I 

powder-processed aluminum alloys is affected by the alloy composition, i 

and the alloying, consolidation, and post-consolidation processing 

approach. 

Vacuum degassing lowers the hydrogen content of powder-processed a1 uminum 

alloys, and its efficiency generally increases with increasing 

temperature. 

The hydrogen contents of powder-processed aluminum alloys, Li-Containing 

alloys, and the alloys fabricated by containerless vacuum hot pressing 

are higher than those of ingot-processed alloys, non-Li alloys, and the 

alloys fabricated by conventional consolidation, respectively. 

I 

The range over which the vacuum degassing parameters were varied in the 

present study was not sufficiently large to cause significant changes in 

the degassing efficiency. The observed changes in the properties of the 

alloys that were vacuum degassed under varying conditions resulted from 

the varying contributions to strengthening by the sub-structure and the 

dispersoids. 

Strengthening by the fine sub-grains, fine grains, and the dispersoids 

individually or in combination is more effective in the alloys containing 

shearable precipitates such as GP zones. Thus, the strength of all of 

the alloys is higher in the naturally aged condition than in the artificially aged condition. 

Hechanical alloying results in a significant increase of the strength 

over that of alloys of similar compositions prepared from prealloyed 

powder. The extent of this increase is determined by the influence of 

the vacuum degassing parameters on the sub-structural and dispersion 

s t rengtheni ng . 

The homogeneity of mechanically alloyed materials is improved by the use 

of prealloyed powder for the alloying of elements such as Fe and Ce. The 

strength of the alloys prepared by the mechanical alloying of prealloyed 

powder is also higher than those of the alloys prepared by the mechanical 

alloying of elemental powder. 

The mechanical properties of the alloys in the transverse orientation are 

lower than in the longitudinal orientation. This effect is more pronounced in the Li-containing alloys. 

Among all of the alloys that were prepared for this study, the Fe- and 

Ce-containing alloys prepared from mechanically alloyed powder possess 

better combinations of strength, ductility, and toughness. 

62 
7. REFERENCES 

1. J. W. Simon, Direct Processing of Aluminum Powder Metallurgy Alloys, 

Report No. AFWAL-TR-84-4144 on Air Force Contract No. F33615-82-C-5132 

(December 1984). 

2. J. M. Fitzpatrick, R. E. Lewis, and D. D. Crooks, Direct Processing of 

A1 umi num Powder Metal 1 urgy A1 1 oys, Report No. AFWAL-TR-84-4176 on Ai r 

Force Contract No. F33615-82-C-5023 (Apri 1 1985) . 

3. 3. R. Pickens and E. A. Starke, Jr., The Effect of Rapid Solidification 

on the Microstructures and Properties of Aluminum Powder Metallurgy 

Alloys, in Rapid Solidification Processing: Principles and Technologies, 

- 111,R. Mehrabian, ed. (National Bureau of Standards, Gaithersburg, MD, 

1982), p. 150. 

4. T. E. Tietz and I. 6. Palmer, Advanced P/M Aluminum Alloys, in Advances 

in Powder Technology, G.Y.Chin, ed. (American Society for Metals, Metals 

Park, OH, 1982), p. 189. 

5. D. J. Chellman, Development of Powder Metallurgy 2XXX Series A1 Alloy 

Plate and Sheet Materials for High Temperature Aircraft Structural Applications, NASA Contractor Report 172521 on Contract No. NASl-16048 (Apri 1 

1985). 

6. R. E. Sanders, Jr. and G. J. Hildeman, Elevated Temperature Aluminum 

A1 loy Development, Report No. AFWAL-TR-81-4076 on Air Force Contract No. 

F33615-77-C-5086 (June 1981 ) . 

7. R. Ray, Advanced Powder Metallurgy Aluminum Alloys via Rapid Solidification Technology, NASA Contractor Report 172343 on Contract No. NAS1-17578 

(May 1984). 

8. C. M. Adam, R. G. Bourdeau, and J. W. Broch, Application of Rapidly Solidified Alloys, Report No. AFWAL-TR-81-4188 on Air Force Contract No. 

F33615-76-C-5136 (February 1982). 

9. 0. J. Skinner, R. L. Bye, D. Raybould, and A. M. Brown, Dispersion 

Strengthened A1-Fe-V-Si Alloys, Scripta Metallurgica 20, 867 (1986). 

10. 3. R. Pickens, Aluminum-Powder Metallurgy Technology for High-Strength 

Applications, Journal of Materials Science l6, 1437 (1981). 

11. Y. W. Kim, W. M. Griffith, and F. H. Froes, Surface Oxides in P/M Aluminum Alloys, Journal of metals 37, 27 (1985). 

12. J. S. Benjamin and R. D. Schelleng, Dispersion Strengthened Aluminum-4 

Pct Magnesium Alloy Made by Mechanical Alloying, Metallurgical Transactions x. 1827 (1980) I 

13. S. J. Donachie, High Strength, P/M Aluminum Mill Products, Phase I. Low 

Cost Bi 1 let Manufacturing, Report No. AFWAL-TR-81-4090 on Air Force Contract No. F33615-79-C-5160 (November 1981 ). 

63 
14. E. A. Starke, Jr., T. H. Sanders, Jr., and I. G. Palmer, New Approaches 

to Alloy development in the A1-Li System, Journal of Metals 33, 24 (1981). 

15. E, Klar and J. W. Fesko, Gas and Water Atomization, in Metals Handbook 

JNinth Editionl, Volume 7 - Powder Metallurgy (American Society for 

Metals, Metals Park, OH, 1984), p. 25. 

16. H. Matyja, 8. C. Giessen, and N. J. Grant, Effect of Cooling Rate on the 

Dendrite Spacing in Splat Cooled Aluminum Alloys, Journal of the Institute of Metals 96, 30 (1968). 

17. P. 3. Meschter, R. J. Lederich, J. E. O'Neal, and P. S. Pao, Study on 

Effects of Flake Chemistry and Morphology on the Properties of 

Al-Cu-Hg-X-X-X Powder Metallurgy Advanced Aluminum Alloys, NASA Contractor Report 177946 on Contract NASl-17107 (November 1985). 

18. H. K. Hardy and J. M. Silcock, The Phase Sections at 500°C and 350°C of 

Aluminum-Rich Aluminum-Lithium-Copper Alloys, Journal of the Institute of 

Metals 84, 423 (1955-56). 

19. R. F. Singer, W. C. Oliver, and W. D. Nix, Identification of Dispersoid 

Phases Created in Aluminum During Mechanical Alloying, Metallurgical 

Transactions m, 1895 (1980). 

20. J. M. Silcock, The Structural Aging Characteristics of AluminumCopper-Lithium Alloys, Journal of the Institute of Metals 88, 357 

(1 959-60). 

21. B. Noble, I. R. McLaughlin, and 6. Thompson, Solute Atom Clustering 

Processes in Aluminum-Copper-Lithium Alloys, Acta Metallurgica l8, 339 

(1 970). 

22. 8. M. Watts, M. J. Stowell, B. L. Baikie, and D. 6. E. Owen, Superplasticity in Al-Cu-Zr Alloys, Part I: Material Preparation and Properties (p. 189) and Part 11: Microstructural Study (p. 198), Metal Science 

- 10 (1976). 

23. 3. D. Embury, Strengthening by Dislocation Substructures, in 

Strengthening Mechanisms in Crystals,. A. Kelley and R. B. Nicholson, eds. 

(Applied Science Publishers, London, 1971), p. 342. 

24. Rapid Inexpensive Tests for Determining Fracture Toughness, Report of the 

Committee on Rapid Inexpensive Tests for Determining Fracture Toughness, 

National Materials Advisory Board (National Academy of Sciences, 

Washington, D.C., 1976), p. 62. 

25. J. W. Bohlen and S. W. Ping, Development of High Strength Fatigue Crack 

Growth Resistant Aluminum Alloys By Rapid Solidification Technology, 

Final Technical Report on Navy Contract N60921-82-C-0043 (1983). 

64 
8. APPENDIX 

Under authority granted by the NASA Contracting Officer to furnish samples of alloys prepared under the present contract for additional evaluation, 

such samples were used by the contractor for an ongoing independent analytical 

investigation of powder-processed aluminum alloys. This investigation included powder surface characterization and transmission electron microscopy. The 

latter study was conducted by Mr. Robert Wheeler, graduate student, and 

Professor Hamish L. Fraser of the Department of Mining and Metallurgical 

Engineering, University of Illinois - Champaign. Highlights of the results of 

these studies are summarized in the following paragraphs for general 

information. 

Thermogravimetric analysis (TGA) of preal loyed A1-4Cu-lMg-l.5Li-0.2Zr 

powder revealed a continuous loss of weight with increasing temperature up to 

623K (662°F). followed by an increase of weight above this temperature. Mass 

spectrometric analysis of the constituents that are liberated on heating the 

powder revealed the presence of only water. X-ray photoelectron spectroscopy 

of this powder also revealed the presence of only oxide on the surface. 

Because of the very weak signals from the powder surface, infrared spectroscopy could not be used successfully for characterizing the surface. The 

above results suggest the absence of chemically bound water on the powder surface thereby obviating the need for vacuum degassing above the highest postconsolidation processing temperature. 

The weight loss between 323K (122OF) and 623K (662°F) for the A1-4CulMg-1.5Li-0.2Zr powder was 0.76 pct. (Figure Al), between 328K (131°F) and 

708K (815°F) for the A1-4Cu-lMg-1.5Fe-0.75Ce powder was 0.27 pct. (Figure A2), 

and between 328K (131°F) and 828K (1031°F) for the Al-1 .6Fe-0.8Ce powder was 

0.12 pct (Figure A3). The absolute weight loss values for the Fe- and 

Ce-containing alloys were about three orders of magnitude lower than that 

reported for air-atomized 7091 alloy. 

The fine sub-grain structure resulting from the refinement due to the 

AlgZr dispersoids in the alloys made from prealloyed A1-4Cu-lMg-1.5Li-0.2Zr 

powder are shown for alloy 3 in Figure A4. The fine sub-grains and the 

Fe- and Ce-containing dispersoids in the alloys made from prealloyed A1-4CulMg-1.5Fe-0.75Ce powder are shown for alloy 6 in Figure A5. The transmission 

electron micrograph of alloy 9 (Figure A6) reveals the very fine recrystallized grains typical of mechanically alloyed materials. The transmission 

electron micrograph of the hot pressed billet of alloy 15 (Figure A7) shows 

the prior-particle surface oxide film in a bright field and dark field pair, 

and that of the extrusion of the same alloy (Figure A8) shows the development 

of sub-grains following hot working. 

65 
99.2 

300 400 500 600 700 800 900 

TEMPERATURE - K 

M11-7251 

99.6 - 

- 

RUN 2 

99.4 - 

- 

300 400 500 600 700 800 900 

TEMPERATURE - K 

FIGURE Al. THERMOGRAVIMETRIC ANALYSIS SCANS FROM 

PREALLOYED AI - 4 Cu - 1Mg - 1 SLi - 0.2 Zr POWDER 

66 
100.05 

100.00 

99.95 

99.90 

99.85 

99.80 

99.75 

99.70 

300 350 400 450 500 550 600 650 700 750 800 850 

TEMPERATURE - K 

M11-7252 

FIGURE A2. THERMOGRAVIMETRIC ANALYSIS SCAN FROM 

PREALLOYED AI - 4Cu - 1Mg - 1.5Fe - 0.75 Ce POWDER 

300 350 400 450 500 550 600 650 700 750 800 850 

TEMPERATURE - K 

Ml l -7253 

FIGURE A3. THERMOGRAVIMETRIC ANALYSIS SCAN FROM 

PREALLOYED AI - 1.6 Fe - 0.8 Ce POWDER 

67 
11 -6480 

I 0.5 pm I 

a 

I 

I 'Ccm I 

FIGURE A4. TRANSMISSION ELECTRON MICROGRAPH OF ALLOY 3 

IN THE ARTIFICIALLY AGED CONDITION 
I 

ORIGINAL PAGE 

of POOR QuAlsrr 

11 -6481 

FIGURE A5. TRANSMISSION ELECTRON MICROGRAPH OF ALLOY 6 

IN THE EXTRUDED CONDITION 

11 -6482 

0.5prn 

FIGURE A6. TRANSMISSION ELECTRON MICROGRAPH OF ALLOY 9 

IN THE SOLUTION HEAT TREATED CONDITION 

69 
FIGURE A7. TRANSMISSION ELECTRON MICROGRAPH OF ALLOY 15 

FOLLOWING CONTAINERLESS VACUUM HOT PRESSING. 

(a) BRIGHT FIELD AND (b) DARK FIELD 
ORlGtNkL PAGE IS 

OF POOR QUAlfTY 

11 -6484 

’ lJm 

FIGURE A8. TRANSMISSION ELECTRON MICROGRAPH OF 

ALLOY 15 IN THE EXTRUDED CONDITION 

71 
' . Report No. 

NASA CR-178227 

. Title and Subtitle 

CONSOLIDATION PROCESSING PARAMETERS AND ALTERNATIVE 

PROCESSING METHODS FOR POWDER METALLURGY A1-Cu-Mg-X-X 

2. Government Accession No. 

ALLOYS 

20 rity Clayif. of this page) of Pages I iJn3assi fled 

19. Secur'ty Classif.(of this report) Unclassified 

~~ 

I- * --,- . Author(s) Lb-4.' , 

K. K. Sankaran 

22. Price 

. Performing Organization Name and Addrean 

McDonnell Douglas Astronautics Company-St.Louis Divisior 

P. 0. Box 516 

St. Louis, Missouri 63166 

2. Sponsoring Agency Name and Address 

National Aeronautics and Space Admi ni strati on 

Washington, DC 20546 

5. Supplementary Notes 

Langley Technical Monitor: 

Final Report 

Dennis L. Dicus 

3. Recipient's Catalog No. 

5. Report Date 

February 1987 

6. Performing Organization Code 

8?Performing Organization Report No. 

~ 

10. Work Unit No. 

11. Contract or Grant No. 

NASl-16967 

13. Type of Report and Period Covered 

Contractor Report 

14. Sponsoring Agency Code 

505-63-01 -02 

6. Abstract 

A study was conducted to determine the effects of varying the vacuum 

degassing parameters on the microstructure and properties of A1-4Cu-1Mg-X-X 

(X-X = 1.5Li-0.2Zr or 1.5Fe-0.75Ce) alloys processed from either prealloyed 

(PA) or mechanically alloyed (MA) powder, and consolidated by either using 

sealed aluminum containers or containerless vacuum hot pressing. The consolidated billets were hot extruded for evaluating the microstructure and properties. The MA Li-containing alloy did not include Zr, and the MA Fe- and Cecontaining alloy was made from both elemental and partially prealloyed 

powder. The alloys were vacuum degassed both above ana below the solution 

heat treatment temperature. While vacuum degassing lowered the hydrogen content of these alloys, the range over which the vacuum degassing parameters 

were varied was not large enough to cause significant changes in degassing 

efficiency, and the observed variations in the mechanical properties of the 

heat treated alloys were attributed to varying contributions to strengthening 

by the sub-structure and the dispersoids. Mechanical alloying increased the 

strength over that of alloys of similar composition made from PA powder. The 

inferior properties in the transverse orientation, especially in the Licontaining alloys, suggested deficiencies in degassing. Among all of the 

alloys processed for this study, the Fe- and Ce-containing alloys made from MA 

powder possessed better combinations of strength and toughness. 

L7. Key Wordn (Suggested by Authom(8)) 

Powder- Processed A1 umi num A1 1 oys 

Rapid Solidification 

Mechanical Alloying 

Powder Vacuum Degassing 

A1 -Cu-Mg A1 lOyS 

18. Distribution Statement 

Unclassified - Unlimited 

Subject Category 26 

For sale by the National Technical Information Service, Springfield, Virginia 22161 

NASA Langley Form 63 (Junm 1085)