NASA Contractor Report 177946 

NASA-CR-177946 

19860006863 

Study on Effects of Powder and Flake 

Chemistry and Morphology on the Properties of 

AI-Cu-Mg-X-X-X Powder Metallurgy Advanced 

Aluminum Alloys 

P. J. Meschter, R. J. Lederich, J. E. O'Neal, and P. S. Pao 

McDonnell Douglas Research Laboratories 

St. Louis, Missouri 

Contract NASl-17107 

November 1985 

NI\SI\ 

National Aeronautics and 

Space Administration 

Langley Research Center 

Hampton, Virginia 23665 

FOR P1EFERENCE 

,.01 .0 6E TAKEN FROM nils ~oow 

f \ 

, ~ , 

J It I w • \ .. 1 r,' :ol,~! '\ 

111111111111111111111111111111111111111111111 

NF00705 
NASA Contractor Report 177946 

Study on Effects of Powder and Flake 

Chemistry and Morphology on the Properties of 

AI-Cu-Mg-X-X-X Powder Metallurgy Advanced 

Aluminum Alloys 

P. J. Meschter, R. J. Lederich, J. E. O'Neal, and P. S. Pao 

McDonnell Douglas Research Laboratories 

St. Louis, Missouri 

Contract NASI-17107 

November 1985 

NI\S/\ 

NatIonal AeronautIcs and 

Space AdmInIstratIon 

Langley Research Center 

Hampton, VIrgInIa 23665 
PREFACE 

ThlS report presents the results of research performed from 21 September 

1982 through 20 June 1984 by the McDonnell Douglas Research Laboratories under 

the National Aeronautics and Space Administration Contract NASI-17107 entltled 

"Study of Effects of Powder and Flake Chemistry and Morphology on the Propertles of Al-Cu-Mg-X-X-X Powder ~etallurgy Advanced Aluminum Alloys." The 

obJectlves were to (1) investlgate systematlcally the effects of alloy chemlstry and control, and particulate morphology on the resultant consolidated 

product properties for the Al-Cu-Mg-X-X-X alloy family, and (2) to identify 

the principal contributing chemical/metallurgical/processing factors which 

affect each property of interest. The results provlde a better fundamental 

understanding of the processing-structure-property relationships for this 

alloy class and a start for defining the degree of process control necessary 

to accomplish commercialization. 

The research was performed ln the Solid State Sciences Department, 

managed by Dr. C. R. Whitsett. The principal investigator was 

Dr. P. J. Meschterj co-investigators were Dr. P. S. Pao, Mr. J. E. O'Neal, and 

Mr. R. J. Lederich. The technical manager was Mr. W. Barry Lisagor, NASA, 

Langley Research Center. 
This Page Intentionally Left Blank 
1. 

2. 

3. 

4. 

5. 

6. 

CONTENTS 

Page 

INTRODUCTION •••••••••••••••••••••••••••••••••••••••••••••••••••••• 1 

RESEARCH OBJECTIVES AND APPROACH •••••••••••••••••••••••••••••••••• 2 

2.1 

2.2 

ObJect1ves ••••••••••••••••••••••••••••••••••••••••••••••••••• 

Selection of Alloy Compos1t1ons •••••••••••••••••••••••••••••• 

2 

3 

EXPERIMENTAL PROCEDURE •••••••••••••••••••••••••••••••••••••••••••• 5 

3.1 Part1culate Product10n ••••••••••••••••••••••••••••••••••••••• 5 

3.2 Part1culate Character1zation ••••••••••••••••••••••••••••••••• 9 

3.3 Consolidat10n and Post-Consol1dation Processing •••••••••••••• 9 

3.4 Microstructures and properties of Consolidated Forms ••••••••• 11 

3.4.1 M1crostructure, Texture, and Dens1ty •••••••••••••••••• 11 

3.4.2 Mechanical Properties ••••••••••••••••••••••••••••••••• 11 

TASK 1: PRODUCTION AND CHARACTERIZATION OF RAPIDLY SOLIDIFIED 

POWDERS AND FLAKES................................................ 13 

4.1 Evaluation of Dispersoid-Forming Systems for Non-Li4.2 

4.3 

Containing Alloy ............................................ . 

Product10n of Rapidly Solid1f1ed Particulates •••••••••••••••• 

4.2.1 Vacuum Atomization ..••.••.•....•...•••.....•••........ 

4.2.2 Ultrasonic Atomization •••••••••••••••••••••••••••••••• 

4.2.3 Twin Roller Quenching .•........•••.•.......•.•......•. 

Part1culate Characterization ••••••••••••••••••••••••••••••••• 

4.3.1 

4.3.2 

4.3.3 

4.3.4 

Chemi cal Anal y s e s ....••.••...••..•.........•.......•.. 

Particle Size Distr1bution •••••••••••••••••••••••••••• 

Metallographic Powder Characterization •••••••••••••••• 

SEM Characterizat10n and Dendrite-Arm-Spac1ng 

13 

15 

15 

15 

16 

17 

17 

19 

21 

Measurements. • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • •• 24 

4.3.5 Oxide Layer Thickness Characterizat10n •••••••••••••••• 29 

TASK 2: CONSOLIDATION AND POST-CONSOLIDATION PROCESSING OF 

PARTICULATES. • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • . • • •• 35 

5.1 particulate Degassing Exper1ments •••••••••••••••••••••••••••• 35 

5.1.1 Removal of Hydrogen and Oxygen •••••••••••••••••••••••• 35 

5.1.2 OX1de Layer Mod1fication During Degassing ••••••••••••• 37 

5.2 Characterizat10n of Extrusions and Forgings •••••••••••••••••• 37 

5.2.1 Chemical Analyses..................................... 37 

5.2.2 Microstructure ........................................ 39 

5.2.3 Texture ..•.....•...•...........•...........•.......... 43 

5.3 Heat Treatment of Extrus10ns and Forgings •••••••••••••••••••• 47 

TASK 3: MICROSTRUCTURE AND PROPERTIES OF CONSOLIDATED FORMS •••••• 50 

6.1 

6.2 

Microstructure .•..•.••..........•............•.••............ 

6.1.1 

6.1.2 

Solution-Treated Condition •••••••••••••••••••••••••••• 

Aged Cond1tion •...••..•.....•.....•.....••............ 

Properties ..•....•........•.•.•.•..........•................. 

6.2.1 

6.2.2 

De ns 1 t y •••••••..••••••••••••••••..••••••••••.•...•.••• 

Elastic Modulus .................•..................... 

111 

50 

50 

56 

62 

62 

63 
7. TASK 

7.1 

7.2 

7.3 

7.4 

7.5 

6.2.3 

6.2.4 

6.2.5 

6.2.6 

6.2.7 

CONTENTS (Continued) 

Tensile Mechanical Properties ••••••••••••••••••••••••• 

6.2.3.1 Non-Li-Conta1ning Alloys ••••••••••••••••••••• 

6.2.3.2 

Fracture 

Li-Conta1n1ng Alloy ......................... . 

Toughness ................................... . 

Fat 19ue-Crack-Growth Rate ••••••••••••••••••••••••••••• 

Creep Res 18 tanee •••••••••••••••••••••••••••••••••••••• 

Stress-Corros10n-Cracking Res1stance •••••••••••••••••• 

4: CORRELATION OF PROCESS/STRUCTURE/PROPERTIES RELATIONS •••• 

Effect of Particulate Type on Propert1es ••••••••••••••••••••• 

Effect of Processing Var1ables on Propert1es ••••••••••••••••• 

7.2.1 processing Sequence .•..........•.•.........•.......... 

7.2.2 Texture Effects ...........•........•.................. 

Effect of Alloy Chem1stry and Microstructure on Properties ••• 

7.3.1 Non-L1-Contain1ng Alloys •••••••••••••••••••••••••••••• 

7.3.2 Li-Conta1ning Alloy .....•....•...••..•.••...•..•....•. 

Effect of Heat Treatment Var1ables on Properties ••••••••••••• 

Recommendations for Optimal Composition/Processing/ 

Structure Comb1nat10ns 1n RSP Alum1num Alloys •••••••••••••••• 

Page 

64 

64 

70 

76 

77 

80 

8t 

83 

83 

84 

84 

87 

87 

87 

89 

90 

90 

8. CONCLUSIONS ••••••••••••••••••••••••••••••••••••••••••••••••••••••• 92 

9. REFERENCES. • • • • • • • • • • • . • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • . • • • . • • . .. 95 

IV 
LIST OF ILLUSTRATIONS 

F1gure Page 

l. Homogeneous Metals, Inc. vacuum atom1zat10n apparatus •••••••••••• 6 

2. MDRL ultrason1c atom1zat10n apparatus ••••••••••••••••••••••••.••. 7 

1. MDRL roller-quenching apparatus •••••••••••••••••••••••••••••••••• 8 

4. Optical micrograph of AI-1Fe-1Ni flake ••••••••••••••••••••••••••• 13 

5. TEM micrographs of roller-quenched Al-1Fe-1Ni flakes ••••••••••••• 14 

6. Particle size distributions of AI-4.4Cu-1.5Mg-1Fe-1Ni 

powders •••••••••••••••••••••••••••••••••••••••••••••••••••••••••• 19 

7. Particle size distributions of Al-3L1-1.5Cu-1.0Mg-0.5Co-0.2Zr 

powders ••••••.•.•••.•••••••.•••••••••.•••.••.•.•.•••.••.•••••.••• 20 

8. Opt1cal micrographs of Al-4.4Cu-1.5Mg-1Fe-1Ni-0.2Zr powders •••••• 22 

9. Optical micrographs of Al-3Li-1.5Cu-1Mg-O.5Co-0.2Zr powders •••••• 23 

10. Scanning Electron Micrographs of Al-4.4Cu-1.5Mg-1Fe-1Ni-O.2Zr 

particles. . . . . . . . . • . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . • . . . . . . . .. 24 

11. Dendrite arm spacing as a function of particle d1ameter 

for Al-4.4Cu-1.5Mg-1Fe-1Ni-O.2Zr vacuum atomized powder •••••••••• 25 

12. Dendrite arm spacing as a function of particle diameter for 

Al-4.4Cu-1.5Mg-IFe-1Ni-O.2Zr ultrasonically atomized in argon •••• 26 

13. Dendrite arm spacing as a function of particle diameter for 

Al-4.4Cu-1.5Mg-1Fe-1Ni-O.2Zr ultrasonically atomized in helium ••• 26 

14. Dendrite arm spacing as a function of part1cle d1ameter for 

Al-3Li-1.5Cu-1Mg-O.5Co-O.2Zr vacuum atom1zed powder •••••••••••••• 27 

15. Dendr1te arm spacing as a function of part1cle diameter for 

Al-3L1-1.5Cu-1Mg-O.5Co-0.2Zr ultrason1cally atomized in argon •••• 27 

16. Dendr1te arm spacing as a function of particle diameter for 

Al-3Li-1.5Cu-1Mg-0.5Co-0.2Zr ultrasonically atomized 1n heI1um ••• 28 

17. Scann1ng Auger m1croscope (SAM) results for as-vacuum-atom1zed 

Al-4.4Cu-1.5Mg-1Fe-1Ni-0.2Zr part1cle •••••••••••••••••••••••••••• 30 

18. SAM results for as-vacuum-atomized Al-3L1-1.5Cu-1Mg-0.5Co-0.2Zr 

part1cle •••••••••••••••••••••••••••••••• ~ ••••••••••••• •••••••.••• 31 

19. Depth profiles of oxide f11ms on helium atomized particles ••••••• 13 

20. Variat10n w1th degass1ng temperature of hydrogen concentrat10n 

1n vacuum-atom1zed alloy powders ••••••••••••••••••••••••••••••••• 36 

21. Variation with degassing temperature of oxygen concentration 

in vacuum-atomized alloy powders ••••••••••••••••••••••••••••••••• 36 

22. Depth profiles of oxide films on vacuum-atom1zed particles, 

degassed 530 oc/2h in vacuum •.•••••.•••.•••.•••••••.••••....•..... 38 

23. Optical micrographs of non-Li-containing alloy extrusions •••••••• 40 

v 
24. 

25. 

26. 

27. 

28. 

Optical m1.crographs 

Opt1.cal m1.crographs 

( 111) pole f1.gures 

(111) pole f1.gures 

(111) pole f1.gures 

of Al-3L1.-l.5Cu-lMg-0.5Co-0.2Zr extrUS1.ong •.• 41 

of forg1ngs ••••.•.•••••••..•••..•••••••••.... 42 

of VA ext rus lons ••••••••.•••••••.••••.••..•..• 44 

of Al-3L1.-1.5Cu-lMg-O.5Co-O.2Zr extrus~ons •••• 45 

of f orgl.ngs ••••••••••••••••••••••••••••••••••• 46 

29. Ag~ng behav~or at 1900 C of non-L~-conta1.n~ng VA extrus~ons ••...• 48 

30. Ag1.ng behav~or at 1770c of Al-3L~-1.5Cu-1Mg-O.5Co-0.2Zr VA 

ext rUS10ns • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • •• 49 

31. Ag1.ng behav1.or at 177°C of Al-3L1.-1.5Cu-1Mg-0.5Co-O.2Zr 

USGA-He extrusl0ns •••••••••••••••.••••••.•••.••..•••••..••••••••• 49 

32. Optical m1.crographs of solution-treated extrUS1.ons •••••••••••••• 51 

33. Optical micrographs of Al-4.4Cu-l.5Mg-lFe-lNi-0.2Zr extrUS1.ons 

from powders ultrasonically atom1.zed in argon and helium ••••••••• 52 

34. Optical micrographs of forgings solut1.on-annealed 530oC/l h •••••• 53 

35. Optical micrographs of Al-3Li-l.5Cu-lMg-0.5Co-0.2Zr 

extrusions, solution-treated 530oC/1 h ••••••••••••••••••••••••••• 54 

36. Opt1.cal micrographs of Al-3L1.-1.5Cu-lMg-0.5Co-0.2Zr 

extrusions, solution-treated 560oC/1 h ••••••••••••••••••••••••••• 55 

37. TEM micrographs of non-Li-containing extrusions, T6 temper ••••••• 57 

38. TEM micrographs of Al-4.4Cu-l.5Mg-lFe-lNi-0.2Zr extrusions 

and forgings, T351 and T85l tempers •••••••••••••••••••••••••••••• 58 

39. TEM micrographs of Al-4.4Cu-l.5Mg-IFe-1Ni-0.2Zr VA extrUS10ns 

T8S1 temper •••••••••••••••••••••••••••••••••••••••••••••••••••••• S9 

40. TEM micrographs of aged Al-3Li-l.5Cu-1Mg-0.5Co-0.2Zr 

extrusions ••••••••••••••••••••••••••••••••••••••••••••••••••••••• 61 

41. Ambient temperature tens1le properties of non-L1-containing alloys, T6 temper................................................ 64 

42. Amb1ent temperature tens1le propert~es of non-L1-conta~ning 

alloy ext rusions, T851 temper.................................... 65 

43. Amb1ent temperature tens~le propert~es of non-L~-conta~n~ng 

alloy extrus10ns and forgings, T8S1 temper ••••••••••••••••••••••• 65 

44. Amb1.ent temperature tens1.le propert1.es of non-L1.-conta1.n1.ng 

alloy VA extrus10ns and forg1.ngs, T351 temper ••••••••••••••••.••• 66 

45. Tens1le propert1.es of non-L~-conta1ning alloy extrus~ons 

and forglngs, T851 temper •••••••••••••••••••••••••••••••••••••••• 66 

46. Var1.ation W1.th temperature of the y~eld stress after 100 h 

exposure. • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • . . • • • • . • • . . • . . • • • • •• 67 

47. Var1.at10n W1.th temperature of the ult1.mate tenslle stress 

after 100 h exposure •••••••••••.•••••••••••••••••••••••••••••.••• 67 

48. Variation w1.th temperature of the duct~l~ty after 100 h 

exposure. • • • • • • • • . • • • . • • • . • • . . • • • • • • • • . • • • . • • • . • • • • . . . • • • . • • . . . .. 68 

VI 
49. SEM micrographs of Al-4.4Cu-l.5Mg-lFe-lNi-O.2Zr VA 

extrusion fracture surfaces •••••••••••••••.•••••••••••••••••••.•• 71 

50. Ambient temperature tens1le propert1es of Al-3Li-l.5Cu-lMgO.SCo-O.2Zr extrus10ns and forgings, T6/S30 temper ••••••••••••••• 72 

51. Long1tud1nal and transverse ambient temperature properties of 

Al-3L1-1.5Cu-lMg-O.5Co-0.2Zr extrus10ns, T6/530 temper ••••••••••• 73 

52. Ambient temperature tens1le properties of Al-3L1-1.5Cu-lMg0.5Co-0.2Zr extrusions, T6/560 temper •••••••••••••••••••••••••••• 73 

53. Amb1ent temperature tensile properties of Al-3L1-l.SCu-lMg0.5Co-0.2Zr extrusions and forgings, T8S1 temper ••••••••••••••••• 74 

54. Tens1le properties at 1nd1cated temperatures of Al-3Li-l.5Cu1Mg-0.5Co-0.2Zr extrusions and forg1ngs •••••••••••••••••••••••••• 75 

55. Fractographs of Al-3Li-1.5Cu-lMg-0.5Co-0.2Zr extrusions •••••••••• 76 

56. Fatigue-crack-growth-rate (FCGR) results on Al-4.4Cu-l.5MglFe-1N1-0.2Zr VA extrusions, T8Sl temper ••••••••••••••••••••••••• 78 

57. FCGR results on Al-4.4Cu-l.SMg-IFe-lNi-O.2Zr VA forgings, 

T8S1 temper...................................................... 78 

58. FCGR results on Al-3Li-l.5Cu-1Mg-0.5Co-0.2Zr VA extrus1ons, 

r851/530 temper .•.•.....••....•...••.•••...•.•....•.•....•....... 79 

59. FCGR results on Al-3Li-l.5Cu-1Mg-0.5Co-0.2Zr VA forgings, 

T851/530 temper ••••••.••••••••••••••••••••.•••••••••••...•...•••• 79 

60. IS00 C creep-test results for T851 temper Al-4.4Cu-l.5MgIFe-lNi-O.2Zr •••....•...•.•.....•.•...........•••................ 80 

61. Time-to-failure in alternate 1mmerS1on C-r1ng tests of 

Al-3L1-1.5Cu-lMg-0.5Co-O.2Zr VA extrusions, T851/530 temper, 

T orientat10n.................................................... 81 

62. Time-to-fa1lure in alternate immersion C-ring tests of 

Al-3Li-l.5Cu-lMg-0.5Co-0.2Zr VA forg1ngs, T851-530 temper, 

S orientat10n.................................................... 82 

ViI 
LIST OF TABLES 

Table Page 

1. Operatlng Parameters for UltrasonlC Atomlzatlon of Al Alloy 

powders. . • • . . • • . • • . . . • . . . . . • . . . . . . . . . . . • . . . • . . . . . • . . . . . . . . . . . • . .. 16 

2. Chemlcal Analyses of Non-Li-Conta1ning Alloy Partlculates •.•••••. 17 

3. Chemlcal Analyses of Ll-Contaln1ng Alloy Partlculates •••••••••.•• 17 

4. Hydrogen and Oxygen Analyses of Partlculates and Extrusl0ns •••••• 18 

5. Mean Partlcle Dlameters, Dendr1te-Arm Spaclngs, and 

Solidlflcatlon Rates for Al-4.4Cu-l.5Mg-IFe-lNl-0.2Zr 

Part l.cles. . . . . . . . . . . . . . • . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .. 20 

6. Mean Particle Diameters, Dendrite-Arm Spaclngs, and 

SOlldlflcatl0n Rates for Al-3Ll-1.5Cu-lMg-O.5Co-O.2Zr 

Part1cles ........................................................ 21 

7. Densltles of RSP Alumlnum Alloy Extrusions ••••••••••••••••••••••. 63 

8. Elastic Moduli of RSP Alloy Extrusions •••••••••••••••.••••••••••• 63 

9. Fracture Toughnesses of Extruslons and Forglngs •••••••••••••.•••• 77 

Vlll 
1. INTRODUCTION 

Recent ~esea~ch and development of wrought aluminum alloys through rap1d 

sol1d1f1cat10n and powde~-metallurgical (P/M) process1ng has shown that s1gn1-

ficant imp~ovementg in properties, relative to those of conventionally formed 

alloys, can be expected, particularly 1n alloys w1th novel compos1t10ns. 

Rap1d sol1dification processing (RSP) of aluminum alloys has thus far been 

concerned with the development and screening of cand1date alloy compos1t10ns. 

The effects of powder and flake chemistry, morphology, and metallurg1cal cond1tion on the properties of consolidated alloys have not been adequately 

studied. Particulate cooling rate and shape, oX1de film formation and breakup, phase relat10nships, and recrystallizat10n and texturing all potent1ally 

affect final properties. These parameters must be systematically investigated 

to identify the extent of process control necessary to achieve commerc1al1zation and full exploitation of this promising new class of alum1num alloys. 
2. RESEARCH OBJECTIVES AND APPROACH 

This study is a systematic 1nvest1gat10n of the effects of alloy 

chemistry and particulate morphology on consolidat1on behav10r and consolidated product propert1es for the Al-Cu-Mg-X-X-X alloy fam1ly. The overall 

goal 1S to develop alloy chemistry design methods, part1culate production 

methods, consol1dation methods, and heat-treatment sequences so that rap1d 

solid1ficat10n processed (RSP) aluminum alloys 1n th1s class m1ght real1ze 

their potential for room-temperature and elevated-temperature app11cat10ns. 

2.1 Objectives 

1. Systemat1cally invest1gate the effects of alloy chemistry 1n non-L1-

and Li-containing Al-Cu-Mg-X-X-X alloys on the properties of heat-treated 

extrusions and forgings. Emphas1s was placed on Al-3L1-1.SCu-IMg-O.SCoO.2Zr for 1mproved ambient-temperature propert1es, and on disperso1dform1ng element additions to Al-4.4Cu-I.SMg for 1mproved properties to 

IS00C (3000F). 

2. Systematically invest1gate the effects of part1culate morphology and 

cooling rate on consolidated product m1crostructures and properties of AlCU-Mg-X-X-X alloys. The emphasis was on property d1fferences between 

powder extrusions and forgings, between extruS10ns consol1dated from 

powders and from flakes, and among extrus10ns consol1dated from 

particulates solidified at different rates. 

3. Ident1fy contributing chem1cal, metallurgical, and process1ng factors 

that affect various properties of interest, e.g., the degree of 1rnprovement 

in IS00C propert1es of a d1spersoid-conta1ning RSP alloy compared w1th 

commercial 2024-Al, or the effect of Li and Co add1t10ns on the balance 

between room-temperature strength and duct1lity 1n an Al-Cu-Mg alloy. 

4. Develop cr1teria for the des1gn of chem1stry/particulate-morphology/ 

consolidat10n/heat-treatment combinat1ons to enable powder-processed 

aluminum alloys to compete with commerc1al Al alloys 1n room-temperature 

applicat10ns and w1th titanium alloys 1n ISOoC appl1cat10ns. 

2 
The contract was dlvlded lnto four tasks, as follows: 

Task 1: Preparation and Characterlzatlon of Rapidly Solidlfied Al-Cu-Mg-XX-X Powder and Flakes; 

Task 2: Consolldation and Post-Consolidation Processing of Particulates; 

Task 3: Microstructure and Properties of Consolidated Forms; and 

Task 4: Correlation of Process/Structure/Properties Relations. 

2.2 Selection of Alloy Compositions 

Two alloy types were chosen: an Al-Cu-Mg alloy modlfled by Fe and Ni 

additions for application to IS00C, and an Al-Li-Cu-Mg alloy intended primarlly for ambient-temperature applications. An applications study (1) has shown 

that an advanced aluminum alloy must have a tensile strength of 480 MPa 

(70 ksi) at IS00 C to match the specific strength of Ti-6Al-4V. If aluminum 

alloys can reach strength equivalency with titanium alloys at this temperature, their lower cost will encourage substitution in structural applicatlons. 

The key to iSOoC strength improvement in aluminum alloys lies in the introduction of significant volume fractions of thermally stable, incoherent dis persoids. Since RSP imposes microstructural refinement, wrought parts made from 

RSP powders can have finer dispersoids than equivalent ingot-metallurgical 

(l/M) alloy parts. Fine, closely spaced dispersoids interfere with dislocation motion and contribute to elevated-temperature strength without degradation of other properties (2). In this study, dlspersoid-forming element addltions were made to the 2024-Al base system, Al-4.4Cu-l.SMg. The addition of 

0.2i. Zr was expected to further improve properties by stabilization of a fine 

subgrain structure, as for the Li-containing alloy discussed below. 

Two candidate dispersoid-forming systems, each calculated to produce 4-5 

voli. dispersoids, were investigated: IFe-lNl and 4Ti-i.8B. The former system 

produces Al9FeNi dlspersoids when Fe and Ni are added in equal quantitles (2). 

Durand et ale (3) demonstrated 100-120 MFa (14-17 ksi) improvements in yield 

and tensile strengths, with no loss in ductility, in 707S-Al + (IFe-lNi) extrusions made from roller-quenched flakes, compared to l/M 707S-Al. Thermodynamic calculations predict that the stable dispersoid phase in the Al-4Ti-l.8B 

system is TiB2• Slaughter and Das (4) showed that additlons of Ti and B to 

RSP iron-aluminum alloys produced fine T1B2 dispersoids, which refined the 

grain size and were expected to significantly lmprove mechanical propertieg. 

3 
Alum1num-lith1um alloys have lower dens1t1es and h1gher elast1c modul1 

than commercial 2XXX and 7XXX alloys at attract1ve strength levels, enabl1ng 

significant we1ght savings 1n a1rcraft structural members (5,6). Convent10nal 

11M AI-L1 alloys are 11mited to ( 2.6 wt% L1 by L1 segregat10n to coarse 

1ntermetallics 1n cast ingots. RSP Al-L1 alloys w1th ~ 3 wt% L1 show prom1se 

of be1ng consolidated and heat treated to reta1n m1crostructural homogene1ty 

and properties combinat10ns superior to those of 11M alloys. 

A prom1sing class of AI-L1 alloys conta1ns 3% L1, 1.5-2% Cu, and about 1% 

Mg (6). These alloys have 7-10% lower densit1es, 15-20% h1gher specif1c elastic modul1, and tensile propert1es comparable to those of 7075-AI, w1th 3-5% 

ductility. The (Li+Cu+Mg) level 1S sufficiently low that few 1nsoluble constituent part1cles are present after solution-treatment to act as fracture 

initiation sites and impair duct111ty. Thus a base compos1t10n of AI-3L1-

1.5Cu-1Mg appears prom1sing if further alloying add1t10ns can be made to 

improve ductility and toughness. 

Zr additions to AI-Li alloys are effective 1n increas1ng strength and 

ductility by preserving a dynam1cally recovered, unrecrystall1zed m1crostructure after solution treatment (7,8). Recrystal11zation and gra1n growth are 

inhibited by a homogeneous distribution of coherent, metastable Al3Zr d1spers01ds, which p1n subgrain boundaries and stabilize the substructure formed 

during prior hot working (extrus10n, forg1ng, rolling). 

Low ductility in Al-Li alloys has been attributed to planar slip, which 

causes localized stress concentrations at grain boundar1es and leads to 1ntergranular failure at low strains. Fine, incoherent dispersoids homogen1ze sl1p 

by forcing Orowan bypassing, thus reducing gra1n boundary stress concentrations. Dispers01d-forming elements such as Fe, Co, Cr, Mn, Cd, or Y produce 

significant increases in duct1lity when added to an RSP b1nary AI-L1 alloy 

(7,8). Cobalt 1S read1ly introduced to form a s1gn1ficant volume fract10n of 

1ncoherent d1spersoids, and was therefore selected for th1s study. 

Based on the above cons1derations, the compos1t10n Al-3L1-1.5Cu-1Mg0.SCo-0.2Zr was selected, to combine the low dens1ty and high spec1f1c modulus 

of an Al-3Li alloy with the sol1d-solution strengthen1ng of Mg, duct1l1ty 

enhancement by Co-conta1ning dispersoids, and resistance to recrystallizat10n 

assoc1ated with Al3Zr d1spers01ds. 

4 
3. EXPERIMENTAL PROCEDURE 

3.1 Particulate Production 

Aluminum alloy powders were produced by vacuum atomizatlon (VA) at 

Homogeneous Metals, Inc. (HMI) , Clayville, NY, uS1ng the apparatus shown 1n 

Figure 1. An 8-kg charge was induction melted 1n an argon-f111ed pressure 

chamber. Once the melt had been homogenized and heated to the des1red temperature, the sealing door was opened and the melt was ejected through a transfer 

tube 1nto the evacuated powder collection tank. The powder was SOlldif1ed by 

the escape of the pressurizing gas and cascaded into a collect1on vessel Wh1Ch 

could be valved off from the chamber for transportat1on and further handling. 

Typ1cal powder Y1elds ranged from 4.5 to 6.8 kg. Further information on 

vacuum atomlzation is available in the Ilterature (9,10). 

The MDRL ultrasonic gas atomization (USGA) apparatus, used to produce RSP 

AI-alloy powders with argon (USGA-Ar) or helium (USGA-He) as the atom1z1ng 

gas, is shown in Figure 2. A pre-alloyed charge weighing up to 6.8 kg was 

induction melted in an inert atmosphere. When the desired melt temperature 

was reached, atomization was initiated by simultaneous activation of soleno1ds 

which opened the gas line and an orifice ln the floor of the melting crucible. 

The molten metal was atomized and solidified by a pulsed (80-100 kHz), ultrasonic stream of inert gas emanating from an annular nozzle ring. The atomized 

powder was recovered from the main chamber collection container and from the 

cyclone through which the atomizlng gas exhausted. Details of USGA are 

avallable in the literature (9-13). 

The MDRL twin-roller quenching (RQ) appratus, used to produce RSP Alalloy flakes, 1S shown in Figure 3. A 900-g charge made from pure alum1num 

and master alloys was inductlon melted, homogenized by stirring, and heated to 

the desired temperature. A plunger sealing a 1.SS-mm-dlameter hole in the 

crucible floor was raised sufficiently to allow dropwise flow between two 

stainless-steel rollers rotating in opposite directlons at 1725 rpm. The 

molten drops were SOlld1fied lnto flakes typically 5 to 10 cm long, 0.5 cm 

wlde, and 125 pm thlCk. The flakes fell into a collection chamber, where they 

were retained under lnert gas cover. Typical yields were 400 to 600 g per 

run. The apparatus has been described in detail elsewhere (14). 

5 
Power 

supply 

Meltmg crucIble 

- To vacuum (' , 1_, r----- - pumps 

, I ( 

, ' , , , 

\\\\!/III 

\\ \ \ I ff Ii 

\\\\111/1 

\\\\!iJ/1 

Powder 

collectIOn tank 

}4---";:"r---+- Seahng door system 

"t::;:=:::::+=~ 

~----;'--i- Pressure vessel 

Powder dram 

To vacuum pumps • ___ ., Powder 

contamer 

GP41~21-IO 

Figure 1. Homogeneous Metals, Inc. vacuum atomIzation apparatus. 

6 
3000 

2000 

1000 

o 

Thermocouple leads ~ Control rod weIght 

Control rod solenoId 

CrucIble wIth charge 

Tundlsh 

Water coohng==:::;;::::~=~~~~~~~ -AtomIzIng 

Quartz sleeve Gas SOlenOld

1 gas 100 kVA 

Induction 

heatIng 

power 

Cyclone generator 

fIlter 

FInes 

collector 

IsolatIOn valves - __ 1"----,,1 

------'L--L...-, 

CollectIOn chamber '--__ ~ Mechamcal vacuum pump 

(.P41-0922 II 

Figure 2. MDRL ultrasomc atomization apparatus. 

7 
Stlrnng paddle 

Power supply ... _-- ===-....;;;;.;~~~ 

ObservatIOn 

o 

II 

, 

I 

I 

-:-:.=--

, 

~\ I , 

I .... 

I ' .... 

.. ---~---

I L. ______ _ 

........ 1 ~;::=::;:::=~=~ 

StiffIng handle 

, 

.... I 

J I I , 1 

-----~-i 

I 

-----~ ....... 

Crucible onflce 

o 

Scraper blades 

Vacuum 

chamber 

Flake 

collectIOn 

contaIner 

MechanIcal 

pump 

(.1'41 0922-12 

Figure 3. MDRL roller-quenchmg apparatus. 

8 
3.2 Particulate Characterization 

As-solid1fied particulates were chem1cally analyzed at McDonnell Douglas 

Corporation for Cu, Mg, Li, Fe, Ni, Co, Cr, Ti, Si, and Zn. Part1culates were 

analyzed for hydrogen and oxygen by a fus10n method at Leco, Inc., St. Joseph, 

MI. 

M1crostructures of alloy powders and flakes were examined using opt1cal 

microscopy and scanning electron m1croscopy (SEM). Optical m1croscopy was 

used for characterization of part1culate shapes, as-solidified m1crostructures, and additional features such as poros1ty or splat caps. Scanning 

electron micrographs of individual part1cles were used for determ1nation of 

dendr1te-arm spacings. 

Particle size d1stributions of powders were obtained by sieving 10-g 

samples according to ASTM E161-70, using electroformed sieves manufactured by 

the Buckbee-Mears Co., St. Paul, MN. Electroformed sieves have precise tolerances and are relatively free from clogging and particle entrapment; hence 

they produce particle size distribution data superior to those obta1ned uS1ng 

wire-woven sieves (15,16). Screens w1th apertures between 180 and 20 I.m were 

used. 

Oxide film thicknesses on as-solidified and as-degassed powders were 

determined using the Perkin-Elmer Model 600 Scanning Auger Multiprobe at 

Physical Electronics Laboratories, Eden Prairie, MN. This instrument 1S 

capable of obtaining Auger electron spectra, and hence chemical information, 

over areas as small as 50 nm d1ameter and 0.5-2.0 nm deep. It can also be 

used to obtain depth profiles of concentrations of up to 20 individual 

elements by alternate argon ion sputtering and Auger spectrum determination. 

After obtaining an initial Auger electron spectrum of a particle surface, the 

surface was ion milled until the oxygen Auger signal had decreased to approximately 15% of its 1nit1al intensity. The oX1de layer thickness was obtained 

from calibration curves determined on Ta20S• 

3.3 Consolidation and Post-Consolidation Processing 

Approximately 10-g samples of both non-Li- and Li-containing alloy 

powders were tray degassed 1n a tube furnace evacuated to 0.13 Pa (10-3 Torr), 

heated rap1dly to 400-5300 C, and held at temperature for 2 h. Hydrogen and 

9 
oxygen analyses of degassed part1culates were performed to determ1ne the 

effic1ency of degassing as a function of temperature. 

Samples of AI-4.4Cu-1.5Mg-lFe-1Ni-0.2Zr VA, USGA-Ar, and USGA-He; AI4.4Cu-1.5Mg-2Fe-2N1-0.2Zr VA; and Al-3L1-1.SCu-1Mg-0.SCo-0.2Zr VA, USGA-Ar, 

USGA-He, and RQ part1culates of 2.8 to 9.1-kg mass were consol1dated 1nto 

extrusions by Nuclear Metals, Inc., Concord, MA. The part1culates were 

compacted at ambient temperature into 14-cm-o.d. x 13.3-cm-1.d. x 2S.4-cm-Iong 

alum1num cans at 152 MPa (22 ksi). The cans were welded and evacuated at 

temperatures up to 500°C to < 0.1 Pa for degass1ng, then sealed. The billets 

were reheated to 400°C and extruded through a 1.27 x 3.81 cm (0.5 x 1.S-in.) 

rectangular d1e at 38 cm/m1n. The extrusion ratio was 29:1. Extrusion compos1tions were verif1ed by chemical analys1s, and hydrogen and oxygen analyses 

were performed to determine the effectiveness of degassing. Extrusions made 

from vacuum atomized powder, argon ultrasonically atom1zed powder, helium 

ultrasonically atomized powder, and roller-quenched flakes will be referred to 

by the acronyms VA, USGA-Ar, USGA-He, and RQ, respectively. 

Approximately 9-kg lots of Al-4.4Cu-1.5Mg-1Fe-1Ni-0.2Zr and Al-3Li-1.5CuIMg-0.5Co-0.2Zr VA powders were consolidated 1nto forgings at Alcoa Techn1cal 

Center, Alcoa Center, PA. The powders were canned, vacuum preheated and degassed to 500°C, and hot pressed to 100% dense billets at 500°C. The b11lets 

were upset forged on flat dies.to approximately 55% reduction in thickness at 

400°C. 

The solidus temperature of Al-4.4Cu-l.sMg lies sl1ghtly above 530°C (17); 

thus 530 and 500°C were selected as solution-treatment temperatures for the 

non-L1-containing alloys. The Al-3Li-l.sCu-lMg sol1dus was est1mated by 

extrapolation from the Al-L1, Al-Cu, and Al-Mg b1nary systems (17) to 11e 

below 577°C; thus 560 and 5300C were chosen as solution-treatment temperatures 

for the Li-contain1ng alloy. 

Aging curves were determined for non-L1-conta1n1ng alloy extruS10ns at 

190°C by analogy w1th the known behav10r of 2024-AI. Aging curves were determ1ned for L1-conta1ning alloy extrus10ns at 1770C to yield peak hardness 1n 8 

to 16 h ag1ng time. Natural aging is not appl1cable to the L1-containing 

alloy. The ag1ng behavior of forgings was assumed to resemble that of the 

extrusions. 

10 
3.4 Microstructures and Properties of Consolidated Forms 

3.4.1 Microstructure, Texture, and Dens1ty 

M1crostructures of extruS10ns and forgings were determ1ned by OM and 

TEM. Optical micrographs of all alloys 1n the as-extruded, as-forged, and 

solution-treated condit1ons were obtained to characterize the as-consolidated 

microstructures and the extent of dissolution of major alloying elements and 

recrystallization upon solution treatment. TEM m1crographs of extrusions and 

forg1ngs in var10US tempers were obtained to follow precip1tation and aging 

behavior as a function of particulate type and aging temperature and time, and 

to correlate microstructural variables with mechanical properties data. Since 

OM 1S not sensitive to changes in substructure, {Ill} pole figures of extrusions and forgings in the as-consolidated and solution-treated conditions were 

obtained to characterize changes in substructure during solution treatment. 

Densities of all extrusions were determined using a two-liquid, densitymatching technique developed at Rockwell International Science Center (18). A 

small alloy sample was placed in a mixture of di-iodomethane (CH2I 2, 3.32 

g/cm3) and Neothene (1.30 g/cm3) of approximately the correct density. The 

appropriate 11quid was added until the sample remained suspended. The density 

of the 11quid mixture, equal to that of the sample, was then determined pycnometrically. This method is estimated to yield densities accurate t~ about 

± 0.5%. 

3.4.2 Mechanical Properties 

The elast1c moduli, yield strengths, ultimate tensile strengths, and 

ductilities of alloy extrusions and forg1ngs in various tempers and orientations were determined at temperatures between 25 and 260°C. Determinations of 

tensile properties were made on subsize specimens of 3.lS-mm diameter and 

l5.9-mm gage length according to ASTM ES-Sl. Most data on tensile propert1es 

reported here were based on triplicate tests. 

Fracture toughnesses of peak-aged extrusions and forgings were determined 

in the LT and TL or1entations using subs1ze compact-tension specimens. (The 

symbol L refers to the direction parallel to application of force during 

extrusion or rolling, and T and S refer to the long and short transverse 

directions, respect1vely. In LT and TL, the f1rst letter refers to the 

11 
MDC Q1227 

20 June 1984 

direct~on of appl~cat~on of force during test~ng, and the second letter to the 

direction of crack propagation.) Since the spec~mens were typ~cally less than 

1.27-cm thick, the measured toughness values generally d~d not meet the planestrain cr~ter~on, and are properly referred to as Kq rather than Krc (planestratn fracture toughness) values. Kq values are not directly comparable with 

literature or handbook Krc data. 

Fat~gue-crack-growth-rate (FCGR) data for peak-aged spec~mens from VA 

extrusl0ns and forgings were determined ~n the LT and TL or~entat~ons us~ng 

subsize compact-tension specimens. The test frequency was 5 Hz and the ratio 

of maXlmum to mlnimum load was 0.1. 

Steady-state creep rates and stress-rupture data were determined for 

peak-aged specimens from VA AI-4.4Cu-l.sMg-lFe-lNl-0.2Zr extruslons at IsOoC 

using flat tensile specimens of 2.s-cm gauge length. Tests were conducted 

under loads corresponding to 70, 80, and 90% of the Is00C yield stress of each 

alloy. 

The stress-corrosion-cracking (SCC) behavior of peak-aged VA extruSlons 

and forgings was determined by alternate-immersion testing of bolt-loaded, 

subsize C-ring specimens in 3.5% NaCI solution accordlng to ASTM G38-73 and 

G44-7s. Forging samples were tested with crack growth in the S directlon, but 

the extrusions were not thick enough to permit S-orientation testing and hence 

were tested with crack growth in the T direction. Loads applled to the specimens corresponded to 25, 50, and 75% of the ambient-temperature Yleld 

stresses. 

12 
4. ~rASK 1: PRODUCTION AND CHARACTERIZATION OF 

RAPIDLY SOLID IF IUD POWDERS AND ]i'LAl{ES 

4.1 Evaluation (Jlf Uispers(Jlid-Forming Systems for Non-Li-Containing Alloy 

Roller-quenched flakes were prepared from I-kg heats of the prototype 

dispersoid-forming systems Al-lFe-INi and Al-4Ti-l.8B. Flakes of Al-lFe-1Ni 

were successfully prepared at a pouring temperature of 7300 C (7S0C superheat). 

Chemical analysis of the flakes yielded an actual composition of Al-0.96Fe1.02Ni. Optical microscopy of etched surfaces of as-solidified flakes 

(Figure 4) shows a cellular/dendrit:Lc structure with 10- to 30-)lm diameter 

cells and an average dendrite-arm spacing (DAS) of 0.6 )lm. Using the correlation of Matyja ~ al~ (19) between DAS and cooling rate, the solidification 

rate was estimated to be '" 4 x 105 K/ s. 

'----J 

10 p.,m 

Figure 4. Optical micwgraph of AI-1F'e-lNi flake. 

(;1'41·0922·13 

The evolution of the as-solidified microstructure into one more nearly 

typical of a consolidated form was approximated by heat treating the flakes 

for 1 h at 400, 500, and 550oC. TEM micrographs of as-solidified and as-heat 

13 
treated flakes (Figure 5) show progressive spheroidization as the temperature 

is raised, to a final distribution of approximately 0.2- x 0.3-pm oblate dispersoids at the boundaries of I-pm--diameter subgrains (Figures Sc and Sd). 

Spheroidization is essentially complete at 500°C. The final distribution of 

dispersoids is promising for significant dispersion strengthening. 

(a) (b) 

(c) (d) 

Ip,ill 

1----1 

Ip,ill 

l----1 

Ip,ill 

(;1'41-0922-14 

J<igure S. Microstructures of roilier-quenched AI-lFe-lNi flakes; (a) as-quenched and heat treated 

for 1 h at (b) 400°C, (c) 500°C, and (d) 550°c' 

14 
X-ray and electron d1ffract1on were used to ident1fy the disperso1ds 1n 

AI-lFe-lN1 flakes. Results of X-ray d1ffraction scans show that the dispers01ds 1n both as-solid1fied and as-heat-treated flakes are predominantly 

A1 9FeN1, which has the monoclinic A1 9Co2 structure. Identification of TEM 

electron diffraction spots and comparison of measured and pred1cted angles 

between var10US d1rections in the reciprocal latt1ce support the 1dent1f1cat10n of d1spersoids as AI gCo 2-type. Both X-ray and electron diffraction 

patterns suggest that traces of other phases may be present in as-sol1dified 

flakes. 

Reference to the Al-T1 and Al-B phase d1agrams (17) suggests that a 

temperature of 10000C is required to d1ssolve 4Ti and 1.8B in an alum1num 

melt. Droplets of Al-4Ti-l.8B heated to a pour temperature of 10000 C in the 

roller-quenching apparatus, were not quenched to solid flakes by the rolls. 

Attempting to pour at a melt temperature of 750 to 8000C resulted in retention 

of unmelted master alloy in the bottom of the crucible which blocked the crucible orifice and prevented flake formation. Since the Al-lFe-lN1 system produced satisfactory results, no further tests of the AI-Ti-B system were made, 

and the Fe-Ni system was selected as the d1sperso1d-forming add1tive for the 

non-Li-containing alloys. 

4.2 Production of Rapidly Solidified Particulates 

4.2.1 Vacuum Atomization 

Ingots of 33 kg mass of both Al-4.4Cu-l.5Mg-1Fe-lNi-O.2Zr and Al-3Li1.5Cu-lMg-O.5Co-O.2Zr were vacuum atomized 1n eight heats. The yield of -80 

me~h powder was 18.9 kg for the non-Li-conta1n1ng alloy and 17.7 kg for the 

Li-containing alloy; these were sufficient to manufacture extruS10ns and forgings large enough to perform all required microstructure and mechanical properties evaluat1ons. HMI also produced 5.9 kg of -80 mesh Al-4.4Cu-l.5Mg-2Fe2N1-0.2Zr powder. 

4.2.2 Ultrasonic Atomization 

Suff1cient quantities of Al-4.4Cu-l.5Mg-lFe-lNi-0.2Zr and Al-3Li-l.5Cu1Mg-0.5Co-0.2Zr powders for producing extrusions were manufactured by ultrason1C atomization 1n argon and 1n helium. Operat1ng parameters for atomiza15 
tIon, and powder YIelds, are summarIzed In Table 1. Although no systematIc 

study of powder characterIstics as a function of operatIng parameters was 

attempted, a combInatIon of 1000C superheat and 9.7 MPa (1400 pSIg) delIvery 

pressure yielded satIsfactory results. The YIeld of usable powder was limIted 

by formatIon of a slntered powder mass at the bottom of the atomIzatIon chamber, and attempts to reduce the slnterIng problem by decreaSIng the melt temperature from 820 to 730-750oC were only partIally successful. The powders 

were sIeved to -80 mesh to elIminate splats and stored In sealed containers 

under argon pending further proceSSIng. 

TABLE 1. OPERATING PARAMETERS FOR ULTRASONIC ATOMIZATION OF 

AI ALLOY POWDERS. 

Initial - 80 mesh ( - 180 I'm) 

melt temp- Gas delivery powder yield, 

Atomizmg Run erature pressure (kg [lb], 

Composition gas no. (OC) (MPa[psig)) 070 of ingot weight) 

A1-3L1-1 5Cu-1Mg-0 5Co-0 2Zr Ar 1 820 83 [1200] 

2 770 97 [1400] 54[119],36 

3 750 97 [1400] 

A1-3Ll-l 5Cu-1Mg-0.5Co-0.2Zr He 1 750 97 [1400] 

2 750 97 [1400] 2 8 [6 2], 19 

3 750 97 [1400] 

AI-4 4Cu-l 5Mg-IFe-1NI-0 2Zr Ar 1 750 97 [1400] 

2 750 97 [1400] 4 2 [9 2], 37 

AI-4 4Cu-l 5Mg-IFe-1NI-0 2Zr He 1 730 97 [1400] 

2 730 97 [1400] 30 [6 5], 27 

GP41-0922 I 

4.2.3 Twin Roller Quenching 

Approximately 5.1 kg of roller-quenched Al-3LI-1.5Cu-1Mg-0.5Co-0.2Zr 

flakes were produced for manufacturIng extruSIons. The flakes were chopped to 

particulates approximately 0.2-cm x O.2-m x l25-~m thICk in a commercIal paper 

shredder and blended to ensure compositional uniformIty In the final extrusion. Attempts to produce large quantItIes of Al-4.4Cu-l.5Mg-lFe-lNI-0.2Zr 

flakes for extrusions resulted in low yields because of excessive roll pItting 

and degradation. Production of flakes for extrUSIons was abandoned In favor 

of production of AI-4.4Cu-l.5Mg-2Fe-2Ni-0.2Zr vacuum atomIzed powder. 

16 
4.3 Particulate Characterization 

4.3.1 Chemical Analyses 

Chemical analyses of all particulates are listed in Tables 2 and 3, and 

hydrogen and oxygen analyses in Table 4. Vacuum-atomized powders are generally clean, with no evidence of cross contamination with Ni-base powders (which 

TABLE 2. CHEMICAL ANALYSES OF NON·Li·CONTAINING ALLOY PARTICULATES. 

Composition (wt 070) 

Designation Cu Mg Fe Ni Zr Cr Ti 

Nommal composItion 4.4 1.5 1.0 1.0 0.2 

Vacuum atomized powder, Lot A 386 1 50 0.94 098 0.23 0015 0.004 

Vacuum atomized powder, Lot B 4.28 1.56 0.92 0.94 0.20 0.003 0002 

Vacuum atomized powder, Lot C 409 1.45 0.94 0.86 0.18 0.003 0.003 

Vacuum atomized powder, Lot D 3.91 1.50 0.92 0.95 0.20 0004 0.002 

Ultrasomcally atomized powder - argon 4.16 088 0.96 095 0.18 0004 0.002 

Ultrasomcally atomized powder - helIum 4.24 095 0.92 094 0.20 0004 0.003 

Roller-quenched flakes 4.29 1 51 0.82 1 03 0.18 NA N.A. 

Nommal composItIon 4.4 1.5' 20 2.0 0.2 

Vacuum atomized powder 4.21 1 03 1.76 1 81 0.22 0006 0.003 

SI, Zn < Olin all particulates 

GP410922·2 

TABLE 3. CHEMICAL ANALYSES OF Li·CONTAINING ALLOY PARTICULATES. 

Composition (wt %) 

Designation Li Cu Mg Co Zr Fe Ni 

Nommal composItIon 30 1 5 I 0 05 02 

Vacuum atomized powder, Lot A 310 1.46 106 049 o 19 0018 0005 

Vacuum atomized powder, Lot B 308 146 104 049 o 19 0021 0007 

Vacuum atomized powder, Lot C 320 146 1 03 050 o 19 0021 0030 

Vacuum atomized powder, Lot D 3 18 1 43 1 03 052 o 19 0041 024 

Ultrasomcally atomized powder - argon 276 1 54 072 047 025 0008 0011 

Ultrasomcally atomized powder - helIum 2.72 144 062 0.47 o 17 0003 0008 

Roller-quenched flakes 310 1 59 1 07 040 010 006 0011 

Cr, < 0 014 (VA Lot 0, 0 OS3), n, < 0 01S (VA Lot 0, 0 020), SI, < 0 OS, Zn < 0 10 In all particulates 

GP41·09223 

17 
TABLE 4. HYDROGEN AND OXYGEN ANALYSES OF PARTICULATES AND EXTRUSIONS. 

Hydrogen Oxygen 

Alloy composItIon concentration (wppm) concentratIOn (wppm) 

and atomizatIon method Particulates ExtrusIOns Particulates ExtrusIOns 

AI-4.4Cu-l.5Mg-lFe-lN I-O.2Zr 

Vacuum atomizatIOn 27±2 2 370± 130 91O± 30 

UltrasonIc atomizatIOn - argon 43±2 4 440± 120 570± 60 

UltrasonIc atomization - helium 35±2 I 350± 110 760± 60 

Roller quenchmg (flakes) 5±1 

AI-4.4Cu-l.5Mg-2Fe-2NI-O.2Zr 

Vacuum atomization 5 1280± 10 

AI-3Li-l.5Cu-lMg-O.5Co-O.2Zr 

Vacuum atomization 164±2 <20 200±60 480±60 

UltrasonIc atomization - argon 234±2 2 280± 110 650± 140 

UltrasonIc atomizatIOn - helium 197 ± 5 18 280± 130 1320± 300 

Roller quenchmg (flakes) 31 ± 1 4 15 ± 10 160±40 

(.P41 09224 

are made at HMI in the same apparatus) except for s1gnificant, but not deleterious, concentrations of Ni, Cr, and Ti 1n Lot D of the Li-containing alloy. 

Compositions of vacuum-atomized powders are close to nominal, except for 

approximately 10% low Cu concentrations in the non-L1-contain1ng alloy. 

Both USGA and RQ particulates made at MDRL have low contaminant levels. 

The ingots supplied by HMI for ultrasonic atomization did not have suff1ciently high Li and Mg concentrations to compensate for the 10% Li and 30% Mg 

losses caused by vaporization or oxidation in the conversion to powder. For 

roller-quenched flakes, Li and Mg concentrations in the initial melts were 

increased by 5 to 7% over nominal to produce the correct concentrations of 

these elements in the flakes. 

Hydrogen concentrations in as-solidified Li-containing alloy particulates 

are consistently 5.5-6 times larger than those in the non-Li-containing alloy, 

regardless of particulate type. This result has been attributed to the h1gh 

reactivity of lithium with water vapor. Hydrogen concentrations 1n as-solidified roller-quenched flakes are about one-sixth as large as those in powders, 

because of the smaller surface area per unit volume of the flakes. Oxygen 

contents of the particulates are low compared to similar alloys produced under 

~DRL Independent Research and Development (20), which had initial 

18 
concentrations as high as 4000 ppm, and surpris1ngly slightly lower 1n the Licontaining than in the non-Li-contain1ng alloy. 

4.3.2 Particle Size Distribution 

Particle size distributions of Al-alloy powde~s are plotted 1n F1gures 6 

and 7. All the distributions follow log-normal behavior. Weight-averaged 

part1cle d1ameters correspond1ng to 5, 50, and 95% smaller than a given 

diameter are summarized for each composition and atomizat1on method 1n Tables 

5 and 6. 

Particle size distributions of both alloys are similar for both VA and 

USGA-Ar. The vacuum atomized powder has a smaller mean particle size (38 to 

40 ~m) than the argon atomized powder (62 to 70 ~m), and a narrower Slze distribution. The USGA-He powder has the w1dest size distribution, and a smaller 

average particulate diameter in the Li-containing alloy (29 ~m) than 1n the 

non-Li-contain1ng alloy (52 ~m). This size difference probably results from a 

higher melt temperature for the Li-containing alloy (see Table 1), consistent 

999r---~----~--~--~ 

95 

... 

..!:! ... 80 - Q) CIS ... E Q) 

60 '" E ... CIS 

C - Q)"O 40 u c ... Q) Q) :> 0._ 

... 00 20 ..c c OOcu 10 o..c 

~ ... 2 

o 1 

10 20 50 100 200 

Particle diameter (p.m) 

Figure 6. PartIcle size distributIons of AI-4.4Cu-l.SMg-lFe-lNi-O.2Zr powders: (0) Vacuum 

atomized, (0) argon ultrasonically atomized, and (t;.) helium ultrasonIcally atomized. 

(.1'41-0922 15 

19 
999r---~----'---~---. 

95 

80 

60 

40 

20 

10 

2 

o I 

001~--~----~--~--~ 

10 20 50 100 200 

Partlcle diameter (Jlm) 

Figure 7 Particle size dlstnbutlOns of AI-3Li-l.5Cu-lMg-O.5Co-O.2Zr powders: (0) Vacuum 

atomized, (0) argon uitra'iomcalh atomized, and (t,) helium ultra'iomcall~ atomized. 

(,1'41 0'122 1~ 

TABLE S. MEAN PARTICLE DIAMETERS, DENDRITE ARM SPACINGS, AND 

SOLIDIFICATION RATES FOR AI-4.4Cu-l.5Mg-lFe-lNI-O.2Zr PARTICLES. 

Geometnc mean 

particle diameter (Jlm) 

Standard deviation 

Dendnte arm spacmgs for 

mean and ± one standard 

devIatIOn partIcle SIzes (Jlm) 

ApproxImate solIdIfIcatIon 

rates for mean and ± one 

standard devIatIOn partIcle 

sizes (K/s) 

Vacuum 

atomization 

40 

1 74 

2 57, 3 30, 1 72 

8 1 X 103 

3 9x 103' 

27xIQ4' 

20 

Atomization method 

Ultrasonic atomization Ultrasonic atomization 

- argon - helium 

64 52 

203 244 

2 55, 3 63, 1.38 1 29, 2 53, 0 63 

83 X 103 63x104 

29x103' 85 X 103' 

5 1 x IQ4' 5 3 x lOs' 

GP41-0922 S 
TABLE 6. MEAN PARTICLE DIAMETERS, DENDRITE ARM SPACINGS, AND 

SOLIDIFICATION RATES FOR AI-3Li-1.5Cu-lMg-O.5Co-O.2Zr PARTICLES. 

Geometric mean 

particle dIameter (JLm) 

Standard devIatIOn 

Dendnte arm spacmgs for 

mean and ± one standard 

devIatIOn particle SIzes (JLm) 

ApproxImate solIdIfIcatIOn 

rates for mean and ± one 

standard deVIation partIcle 

SIzes (K/s) 

Vacuum 

atomization 

38 

1 70 

241,351, 1 56 

98 X 103 

3 2x 103' 

3 6x 1()4' 

Atomization method 

Ultrasonic atomization 

- argon 

70 

1.92 

3 09, 3 66, 1 93 

4 7x 103 

28x103' 

19x1()4' 

Ultrasonic atomization 

- helium 

29 

2.43 

o 84, 1 57, 0 45 

2 2x lOS 

3 5 x 1()4' 

1.4 x 1()6' 

GP41-09226 

with observations by Grant (9). The results also suggest that particulate 

size distributions in rapidly solidified aluminum alloys are primarily funct10ns of atom1zation method, atomizing gas pressure, and melt superheat, but 

not of compositional d1fferences. 

4.3.3 Metallographic Powder Characterization 

Optical micrographs of as-atomized alloy powders are shown in Figures 8 

and 9. All powders are predominantly spherical, regardless of composition or 

atomization method, and have dendr1tic or cellular microstructures. The VA 

powders (Figures 8a and 9a) have larger numbers of satellites (small particles 

bonded to larger ones, top center Figure 8a) and splat caps (resulting from 

collision of a sol1d particle with a l1quid droplet, right center Figure 9a) 

than do USGA powders, in agreement with previous observations (9). The presence of satellites and splat caps can impair particle packing during consolidation and decrease microstructural homogeneity in consolidated forms, thus 

degrading mechanical properties (9). 

All the powders have occasional porous particles (upper right in 

Figure 9b). Such particles are rare except in USGA-He Al-3Li-l.5Cu-lMg-O.5Co21 
O.2Zr, where most particles contain internal voids. Completely contained gas 

bubbles are undesirable because they may lead to high concentrations of IT and 

o in consolidated forms. 

20/.-tm 20 p.,m 

(c) 

20 p.,m 

(;1'41-0922-17 

:Figme :~. Optical of AI·4.4Cu··l powders: (a) vacuum 

(b) ultrllsonically atomized in arg:on, and (c) ultrasonically atomized in helium. 

22 
(a) 

L_,.J 

(c) 20 ron 

(b) 

(d) 

LJ 

20 {-tm 

LJ 

250 {-tm 

(;1'41·0922·18 

Figure 9. Optical micrographs of AI-3U-1.SCu-lMg-O.SCo-O.2Zr powders: (a) vacuum atomized, 

(b) ultrasonically :atomized in argon, (c) and (d) ultrasonically atomized in helium. 

23 
4.3.4 SEl1 Characterization and Dendrite Arm Spacing Measurements 

SEM Inicrographs of typical VA, USGA-Ar, and USGA-He particles are shown 

in Figure 10. Cellular microstructures are clearly visible on the unetched 

particle surfaces. Major microstructural features are similar to those 

observed optically (Figures 8 and 9), but since the full particle diameter is 

visible in the SEM micrographs, dendrite-arm spacings and particle diameters 

can be unambiguously related. 

(a) 

(c) 

'--___ -.J 

10 !lm 

L--____ ....J 

10 !lm 

(b) 

10 !lm 

GI'41·0922·19 

Figure ]0. Scanning electron micrographs of AI-4.4Cu-1.5Mg-]Fe-]Ni-0.2Zr particles prepared by 

(a) vacuum atomization, (b) ultrasonic atomization in argon, and 

(c) ultrasonic atomization in helium. 

24 
plots of dendrite-arm spacing (DAS) as a function of particle d1ameter 

for each alloy composition and atomization method are shown in Figures 11-16. 

Each data point represents the average of 12-20 DAS measurements on a single 

particle, and the solid curve on each graph 1S the best three-parameter polynomial fit to the data. S1milar curves have been obtained for aluminum alloy 

powders by Grant (9). 

For each alloy composition, the DAS curve for helium ultrasonic 

atomization falls well below that for argon ultrasonic atomization, wh1ch in 

turn falls somewhat below that for vacuum atomization. Thus a USGA-He powder 

particle of mean diameter has been sol1dified more rap1dly than a similar 

USGA-Ar particle, and the mean USGA-Ar particle has been solid1fied more 

rap1dly than a similar VA particle. This progression is in accord with the 

expected solid1ficat1on rates based on the relative efficiencies of helium, 

argon, and reduced-pressure argon (1n "vacuum" atomization) as heat-transfer 

med1a. 

5r-----~------~-----,------~------r_----_r----__, 

4 

.-. 

e 

5 

co 

c::: 3 13 

tV 

Co 

en 

e 

.. 

tV 

U 

-~ 

'tl 

c::: 

u 

0 

140 

Particle diameter (Jlm) GP41-0922-20 

Figure 11. Dendrite arm spacing as a function of particle diameter for AI-4.4Cu-1.5Mg-IFeINi-O.2Zr vacuum atomized powder. 

25 
5r-----~----~~----~----~------~----_,----__, 

4 

--8 

.=. 

OIl 

:= 3 U 

<II 

0. 

til 

8 

.... 

<II 

I\) 2 

-::: 

.... 

"C 

:= 

I\) 

0 

20 40 60 80 100 120 140 

Particle diameter (JLm) GP41-0922-21 

Figure 12. Dendrite arm spacing as a function of particle diameter for AI-4.4Cu-1.5Mg-IFeINi-O.1Zr argon ultrasonically atomized powder. 

5~-----r------~----~------~----_'------~-----' 

4 

--8 

.=. 

1:)1) 

:= 3 

U 

<II 

0. 

til 

8 

.... 

<II 2 I\) 

..... 

;:: 

"C 

:= 

I\) 

0 

10 20 30 40 50 60 70 

PartIcle diameter (JLm) GP41-0922-22 

Figure 13. Dendrite arm spacing as a function of particle diameter for AI-4.4Cu-1.5Mg-IFeINi-O.1Zr belium ultrasonically atomized powder. 

26 
,-. 

e 

:::t 

'-' 

01) 

= 0 

(\I 

c. 

'" e 

.... 

(\I 

CIJ 

<.::: 

.... 

'0 

= CIJ 

Q 

,-. 

e 

,:; 

01) 

= 0 

(\I 

c. 

'" e 

.... 

(\I 

CIJ 

.. 

t: 

'0 

= CIJ 

Q 

5~----~----~------~-----r----~r-----'---n--' 

2 

140 

PartIcle dIameter (J.Lm) GP41-4911-13 

Figure 14. Dendrite arm spacing as a function of particle diameter for 

AI-3Li-l.5Cu-lMg-O.5Co-O.1Zr vacuum atomized powder. 

5r-----~------~----~------~----~r------r----~ 

4 

3 

2 

20 40 60 80 100 120 140 

PartIcle dIameter (J.Lm) GP41-092l-14 

Figure 15. Dendrite arm spacing as a function of particle diameter for 

AI-3Li-l.5Cu-lMg-O.5Co-O.2Zr argon ultrasonically atomized powder. 

27 
5r-----,------.------r-----1I----~------._----~ 

4 ,-.. 

e 

::t 

'-' 

t:>I) 

c: 

U 3 ~ 

0. 

CIl 

e 

.... 

~ 

11) 2 !: 

.... 

'1:1 

c: 

11) 

0 

140 

Particle diameter (J.tm) GP41-09ll-15 

Figure 16. Dendrite arm spacing as a function of particle diameter for 

AI-3Li-l.5Cu-1Mg-O.5Co-O.2Zr helium ultrasonically atomized powder. 

Average dendrite-arm spacings for particles of mean diameter and one 

standard deviation larger or smaller than the mean diameter are tabulated in 

Tables 5 and 6. (A factor of one standard deviation around the mean covers 

the range between 16 and 84% smaller than a given diameter.) These dendritearm spacings have been converted to approximate solidification rates using the 

correlation developed by Matyja et al. (19). For Al-4.4Cu-1.5Mg-1Fe-1Ni-0.2Zr 

(Table 5), average particle sizes for powders atomized by all three methods 

are roughly similar. The inherently larger solidification rate obtainable 

from USGA-He causes the average USGA-He particle to have a sol1d1fication rate 

almost an order of magnitude larger than that for USGA-Ar or VA powders 

(60 000 Kls vs. 8000 K/s). This difference is more pronounced for fine and 

less pronounced for coarse powders, owing to the w1de S1ze d1str1bution of 

USGA-He powder. The smaller average diameter of VA powder is compensated for 

by the higher solid1f1cation rate of USGA-Ar powder to yield approximately 

equivalent solid1ficat10n rates for particles of mean d1ameter prepared by 

these two methods. 

28 
For Al-3Li-1.5Cu-1Mg-0.5Co-0.2Zr, the USGA-He powder has both a smaller 

average part1cle size and a larger inherent solidification rate than do the 

other powders, and hence has an average solid1fication rate for particles of 

mean diameter about 20 times larger than that of VA powder and about 50 times 

larger than that of USGA-Ar powder. The smaller average d1ameter of VA powder 

more than compensates for the larger inherent sol1d1fication rate of USGA-Ar 

powder to yield a weight-averaged solid1fication rate about twice as large for 

the former. 

4.3.5 Ox1de Layer Thickness Characterization 

An Auger electron spectrum of the surface of an as-received VA Al-4.4Cu1.5Mg-1Fe-1Ni-0.2Zr particle (Figure 17a) shows significant carbon contamination from handling and atmospheric exposure, and an oxide rich in Mg and Al 

and relatively poor in Cu (compare the relative Mg and Cu heights in 

F1gure 18a with the nominal composition of the alloy). After obtaining the 

initial spectrum on this particle, the surface was argon ion sputtered at an 

approximate rate of 0.7 nm/min, and Auger peak intensities for Al, Cu, Mg, 0, 

and C were obtained after each 30 s exposure. The plot of peak intensity as a 

function of depth (Figure 17b) shows that the oxide film is about 1.7 nm thick 

and is strongly enriched in Mg. (Also compare Figures 17a and 17c.) This 

thin oxide layer corresponds well to the < 0.04 wt% oxygen concentration in 

the powder (Table 4). Copper is rejected from the film, in accord with the 

relative thermodynamic stabilities of copper and magnesium oxides. After 

extended sputtering, the Auger spectrum of the matrix (Figure 17c) shows 

strong Al and Cu peaks, a weak Mg peak near 1200 eV, and virtually no C or o. 

(The indium signal originates from In atoms from the substrate, deposited onto 

the particle during sputtering.) 

The initial spectrum of an as-received VA Al-3Li-1.5Cu-1Mg-O.5Co-O.2Zr 

particle (Figure 18a) is similar to that of as-received VA Al-4.4Cu-1.5Mg-1Fe1Ni-O.2Zr (Figure 17a), with a strongly Mg-enriched oxide on the surface. The 

depth profile (Figure l~b) shows that the oxide is about 3.5 nm thick. The 

oxide is enriched 1n Mg throughout and may be enriched in Li, although the 

major Auger peak for Li is weak and d1fficult to evaluate quantitatively. The 

matrix spectrum (F1gure 18c) resembles that of the non-Li-containing matrix 

(Figure 17c) except for a larger residual oxygen concentration. 

29 
7 

~I 

Al C 

o 

o 200 400 

o 1 4 

7 

6 

5 

4 

3 

2 

(s) 

I __ Ill .J. _"'fV.LL.1 .JI.J.. _I ••. M ........ , .. ~.\" rlJr~ ".'" , ..... --? 11'""" -, ~r'N ,. 1 r" ... ··l 

Cu 

0 

. Mg 

Al 

> 

600 800 1000 1200 1400 1600 1800 2000 

28 

Kmetlc energy (e V) 

(b) 

Sputter depth (nm) 

4.2 5.6 7.0 

GP41-0922 26 

84 98 112 

o~-+----~----+---~~---+----~----+---~~---+~ 

o 2 4 6 8 10 12 14 16 

Sputter time (mm) GP41-0922-27 

Figure 17. Scanning Auger Microscope (SAM) results for as-vacuum-atomized AI-4.4Cu-1.5MgIFe-lNi-O.2Zr particle: (a) Auger electron spectrum of as-received particle, (b) depth 

profile of AI, Cu, Mg, 0, and C concentrations, sputtermg rate 0.7 nm/min, and (c) 

Auger electron spectrum of underlying matrix. 

30 
(c) 

7 

.ifJ JJ ~~~~ J • .L. 1V .11 I' .&11 'If . ·ff~1 1~~ r~ ,...- T ''for rrl r 

Fe Cu 

In 

Al 

Al 

> 

Kmetlc energy (e V) GP41.o9ll 18 

Figure 17. (Concluded) 

(a) 

7 

6 

5 

~J ~J !,ill, 1',; AJ IJJ~ f fI"~JJ l~ II "1/111 ,1~ 11'-' ,# I'JIIIIJllf"J~ "111111 I lPllj~\'rI~'1"'ir'lrrrll"ll Ill"·''''1 \tl"'VI)''tI~ 1II'It l'I"1I11"n 

! C o Mg 

4 

3 

Al 

Al - 2 LI--

> 

Kmetlc energy (e V) LP41.0922-29 

Figure 18. SAM results for as-vacuum-atomized AI-3Li-l.5Cu-IMg-O.5Co-O.2Zr particle: 

(a) Auger electron spectrum of as-received particle, (b) deptb profile of AI, Cu, Mg, 

LI, 0, and C concentrations, sputtering rate 0.7 nm/min, and (c) Auger electron 

spectrum of underlying matrix. 

31 
(b) 

Sputtering depth (nm) 

o 1 4 28 42 56 70 84 9 8 11 2 12 6 14 0 

7 

;Z=::-" ~'~. ------~--- Al 

> 

0~~--~--_+---4--~----~--+_--4_--~--~--_4 

o 2 3 4 5 6 7 8 9 10 

Sputter time (mm) GP41-0922lC1 

(c) 

7 

... J. -d .'dI,..}JJI: ,~ ....... I .. ~" ... d., .1 . 

".......... ..uf,"",,~ I rr 'H'"" 'fir -T .... ... TO t". 1"1/1' T 1 , 

In r I 

CU Mg 'I 

6 

5 

4 

0 

Al 

3 

2 -AI 

> 

o 

KmetIc energy (eV) (,P41-0922-31 

Figure 18. (Concluded). 

32 
Depth profiles of ox~de films on USGA-He alloy part~cles are shown ~n 

Figure 19. The oxide layers are about 3 nm thick for Al-4Cu-1.5Mg-IFe-1Ni7 

6 

5 

4 

3 

2 

o 

7 

6 

5 

4 

3 

2 

o 1 5 

lAI 

/ 

:/ 

V 

~ leu 

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

(a) 

Sputtermg depth (nm) 

3 0 4.5 

, 

~ , 

I I 

~ t-... -----~ 1 I IMg 

o 

o 1 5 

2 3 

Sputter time (mm) 

(b) 

Sputtenng depth (nm) 

3.0 4.5 

60 7.5 

I 

I 

J. 

4 5 

GP41~12-32 

60 75 

o ~-+--------~--------~------~--------+--------4 

o 2 3 4 5 

Sputter time (mm) GP41~22-33 

Figure 19. Depth profiles of oxide films on helium atomized particles. Sputter rate 1.5 nm/min: 

(a) AI-4.4Cu-1.5Mg-IFe-1Ni-O.2Zr and (b) AI-3Li-1.5Cu-1Mg-O.SCo-O.2Zr. 

33 
O.2Zr and 4.2 nm thick for Al-3Li-l.5Cu-lMg-O.5Co-O.2Zr. Ox~de f~lm 

thicknesses are only sl~ghtly greater in Li-conta~ning than ~n non-L~­

conta~n1ng alloy powders, suggest~ng that the in~tially formed ox~de layer ~s 

passive and does not grow rapidly at amb~ent temperature, even when the 

react~ve element L1 is present. 

34 
5. TASK 2: CONSOLIDATION AND POST-CONSOLIDATION 

PROCESSING OF PARTICULATES 

5.1 Particulate Degassing Experiments 

S.l.l Removal of Hydrogen and Oxygen 

Hydrogen and oxygen concentrat10ns in powder samples, tray degassed in 

vacuum at 400-S30oC, are shown in Figures 20 and 21. Efficiency of hydrogen 

removal from the non-L1-containing alloy 1ncreases with increasing degassing 

temperature, but res1dual hydrogen concentrations in the Li-containing alloy 

pass through a minimum and 1ncrease sharply above 4S00C (Figure 21). This 

result suggests competition between two processes in vacuum degassing of the 

Al-3Li-1.SCu-1Mg-0.SCo-0.2Zr alloy: removal of initially adsorbed hydrogenbearing species by the vacuum, and reaction of such species with Li in the 

alloy to chemically fix hydrogen. If the temperature dependence of the rate 

of the hydrogen-fixing chemical reaction exceeds that of physical degassing, 

the latter reaction is more rapid at lower temperatures and the H concentration decreases as the degassing temperature is raised. At still higher temperatures, more of the hydrogen is chemically fixed than physically removed, 

and the H concentration increases with increasing temperature. The non-Licontaining alloy lacks a reactive chemical species to fix H, and degassing 

becomes continuously more efficient as the temperature is raised. 

Oxygen concentration data (Figure 21) show that oxygen is not effectively 

removed in tray degassing of either alloy powder. Within large uncertainty 

limits, the final 0 content of the non-Li-containing alloy is independent of 

degassing temperature and somewhat larger than that in the original powder. 

This result is consistent with scanning Auger microscopy (SAM) stud1es d1scussed in Section S.1.2. Degassing at temperatures below soooe approximately 

doubles the oxygen content in the Li-containing alloy powder. The sharp 

increase in oxygen concentration during 5300 C degassing suggests that the 

powder reacts rapidly with oxygen at this temperature, lead1ng to a rapid 

1ncrease 1n the oxide layer thickness, a result in agreement with the SAM 

f1ndings d1scussed below. 

35 
E 

c. 

c. 

'-' 

s:: 

9 

... 

t1:S 

... 

... 

s:: 

II) 

u 

s:: 

0 

u 

s:: 

II) 

tID 

>. 

~ 

0 

-; 

... 

0 

EE 

c. 

c. 

'-' 

s:: 

9 

... 

t1:S 

... 

... 

s:: 

II) 

u 

s:: 

0 

u 

s:: 

II) 

tID 

0 

... 

'0 

>. 

..s:: 

-; 

... 

0 

E280 

250 

200 

150 

100 

50 

500 

Degassmg temperature (0C) 

525 550 

GP41~21-J4 

Figure 20. Variation witb degassing temperature of bydrogen 

concentration in vacuum-atomized alloy powders: 

(0) AI-4.4Cu-l.5Mg-lFe-lNi-0.2Zr and 

([J) AI-3Li-l.5Cu-lMg-O.5Co-O.2Zr. 

l~r--T~~----~------~-----r----~r-----,,----~ 

800 

600 

400 

200 

o~~~~------~------~----~------~------~----~ o 550 

Degassmg temperature (0C) GP41~22-35 

Figure 21. Variation witb degassing temperature of oxygen concentration in vacuum-atomized 

alloy powders: (0) AI-4.4Cu-l.5Mg-lFe-lNi-0.2Zr and 

(D) AI-3Li-l.5Cu-lMg-0.5Co-O.2Zr. 

36 
5.1.2 OX1de Layer Mod1f1cation Dur1ng Degassing 

Depth prof1les of oxide films on non-Li- and Li-conta1ning VA particles, 

degassed for 2 h at 5300 C, are shown ~n Figures 22a and 22b. Comparison of 

the depth prof~le of a degassed non-Li-containing alloy particle (Figure 22a) 

with that of a corresponding as-received particle (F~gure 17b) shows that the 

oxide layer thickness has increased from about 1.7 nm to about 2.3 nm, presumably by chemical reaction of oxygen from adsorbed water vapor. Th~s result is 

in agreement with the modest increase in oxygen concentration upon 5300 C 

degassing shown in Figure 21. 

The Li-conta~n~ng particle shows an increase in oxide thickness upon 

degassing from ~ 3.5 to ~ 19 nm (compare Figure 22b with Figure 18b). This 

result is in agreement with the large rise in oxygen concentrat~on in the 

5300 C-degassed powder shown in Figure 21. Although no Li signal was obtained 

in the depth profile, the greater reactivity of the Li-containing alloy is 

presumably the result of the high chemical affinity of Li for oxygen and the 

high mobility of the Li ion in diffusing through an oxide layer to the free 

surface to enable further oxidation. In both alloys, the oxide layer after 

degassing remains rich in Mg, while Cu is rejected into the matrix. 

Oxygen and hydrogen concentration results and the SAM data have important 

implications for powder degassing. It is important, and vital in the Li-containing alloy, to remove as much adsorbed gas at as low a temperature as possible. Rap~d heating of the powder to the degassing temperature results in 

reaction of significant amounts of adsorbed gases such as H20 with the powders, liberating H2 and increasing the oxide layer thickness. The resulting 

thick oxide layers are d~ff~cult to fracture during consolidation, thus increasing the probability of poor interparticulate bonding. Degassing cycles 

should thus include extensive exposure to vacuum at 25-3000 C, where desorption 

rather than chemical reaction can be expected to occur. 

5.2 Characterization of Extrusions and Forgings 

5.2.1 Chemical Analyses 

Chemical analyses of extrusions made from various particulate types are 

consistent with analyses of the corresponding powders (Tables 2 and 3), indicating little additional loss of volatile or reactive elements and negligible 

37 
(a) 

Sputter depth (nm) 

o 1 5 3.0 4.5 60 7 5 90 10 5 120 13 5 15 0 

7 

6 

o L-~ ____ ~ ____ +-__ ~ ____ ~ ____ ~ __ ~ ____ ~ ____ +-__ ~ ____ -+ 

o 2 3 

o 3 6 9 

o 2 4 6 

4 5 

Sputter time (man) 

(b) 

6 

Sputter depth (nm) 

7 8 9 10 

('P41-0922 36 

12 15 18 21 24 27 30 33 36 

8 10 12 14 16 18 20 22 24 

Sputter time (man) ('P41-0922 37 

Figure 22. Depth profiles of oxide films on vacuum-atomized particles, degassed 530°C/2 h ID 

vacuum. Sputter rate 1.5 nm/min: (a) AI-4.4Cu-l.SMg-1Fe-1Ni-O.2Zr and (b) AI3Li-l.5Cu-lMg-O.SCo-O.2Zr. 

38 
1ntroduction of impurities during the consolidation process. The chemical 

analyses shown in Tables 2 and 3 are thus valid for the extrusions as well as 

the part1culates. 

Hydrogen analyses of the extrusions (Table 4) show that degassing was 

generally efficient and reduced hydrogen concentrations to acceptable levels 

except in the case of the extrusion made from the USGA-He Li-containing alloy. 

The average particle size of the latter alloy was smaller and the weight percent of < 10-~m-diameter particles was larger (12% versus 1% for the VA and 

USGA-Ar powders, see Figure 7) than in any other alloy. The large surface-tovolume ratio in the USGA-He powder may have fixed larger amounts of hydrogenand oxygen-bear1ng species than for other alloy powders prior to consolidat1on, leading in turn to larger concentrations of Hand 0 after degassing. 

Oxygen contents in the extrusions were generally larger than those in the 

powders, suggest1ng that some chemical fixing of oxygen occurred during degassing. Oxygen levels were comparable to those in other RSP Al-Li alloys produced at MDRL (21), and considerably lower than the typical value of 4000 ppm 

found in commercial 7091-Al, wh1ch is consolidated from air-atomized powder. 

5.2.2 Microstructure 

Optical micrographs of as-rece1ved Al-4.4Cu-l.5Mg-IFe-1Ni-0.2Zr and Al3L1-1.5Cu-1Mg-0.5Co-O.2Zr extrusions and forgings (Figures 23-25) show microstructural refinement with respect to 11M alloys, unrecrystallized grain 

structures, and microstructural features not exceeding 10 ~m in diameter. 

Differences among m1crostructures of as-extruded non-Li-containing alloys 

(F1gure 23) parallel differences in composition. The 1- to 3-~m-diameter 

particles commonly observed 1n the AI-4.4Cu-l.5Mg-IFe-1Ni-0.2Zr VA extrusion, 

and occasionally observed in the USGA extrusions, are larger than the Al9FeNi 

d1spersoids (compare with Figure 5), and therefore must be Al2Cu or Al2CuMg 

const1tuent particles. The phase in equilibrium with the matrix in AI-3.94Cu1.54Mg, the analyzed compos1tion of the VA extrusion, is Al 2CuMg (17). This 

phase 1S known to be strongly etched by Keller's reagent in 2XXX series alloys 

(22); thus the 1- to 3-~m constituent particles are identified as Al2CuMg. As 

the Mg concentration is reduced, as in the USGA extrusions and 1n the Al4.4Cu-l.5Mg-2Fe-2Ni-O.2Zr extrusion, Al2CuMg is predicted by the phase diagram 

(17) to be replaced by Al2Cu, wh1ch is not strongly attacked by the etch 

39 
ST 

l 

(a) 

1--------- L T ---~'_w4 

10 flm 

10 flm 

l'igllre 23. Optic:al of 

IFe··]Ni-O.2ZI' made from 

powder, (d) 1'''''.,''11 .... ''' 

10 flm 

10 flm 

(;1'41·0922·38 

extnlsiol1s: (aHc) 

lJSGA-AI' (c) lUSGA-He 

40 
(a) (b) 

I--'---L T---------I 

~ 

10 fJ.ID 

10 fJ.ID 

(d) 

L~ 

10 fJ.ID 

~ 

10 fJ.ID 

GI'41'()922-39 

Figure 24.0ptic:al microl~raphs of AI-3Li··1.5Cu-lMg-O.5Co-O.2Zr extrusions made from (a) VA 

powdler, (b) USGA··Ar powder, (c) lJSGA-He powder, and (d) RQ flakes. 

41 
IO,,l,m 

10 f,l,m 

(;1'41·0922·40 
medlum and appears as smaller, outlined light particles (22). These predlct10ns are consistent with the appearance of Figures Z3b-d in comparison with 

23a. F1nally, the Al9FeNi dlsperso1ds are not strongly attacked by the etch 

medium and appear as numerous flne, Ilght partlcles ln all the micrographs in 

Flgure 23. The dispersoid density lS noticeably greater in Figure 23d, conslstent with the larger volume fraction of dispersoids in the 2Fe-2Ni alloy. 

Major mlcrostructural features of as-extruded AI-3Li-l.5Cu-lMg-O.5CoO.2Zr are 3- to 8-um-diameter, Jagged constituent particles (Figure 24). The 

phases in equilibrium with the matrix for the overall composition Al-3Li-l.5Cu 

are 6 (AILi) and T2 (AI5Li3Cu). Prior experience at MDRL (20) has shown that 

6 is not strongly attacked by the Keller's reagent used to etch the metallographic samples, so the dark constituent particles here are probably TZ' The 

USGA extruslons contain 2.7Li and O.7Mg and should therefore have smaller 

volume fractions of AI-Li-Cu-Mg constituent phases than the VA or RQ extrusions, which contain 3Li and IMg. Comparing Figures 24b and 24c with 24a and 

Z4d shows this to be the case. The smaller particles, which are not strongly 

etched, may be A1 9C02 dispersolds, or in the case of dispersoids along prior 

particulate boundaries (e.g., longitudlnal sections of Figures Z4b and Z4c), 

may be oxides. 

Optical micrographs of forgings (Figure 25) resemble those of the 

corresponding VA extrusions except for the weaker texture. Distributions of 

constituent particles reflect prior segregation within individual particles. 

The non-Li-contalning forging shows particle-to-particle variation of concentratlon of fine, lightly etched dispersoids, probably Al 9FeNi (Figures 25a 

and Z5b). 

5.2.3 Texture 

Flgures 26-28 show {Ill} pole f1gures of extrusions and forgings. Pole 

f1gures of extruslons are typical of a hot-worked material which 1S fully 

recovered but not recrystalllzed. The forgings have a weaker fiber texture 

typical of axisymmetric deformation. 

The pole figure of the Al-4.4Cu-l.5Mg-IFe-lNi-O.2Zr VA extrusion 

(Figure Z6a) is also typical of the USGA extrusions. As the subgrain size 

decreases and the subgraln density increases ln a fully-recovered microstructure of this type, the variation in subgrain orientation from those 

43 
<a) 

PH I A AL ALLOY (I I I ) E D 

CONTOUR VALUES 

I 4 00 

2 8 00 

3 16 00 

4 24 00 "" 5 32 00 

6 40 00 

7 46 00 

6 56 00 I I 

/ 

i I, 0 } tD \\ ;/ ~"-"" I ____ ""/ 

~~ 

(e) 

"-

\ 

(b) 

PMIAST AI... ALLOt cllt) ED. 

ccwc JP I, AL'JES 

(d) 

3 2 

....L.._ ----

1T D 

/ 

/ ;' 

/ // 

/ 

/' 

GP41~-41 

Figure 26. (111) pole figures of VA extrusions: (a), (b) AI-4.4Cu-l.5Mg-IFe-lNi-O.2Zr, (a) asextruded, (b) solution-treated 530°C/l b; (e), (d) AI-4.4Cu-1.5Mg-2Fe-2Ni, (c) asextruded, (d) solution-treated 530°C/l b. 

44 
Pt'':., AL ALLOY 

SO" 'O'JR IftL:.JES 

I 3 00 

2 60:'1/ 

3 9 n~ 

4 I? 00 /' 

// 1/ 

/ / / 

I 

! 

\ 

\ 

\ 

(: II ) 

(a) 

E.D. 

-----.... -- "'----L...-.--

\ 

/ 

\ 

\ 

\ 

JT D. 

\ 

'\ 

\ 

" <3 

, 

8 

(b) 

,.-L-

_'l~ --

~---_J ___ ~- -

(d) 

, ) ~~ 

-- -- ... - .... "" .3 ~~ ,-

---- , '" "n , "- ?~ ,~ " "13 (:'~ \ 

33 nn , 

, 

:--__ 1 

\ ~ / 

........... ~ .... - . ----... --\' " - __ .... _ ' ..... __ --- ® ,.r-.... ..,---- ... .J"- ______ r-" - _ ..... 1 

--'--~--...--

\ 

" I 

/ 

" 

, 

\ 

lTD 

I 

·TD 

GP41..,U-41 

Figure 27. (111) pole figures of AI-3Li-1.5Cu-lMg-O.5Co-O.2Zr extrusions: (a), (b) V A extrusion, 

(a) as-extruded, (b) solution-treated 530°CIl h; (c), (d) USGA-He extrusion, (c) asextruded, (d) solution-treated 530°CIl h. 

45 
(a) 

E.D 

PH204VAF AS FORGED AL (111) CONTOUR VALUES _ - - 1 1 1313 /' 

2 31313./ 

3 5 1313 

4 7 1313 

5 8 1313 

6 91313 

7 Ie 1313 

8 11 1313 

\ 

-... "-

"-

"-

\ 

\ 

\ 

I 

I 

/ 

/ \ / 

'-'" 

~/ 

- - ---

(c) 

E.D. 

PH27VAF AS FORGED AL (111) 

CONTOUR VALUES""-

1 2 00' ".,. - - - -... 

2 31313./ "'- 3 41313,.. "- 4 S00 

5 61313 "- 6 700 \ 

7 600 

\ 

\ / 

\ 

\ 

/ 

/ 

/ 

/ 

........ '--._---..,., 

./ 

TD 

TD. 

(b) 

ED. 

PM24FST AL ALLOY (111) 

CONTOUR VALUES 

1 2 1313 

2 4 1313 ./ 

3 6 1313 

4 6 1313 

5 9 1313 

6 113 1313 

7 11 1313 

8 121313 

I 

I 

I 

\ 

\ 

\ 

...,...------.. "-

/ 

./ "-

-------~ 

(d) 

ED 

PH27FST AL ALLOY (111) 

CONTOUR VALUES 

1 1 1313 --- - - 2 2 1313 /' 

3 3 013 ./ 

4 4 1313 

5 5 1313 

6 6 013 

7 7 1313 

8 8 1313 

I 

I 

\ 

\ 

"-

"- "- ...... 3 

--- "-

"-

./ 

----

/ 

/ 

I 

/ 

Figure 28. (111) pole figures of forgmgs: (a), (b) AI-4.4Cu-1.5Mg-lFe-1Ni-O.2Zr, (a) as-forged, 

(b) solution-treated 530°C/l h; (c), (d) AI-3Li-1.5Cu-1Mg-O.5Co-O.2Zr, (c) as-forged, 

(d) solution-treated 530°C/l h. 

46 

TD 

TD 

GP41~11-4J 
or1entations typical of the extrusion increases, and the maximum pole dens1ty 

or X-ray 1ntensity decreases. The texture of the AI-4.4Cu-1.5Mg-2Fe-2N1-0.2Zr 

alloy (Figure 26c) is relatively diffuse, suggesting that the larger dens1ty 

of Al9FeN1 d1sperso1ds stabilizes a finer substructure than 1n the 1Fe-1N1 

extrus10ns. 

Pole f1gures of the Al-3L1-1.5Cu-lMg-0.5Co-0.2Zr extrusions (Figure 27) 

resemble those in Figure 26, but the maximum 1ntensity is smaller, indicating 

a smaller subgrain size and a larger variation in subgrain orientat10n. The 

general appearance of the pole figures is also typ1cal of USGA-Ar and RQ extensions, showing that the as-extruded subgrain microstructure is influenced 

primarily by extrusion parameters and not by particulate type. 

Pole f1gures of the forgings (Figure 28) all show weak textures result1ng 

from the relatively small deformations required to produce the forgings. The 

texture 1n the Al-4.4Cu-1.5Mg-1Fe-1Ni-0.2Zr forging is somewhat sharper than 

that in the AI-3Li-1.5Cu-1Mg-0.5Co-0.2Zr forging. 

5.3 Heat Treatment of Extrusions and Forgings 

Aging curves of all extrusions, solution treated and aged at various 

temperatures, are shown in Figures 29-31. Two solution-treatment temperatures 

were chosen for each alloy: the estimated highest safe temperature, and a 

temperature 30°C lower. Hardness curves were obtained for 1900C aging of the 

non-Li-containing alloys and 1770C aging of the Li-contain1ng alloy. The 

ag1ng behavior of forgings was assumed to resemble that of VA extruS10ns. 

The non-Li-containing VA extrusions attain maximum hardness in 8 to 16 h 

at 1900C as expected, followed by a sharp decrease (Figure 29). Increasing 

the solution-treatment temperature of Al-4.4Cu-1.5Mg-1Fe-1Ni-0.2Zr from 500 to 

5300C solution1zes larger amounts of Cu and Mg, permitting a larger density of 

strengthening precipitates and 1ncreasing the hardness. The effect of 

1ncreasing the dispersoid-forming element addition from 1Fe-1Ni to 2Fe-2Ni 1S 

to increase the hardness by 6 to 7 Rockwell B (RB) points. Aging curves of 

USGA extrusions are similar to those in Figure 29, but peak hardnesses are 72 

to 74 RB, compared with 81 RB for the VA extrusion, probably as a result of 

the lower Mg content in the VA extrusions and the resulting lower volume fract10n of AI-Cu-(Mg) precipitates. No variations 1n hardness or aging behavior 

traceable to particulate type were noted. 

47 
90 

80 

/ 

'" 70 I II 

'" ~/I (1) 

c 

-0 

..... 

/ ~ 

..c 

CO I 

"0 

~ 

..>0: 

u 

0 50 IX 

40 

30 

01 02 05 2 5 10 20 50 100 

Agmg tIme (h) 

(.P41-092244 

FIgure 29. Agmg behavIor at 190°C of non-LI-contammg V A extrusIOns: ( 0,0) 

AI-4.4Cu-1.5Mg-IFe-lNi-O.2Zr; (0) solutIOn-treated 530°CIl h, (0) solution-treated 

500°C/I h, (/!;.) AI-4.4Cu-1.5Mg-2Fe-2NI-O.2 Zr, solution-treated 530°C/I h. 

Figures 30 and 31 show typical aging behavior of the Al-3Li-1.5Cu-1Mg0.5Co-0.2Zr extrusions. Hardness curves of samples previously solution 

treated at 530 and 560°C were virtually identical for each particulate type. 

Peak hardness was uniformly reached 1n 8 to 16 h, and little variation in 

hardness was noted in the vicinity of the maximum. The peak-aged hardness of 

VA and RQ extrusions, which contained 3Li-1Mg, was 87 to 88 RB (Figure 30); 

while those of USGA extrusions, which conta1ned 2.7L1-0.6Mg, was 83-85 RB 

(F1gure 31). These results suggest that the ultimate tensile strengths of the 

VA and RQ extrusions as peak aged should be somewhat higher than those of USGA 

extrusions. 

48 
90 

80 

'" 70 II) '" 

c:: 

"0 

... 

C<l 

..c:: 

"0 

a:l -

~ 

~ 

~ 

u 

0 50 cG 

40 

30 

01 02 05 2 5 IO 20 50 100 

Agmg time (h) (,1'41-0921-45 

Figure 30. Agmg behavior at 177°C of AI-3LI-1.5Cu-lMg-0.5Co-0.2Zr VA extrusIOns, solutIOntreated at ( 0) 530°CIl hand ( 0) 560°CIl h. 

90~""'~""''''''~'''''~'''''~''''''''''~''''''''''r-'''''~''''''''''~'''''~ 

80 

I " 

I '" cr 70 

I I cr I II) I, c:: 

"0 ~ ... 

C<l t ..c:: I 

a:l 60 

"0 /1 

~ I 

~ 

u 

0 50 cG 

40 

30 

o I 02 05 2 5 10 20 50 100 

Agmg time (h) 

(.1'41092246 

Figure 31. Agmg behavior at 177°C of AI-3LI-1.5Cu-lMg-0.5Co-0.2Zr USGA-He extrusIOns: (0) 

solution-treated 530°CIl hand (D) solutIOn-treated 560°C/l h. 

49 
6. TASK 3: MICROSTRUCTURE AND PROPERTIES OF CONSOLIDATED FORMS 

6.1 Microstructure 

6.1.1 SolutLon-Treated CondLtLon 

OptLcal micrographs of non-LL-contaLnLng alloy extrusLons and forgLngs, 

solutLon treated at 500 and 5300C, are shown Ln FLgures 32-34. In accord wLth 

the Al-Cu-Mg phase dLagram (17), the A1 2CuMg and A1 2Cu constLtuent part1cles 

v1sible in the extrusions (FLgure 23) have dissolved, except Ln the Al-4.4Cu1.5Mg-IFe-INi-0.2Zr extrusion solut1on treated at 5000C (Figure 32b). The extrusions have not recrystallized, and show a pronounced "pancake" graLn structure. The A1 9FeNi dLspersOLds are present in greater density Ln the 2Fe-2NL 

alloy (Figure 32c) than in the IFe-INL alloy. Minor microstructural dLfferences are attributable to composLtional variations rather than dLfferences Ln 

particulate type. The forging (FLgure 34a) shows evidence of prLor particulate boundaries and an unrecrystallized graLn structure, as 1n the as-receLved 

condit10n. 

Optical m1crographs of Li-containing alloy extrusions and forg1ngs, 

solutLon treated at 530 and 560oC, are shown in Figures 34-36. The Al-L1-Cu 

phase dLagram (17) predLcts that 2.7Li-I.SCu (the actual compos1tLon of USGA 

extrusions) should be almost completely dLssolved at 5300C, whLle 3LL-I.5Cu 

(the actual composition of VA and RQ extrusions) should be dissolved at 560oC. 

The appearance and volume fractions of Al-LL-CU-(Mg) constLtuent particles 1n 

Figures 34-36 are in accord with these predict10ns and the chem1cal analyses 

(Table 3). The concentrat1ons of Li, Cu, and Mg in sol1d solut10n, and hence 

avaLlable for precipitation, are greater for the VA and RQ extrus10ns than for 

USGA extrusions (Section 4.3.1), particularly after 5600 C solutLon treatment. 

Thus, it is expected that the former extrusions w1ll be stronger and have a 

greater var1ation of strength w1th prior solution-treatment temperature than 

the latter. 

Also vLsible Ln the extrusions are f1ne, lLghtly etched partLcles on 

(highly distorted) prior partLculate boundarLes, most of whLch must be oXLdes. 

They are less numerous in the RQ extrusion (Figures 3Sd and 36d) than Ln the 

other extrusLons, Ln accord w1th the smaller part1culate surface area and 

lower oxygen content of thLs extrusion (Table 4). Some of the lLghtly etched 

50 
(1:11) (b) 

f----LT ----1 L_~ . 10 /lm 

(c) 

L-J 

10/lm 

Figure :32. Optical micrographs of solution-treated extrusions: (a), 

(b) AI-4.4Cu-1.5Mg-llFe-lNi-O.2Zr, (a) 530°C/l h, (b) 500°C/l h; (c) 

AI-4.4Cu-l.5Mg-2Fe-2Ni-O.2Zr, 530°CI1 h. 

51 

GP41-0922-47 
10 ~Hn 

33, 
(a) 

L....J 

lO ""m 

(b) (c) 

L-J 

10 ""m GP41·0922·49 

Figure 3i4. Optical micrographs of forgings; solution-annealed 530°C/lh: (a) AI.4.4Cu-1.5Mg-IFeINi-O.2:Zr, perpendicular to forgings direction, (b), (c) AI-3Li-l.5Cu-lMg-O.5Co-O.2Zr, 

(b) perll)endiculllr to forging direction lind (c) parallel to forging direction. 

53 
ST 

1 

(a) 

~---

10 fJ-m 

(c) 

L_I 

IOfJ-m 

(b) 

L-l 

10 fJ-m 

(d) 

L-I 

10 fJ-m 

GI'41-0922-50 

Figure 35. Optical micrographs of AI-3U-1.5Cu-lMg-O.SCo··O.2Zr extrusions, solution-treated 

530°C/I h, prepared from (a) vacuum-atomized powder, (b) argon ultrasonically 

atomized powder, (c) helium ultrasonically atomized powder, and 

(d) roller-quenched flakes. 

54 
ST 

L 

(a) 

~-LT.--~ 

L--.J 

10 fJ..m 

(c) 

L--.-J 

10 fJ..m 

(b) 

L-l 

10 fJ..m 

(d) 

L-l 

10 fJ..m GP41·0922·51 

Figure 36. Optical micrographs of AI-3U-1.5Cu-lMg-O.5Co-O.2Zr extrusions, solution-treated 

5600C/l h, prepared from (a) vacuum-atomized powder, (b) argon ultrasonically 

atomized powder, (c) helium uldrasonically atomized powder, and (d) roller-quenched 

flakes. 

55 
particles at prior particulate boundar1es may be Al 9Co 2 d1sperso1ds. The L1-

containing forging (F1gure 34b) shows const1tuent particles and oX1des s1m1lar 

to those in the extrusions. Former powder particle boundaries are outl1ned by 

oxides as a result of the smaller degree of deformat1on in the forging than 1n 

the extrusions. 

Typical {Ill} pole figures of solution-treated extrusions and forg1ngs 

are shown in Figures 26-28. None of the consolidated forms has recrystallized 

after solution treatment, consistent with the microstructural observat1ons, 

and the pole figures are still typical of the extrus10n or forging processes. 

The textures have sharpened noticeably compared with those in the consolidated 

forms, indicating subgrain coalescence and motion of high-energy subboundaries 

during solution-treatment. Texture sharpening is particularly noticeable in 

the AI-4.4Cu-l.5Mg-lFe-lNi-0.2Zr extrusion (compare F1gures 26a and 26b), and 

in the Al-3Li-l.5Cu-lMg-0.5Co-0.2Zr USGA-He extrusion (compare Figures 27c and 

27d). 

6.1.2 Aged Cond1tion 

TEM micrographs of variously aged extrusious and forgings are shown 1n 

Figures 37-40. The non-Li-containing alloy extrusions (Figures 37-39) have 

modified mosaic structures consisting of 1.8- to 2.5- x 3- to 4-~m, unrecrystallized subgrains. The Al 9FeNi dispersoids are fairly homogeneously dispersed and typically 0.4- to 0.5-~m in diameter. Subgrain boundaries are 

decorated with incoherent e (AI 2Cu) precipitates in artificially aged samples. 

All extrusion microstructures are similar, regardless of particulate type 

(compare Figures 37a-c and Figures 38a-b), further conf1rm1ng that microstructures of aged alloys are not dependent on the microstructures of the or1g1nal 

particulates. 

Micrographs of non-Li-containing extrusions in the T6 temper (solution 

treated at 530oC/1 h + aged 190oC/8 h, Figure 37) show coarse (1 ~m in 

length) 9' precipitates, owing to the small number of nucleation sites for 

this semi-coherent phase. The microstructure of the 2Fe-2Ni alloy 

(Figure 37d) is similar to that of the lFe-lN1 alloy (F1gures 37a-c), the 

added Fe+Ni being accommodated by an increase 1n the density of A1 9FeNi 

dispersoids. 

56 
(a) (b) 

L--.J 

If.-tm 

(c) (d) 

GI'4I-0922-52 

Figme 37. TEM micrographs of non-U containing extrusions, T6 temper (solution-treated 

530°C/1 h + aged 190°C/8 h): (a) - (c) AI-4.4Cu-1.5Mg-IFe-lNi-O.2Zr from (a) 

vacuum atomized powder, (b) argon ultrasonically atomized powder, (c) helium 

uitmsonicallly atomized powder; (d) AI-4.4Cu-l.5Mg-2Fe-2Ni-O.2Zr from vacuum 

atomiz4~d powder. 

57 
lJhffi 

IJhffi 

GI'4I·0922·53 
(a) (b) 

L_~ 

0.2 p.,m 0.2 p.,m 

(c) (d) 

L __ -, 

0.2 p.,m 0.2 p.,m 

(;1'41·0922·54 

Figure 39. TEM miclI'Ographs of AI-4.4CUI-1.SMg-U'e-lNi-0.2Zr V A extrusions, T851 temper, 

exposled for 100 h at (a) 25°C, (b) 150°C, (c) 204°C, and (d) 260°C. 

59 
Mlcrostructures of non-Ll-contalnlng extrusl0ns and forglngs 10 the T851 

(Solutl0n treated at 530oC/1 h + 2% stretch + 190oC/8 h) and T351 (Solut1on 

treated at 5300 C/1 h + 2% stretch + 250 C/168 h) tempers are shown In 

Flgure 38. GUlnler-Preston (GPl) zones ln the naturally aged alloy are too 

small to be imaged In Figure 38a. The effect of losertlng a 2% stretch between solution treatment and aglng 18 to reflne and lncrea8e the denslty and 

homogenelty of e~ precipltates (compare Flgures 38b-c wlth 37a-b, respectlvely). The increase ln preclpitate denslty lS expected to produce a 

substantlal lncrease In strength In the T85l temper, compared wlth T6. The 

forglng mlcrostructure (Flgure 38d) lS simllar to that of the extruSl0ns 

except for the unlform 4- to 5-~m subgrain dlameter, compared wlth 2- to 3-~m 

In the extrusions. 

High-magnlfication micrographs of Al-4.4Cu-l.5Mg-lFe-lNl-0.2Zr VA 

extrusions, T85l temper, exposed to 150, 204, and 2600 C for 100 h, are shown 

in Figure 39. Exposure to l500 C/100 h (Flgure 39b) has llttle effect on 

preclpitate slze, but promotes growth of subgrain-boundary e and formatlon of 

a PFZ - 0.2 ~m wide. After 100 h at 2040 C, e plates (thicker plates ln 

Figure 39c) have begun to form at the expense of e~. As the temperature lS 

raised, these plates grow faster wlth concomitant dlSSolutl0n of e~, until 

after 100 h at 2600C only coarse e plates are vlsible. Mlcrostructural 

changes with increasing exposure temperature suggest that noticeable decreases 

in strength should be apparent at 204°C, and that the preclpitatlon 

strengthening contribution should be absent after 260°C exposure. Mlcrostructural development upon aging and high-temperature exposure ln the non-Llcontalning alloys resembles that In l/M 2024-Al. Thus, any lncrease III 

strength of the RSP alloys, compared wlth thelr l/M counterparts, lS the 

result of Hall-Petch strengthenlng from the flne subgraln Slze or dlsperslon 

strengthenlng from the Al 9FeNl dlspersoids. 

General microstructural features of peak-aged Al-3Ll-l.5Cu-1Mg-0.5Co0.2Zr are revealed ln Figure 40a. The structure lS unrecrystalllzed, wlth 2-

to 4-~m diameter subgrains. Numerous low-angle and several hlgh-angle boundaries among subgrains are observed, as well as a sparse d~strlbutlon of 0.3-

to 0.4-~m diameter Al9Co2 dispersolds. Rows of flne, well-s~parated oXlde 

particles outline th~ prior partlcle boundarles. Subgrain boundarles sometimes coinclde with the oxides, but they have more often migrated from the 

60 
(a) (b) 

0.2 p,m 

(c) (d) 

L-.. __ , • ..J 

2 p,m 2p,m 

(c~) (f) 

2 p,m 

(;\'41-0922-55 

Figure 40. 1'EM micrographs of AI-3Li-1.5Cu-lMg-0.5Co-0.2Zr extrusions: (a) " (c) from vacuum 

atomized powder, (a), (b) 1'6-530 temper (solution-treated 530°C/l h + aged 177°C/8 

h), (c) 1'6-560 temper (solution-treated 560°C/l h + aged 177°C/8 h); (d) from argon 

ultrasonically atomized powder, 1'6-530 temper; (e) from helium ultrasonically 

atomized powder, 1'6 .. 560 temper; (0 from roller-quenched flakes, 1'6-530 temper. 

61 
oX1des during extrus10n, golut10n treatment, or both. Rows of oX1des are less 

f~equently observed 1n the non-L1-conta1n1ng alloy. At h1gher magnificat10n 

(Figure 40b), a homogeneous d1str1bution of 15- to 20-nm diameter o~ 

prec1p1tates is observed. 

M1crostructures of USGA extrusions are sim1lar to that of the VA 

extrus10n (compare F1gures 40d and 40e with 40a). The RQ extrusion 

(F1gure 40f) has a larger average subgrain size, a more homogeneous d1str1bution of Al9Co2 disperg01ds, and fewer oxides (owing to the greater thickness 

of the or1ginal part1culates) than the powder extrusions. Th1s extrusion is 

the only one 1n Wh1Ch differences in particulate type produce sign1ficant 

d1fferences in the peak-aged m1crostructure. 

Upon solution treatment at 560oC, the Al-3Li-l.5Cu-lMg-O.5Co-O.2Zr 

extrusion exhibits equiaxed subgrains rather than the mosaic structure 

observed after 5300 C solution treatment (compare Figure 40c with Figure 40a). 

Since both structures have subgrains of similar diameter, microstructural 

modification has probably occurred by subgrain boundary motion at the higher 

solution-treatment temperature, rather than by recrystallization. The subtle 

change 1n microstructure is unlikely to significantly affect mechanical 

properties. 

6.2 Properties 

6.2.1 Density 

Densities of extrusions are listed in Table 7. Increases in density of 

the non-Li-containing alloys relative to 2024-Al are in good agreement with 

predict10ns based on density contributions from the Al 9FeNi dispersoids. 

Density decreases of the L1-containing alloys relative to that of 7075-Al are 

in accord with those expected from alloys containing 2.7-3Li. Densities of 

Li-containing USGA extrusions are slightly greater than those of VA or RQ 

extrusions, in accord with their lower Li and Mg concentrations. 

62 
TABLE 7. DENSITIES OF RSP ALLOY EXTRUSIONS. 

Alloy and 

partIculate type 

AI-4.4Lu-l.SMg-lFe-lNI-O.2Zr 

Vacuum atomized powder 

Argon ultrasonically atomized powder 

Helium ultrasonically atomized powder 

AI-4.4Cu -1.SMg-2Fe-2N i-O.2Zr 

Vacuum atomized powder 

2024-AI (11M) 

AI-3Li-l SCu-lMg-O.SCo-O.2Zr 

Vacuum atomized powder 

Argon ultrasonically atomized powder 

Helium ultrasonically atomized powder 

Roller-quenched flakes 

707S-AI (11M) 

6.2.2 Elastic Modulus 

Density 

(g/cm3) 

2806 

2813 

2823 

2861 

277 

2489 

2530 

2530 

2492 

280 

070 increase 

relative to 

2024-AI 

1 3 

1 6 

1 9 

3 3 

070 decrease 

relatIve to 

707S-AI 

11 1 

96 

96 

110 

GP41 09227 

Elast1c moduli of heat-treated extrusions are given in Table 8. The 

relatively small volume fractions of incoherent d1sperso1ds in the non-Licontain1ng alloy increase its elastic modulus by only 2 to 3% relative to that 

TABLE 8. ELASTIC MODULI OF RSP ALUMINUM ALLOY EXTRUSIONS. 

Elastic 070 increase 070 increase 

Alloy composition and Heat modulus, relative to relative to 

particulate type treatment (GPa(loti psi)) 2024-AI 707S-AI 

AI-4.4Cu-l.SMg-lFe-lNi-O.2Zr 

Vacuum atomIzed powder T851 739(10.72) 2.1 

Argon Ultrasonically atomIzed powder T851 747 (10 85) 33 

Helium ultrasonically atomIzed powder T851 737 (10 68) 1 8 

2024-AI(I/M) 72 4 (10 50) 

AI-3Li-l.SCu-lMg-O.SCo-O.2Zr 

Vacuum atomIzed powder T61530 78.5 (II 39) 106 

Argon ultrasonIcally atomIzed powder T61530 760 (II 02) 70 

HelIum ultrasonIcally atomIzed powder T61560 763 (II 06) 75 

Roller-quenched flakes T61530 80 1 (11 62) 128 

707S-AI (11M) 71 0 (10 30) 

GP41 09228 

63 
of 2024-Al. The percent modulus increase in Al-Li alloys, compared with 11M 

7075-Al, has been found in previous work to be approximately equal to the 

percent density decrease, and this empirical rule is followed for the data 

reported here. The moduli of Al-3Li-1.5Cu-lMg-O.5Co-O.2Zr VA and RQ extrusions are 17 to 20% larger than that of pure aluminum [67 GPa [9.7 x 106 psi] 

(22)], in accOird with thE! rule-of-mixtures calculation that the modulus in 

Al-Li alloys should increase by about 6% for every wt% Li added. 

6.2.3 Tensile Mechanical Properties 

6.2.3.1 Non-Li-Containing Alloys. 

Mechanical propertiE!s of non-Li-containing alloy extrusions and forgings 

at temperatures between 25 and 2600C are shown in Figures 41-48. Comparisons 

between results for the RSP alloys and those for commercial 2x.xX alloys such 

as 2024, 2124~ and 22.19 (Figures 46-48) illustrate the improvements in properties obtained by RSP and addition of fine, incoherent dispersoids to an 

AI-4. 4Gu-1. 5Mg mat rh:. • 

408 393 

Yield stress 

(MPa) 

504 

482 476 

Ultimate tensile stress 

(MPa) 

AI-4.4Cu-l.5Mg-l Fe-I Ni-0.2Zr 

Vacuum atomized powder 

Helium ultrasonically atomized powder 

AI-4.4Cu-l.5Mg-2Fe-2Ni-0.2Zr 

Vacuum atomized powder 

2024-T6 

10 

Elongation 

(070 ) 

GP4I·()922·56 

Figure 41. Ambiell1lt temperature tensile properties of non-Li-containing alloys: T6 temper 

(solution tn~ated 530°C/1 h -I- aged 190°C/8 11). 

64 
Yield stress (MPa) Ultimate tensile stress 

(MPa) 

AI-4.4Cu-I.5Mg-IFe-INi-O.2Zr 

V A (solution treated 530°C/l h) 

V A (solution treated 500°CIl h) 

USGA-Ar 

USGA-He 

AI-4.4Cu-l.5Mg-2Fe-2Ni-O.2Zr 

VA (solution treated 530°C/l h) 

2024-T851 

13.1 

7 

Elongation (Olo) 01'41·0922·57 

Figure 42. AJllllbient temperature tensile l)Il'operties of lIlolll-Li-containing alloy extrusions: T851 

temper (sohlltion treated 530°1[11 h [except as shown] + 21170 stretch + 190°C/8 h). 

439440 448 

Yield stress (MPa) Ultimate tensile stress (MPa) 

AI-4.4Cu-l.5Mg-l Fe-l Ni-0.2Zr 

Extrusion, L 

Extrusion, T 

Forging, L 

AI-4.4Cu- I .SMg-2Fe-2Ni-O.2Zr 

Extrusion, L 

Extrusion, T 

2024-T85J 

10.0 

7 

Elongation (Olo) 

OP41·0922·5H 

Figllue 43. Ambient temperature tensile properties of non-U containing alloy extrusions and 

forgings: TH51 temper. 

65 
Yield stress (MPa) Ultimate tensile stress 

(MPa) 

AI-4.4Cu-I.5Mg-l Fe-I Ni-O.2Zr 

Extrusion ~ 

Forging ~ 

AI-4.4Cu-I.5Mg-2Fe-2Ni-O.2Zr 

Extrusion 

2024-T35I 

19 

Elongation (070) 

(,1'41-0922-59 

Figure 44. Ambi(mt temp(~rature tensile properties of non-U -containing alloy V A extrusions and 

forgings: T351 temper (solution treated 530°CI1 h + 2070 stretch + aged 25°C/168 11). 

440 

417 

Yield stress (MPa) Ultimate tensile stress (MPa) 

AI-4.4Cu-l.5Mg-lFe-1Ni-0.2Zr 

V A extrusion 

VA forging 

USGA-Ar extrusion 

USGA-He extrusion 

AI-4.4Cu-1.5Mg-2Fe-2Ni-0.2Zr 

V A extrusion 

2124-T851 

10 

Elongation (070) 

(ZJ 

rm 

fill 

f&J 

GP41-0922-60 

Figure 4:5. Tensile properties of non-Li-containing alloy extrusions and forgings: T851 temper, at 

150°C following 100 11 exposure at 150°C. 

66 
5OO~--------~------~--------~------~--------~~ 

400 ---------

...... 

tIS 

~ 300 ::E --------- '-' 

'" '" CI) 

... 

... 

'" "0 200 U 

>= 

100 

o~--------~------~------~------~--------~~ o 50 100 150 200 250 

Test temperature (OC) GP41 092261 

Figure 46. Variation with temperature of the yield stress after 100 h exposure at temperature of 

( 0) AI-4.4Cu-l.5Mg-IFe-1Ni-O.2Zr V A extrusion, (0) AI-4.4Cu-l.5Mg-IFe-1Ni-O.2Zr 

forging, ( - - - ) 2124-AI plate, and (-- -) 2219-AI(23). Original temper T851. 

600 

500 ------ ...... ---- tIS 

~ 

::E 

'-' 400 

'" '" CI) 

... 

... 

'" 

~ 300 en 

c:: 

CI) 

... 

CI) 

... 

tIS 

e 200 

.;:: 

5 

100 

50 150 200 250 

Test temperature (0C) GP41-0922-62 

Figure 47. Variation with temperature of the ultimate tensile stress after 100 h exposure at 

temperature of (0) AI-4.4Cu-1.5Mg-IFe-1Ni-O.2Zr V A extrusion, (0) 

AI-4.4Cu-l.5Mg-IFe-1Ni-O.2Zr VA forging, ( - - - ) 2124-AI plate, and (---) 

2219-AI(23). Original temper T851. 

67 
25r----r------~------~------~--~--~ 

20 

/ ,.-. 

~ '-" '" Q) / ... '" 

-= 

'-

; 15 ~ 

.9 ...- ~ 

I: 

9 

... 10 (Ij 

tlI) 

I: 

0 --------- ---- m 

5 

150 250 

Test temperature (0C) GP41-0922-63 

FIgure 48. Variation with temperature of the ductilIty after 100 h exposure at temperature of 

(0) AI-4.4Cu-l.SMg-IFe-1Ni-O.2Zr VA extrusion, (0) AI-4.4Cu-1.SMg-lFe-1Ni-O.2Zr 

V A forging, ( - - - ) 2124-AI plate, and (---) 2219-AI(23). Ongmal temper T8Sl. 

Results of a pre11minary study on extrus10ns 1n the T6 temper (Figure 41) 

show that properties of the VA extrusions are similar to those of 2024-Al, 

while the propert1es of the USGA-He extrusion are significantly lower, probably as a result of the smaller (Cu+Mg) concentration available for precipitation. The effect of adding 2Fe-2Ni instead of 1Fe-1N1 is to increase the 

Y1eld stress about 35 MPa and slightly reduce the duct1l1ty. 

As expected from the microstructures (compare Figure 37 with Figure 38), 

yield stresses of 1Fe-1Ni extrusions 1n the T851 temper are 70-90 MPa larger 

than in the T6 temper, with little change in tensile strength or ductility. 

Changes in the properties of the 2Fe-2Ni alloy are smaller and result in 

properties similar to those of the 1Fe-1Ni alloy in the T851 temper. Variations in properties from the T6 to the T851 tempers for the VA extrUS10ns 

strongly resemble those of 2024-Al, and no advantages result1ng specifically 

from RSP are observed. 

68 
The T8S1 properties of IFe-INi extrusions from various particulate types 

follow compos1t1onal d1fferences (Figure 42). In particular, the USGA-He 

extrusion has lower strength than the VA extrusion, 1n keeping w1th 1tS lower 

Mg level, but the strength 1S independent of the high solidification rate of 

He-atom1zed powder. These observations are conS1stent with the microstructures (Figures 37 and 38), wh1ch show little d1fference among VA and USGA 

extrusions. 

The quality of consol1dation process1ng and d1fferences in consolidat1on 

process1ng between forgings and extrus10ns are reflected 1n the mechanical 

propert1es results shown in Figure 43. Transverse yield and tensile strengths 

of IFe-INi extrusions are 9S-100% of the corresponding longitudinal properties, indicating that the extrusion processing (Section 3.3.2) was effective 

in d1sintegrating particulate oxide layers and obtaining good interparticulate 

bonding. Y1eld and tensile strengths of the VA forging are 70 to 80 MPa lower 

than in the corresponding extrusion; however, the ductilities are similar, 

suggest1ng that the lower properties of the forging result from its weaker 

texture (compare Figures 26b and 28b), rather than deficient consolidation 

processing. 

Yield and ultimate tens1le strengths of non-Li-containing alloy 

extrusions in the T3S1 (stretched and naturally aged) temper are 70 to 100 MPa 

larger than in 2024-T3S1, with reasonable ductilities (Figure 44). The nearequivalence of properties of the IFe-l~i and 2Fe-2Ni extrusions indicates that 

the source of higher strength is substructural ref1nement or texture effects, 

rather than dispersion strengthening. 

D1fferences in tensile properties among extrusions from various 

part1culate types, tested at IS0oC, generally follow differences in Mg concentrat10n (F1gure 4S), except that the penalty for a shortage of Mg is more pronounced than at 2SoC. For instance, the T8S1 yield strength of the USGA-He 

extrusion at 2SoC is 91% of that of the VA extrusion (Figure 43), but only 77% 

as large at IS00C (F1gure 4S). The larger dispersoid concentrat10n in the 

2Fe-2Ni alloy is no more important at IS00 C than at 2SoC, and the 1.IMg level 

in this extrusion (compared w1th 1.SMg in the IFe-INi extrusion) causes yield 

and tensile stresses to be 30-40 MFa lower. 

The var1ation 1n T8S1 tens1le properties of VA extrusions and forgings 

with temperature, 2S to 260oC, 1S shown 1n Figures 46-48. The RSP materials 

69 
have Y1eld strength retent10n at I500C super10r to those of 2124-Al or 2219-Al 

(96 and 100% of room temperature strength at 1500C for VA extrus10ns and forgings, respectively, compared w1th 88% for 2124-Al and 79% for 2219-Al). The 

absolute Y1eld and tens11e strengths of the IFe-IN1 extrus10n are 5-8% larger 

than those of 2124-Al at 100 to I500C, but the strength advantages van1sh 

above about 180°C, and 2124-Al 1S stronger at 2600C. Extrus10n and forg1ng 

propert1es above 200°C are similar, probably as the result of the 

d1sappearance of texture strengthen1ng as a major contr1butor to overall 

strength above ~ 175°C. Finally, ducti11t1es of the RSP mater1als are superior 

to those of 2124-Al up to about 2000C. This improvement 1S most probably 

connected with the absence in RSP materials of coarse (20-~m diameter), 

1mpur1ty-containing part1cles which can act as crack in1t1at10n sites. 

SEM micrographs of typical fracture surfaces are shown 1n Figure 49. 

Most fractures are at 450 to the tens11e axis, follow1ng the d1rect10n of 

maximum resolved shear stress (Figure 49a). Ductile tear1ng 1S the un1versal 

fracture mode, as evinced by the presence of 0.5- to 2-~m d1ameter d1mples on 

all fracture surfaces (Figures 49b-f). Transverse and 10ngitud1nal specimens 

have similar fracture surfaces, except for striat10ns on transverse specimen 

surfaces which reflect the extrusion texture (F1gures 49b-c). Occasional 

coarse inclusions on the fracture surface (Figure 49d) conta1n Ai and Cu, and 

thus are not foreign particles. The sample-to-sample cons1stency of ducti11ty 

is good, suggesting that the presence of coarse 1nclusions in some samples 

does not significantly affect mechanical properties. As the test temperature 

is raised, the average dimple size increases and plastic tear1ng becomes more 

eV1dent, but the fracture mode does not change (Figures 4ge-f). The fracture 

behav10r of the forgings is sim11ar to that of the extrus10ns. 

6.2.3.2 L1-Containing Alloy 

Properties of L1-containing extrusions are largely 1ndependent of 

part1culate type (Figure 50). T6 yield stresses of RQ extruS10ns are, 

however, 40 to 60 MPa lower than those of VA extrus10ns of s1m1lar compos1-

tion, and tenS1le stresses are 12-20 MPa lower (Figures 50 and 52). T6 

duct1l1ties of extrUS10ns are uniformly 3 to 4%, regardless of yield stress 

level. Increas1ng the solut10n-treatment temperature to 5600C to solut10n1ze 

larger amounts of Li, Cu, and Mg and provide a larger dens1ty of prec1p1tates 

70 
(a) 

1000 p.,m 

(c) 

10 p.,m 

(e) 

10 p.,m 

(b) 

(d) 

(f) 

10 p.,m 

100p.,m 

U 

Ip.,m 

GP41-0922-64 

Figure 419. SEM of AI-4.4Cu-1.5Mg-1Fe-1Ni-O.2ZI' V A extrusion fracture surfaces: 

(a) 1'851, general apl)earance; (b) 1'851, longitudinal, 25°C test; (c) 1'851, transverse, 

25°C test; (d) blrge inclusion on fracture surface (1'351, 25°C test); (e) 1'851, 

longitmlimlli, 150°C ~e§t; (f) 1'851, 204°C test. 

71 
results in a 20 to 30 MPa increase in T6 yield stress and a 30 to 50 MPa 

increase in ultimate tensile stress. T6 strengths of extrusions are generally 

equal to or greater than typical properties of 7075·-T73, but ductilities fall 

well short of the 13% typical of the 11M alloy. 

Figures SO and 51 reveal deHciencies in consolidation processing of Licontaining extrusions and forgings. T6 yield and ultimate tensile stresses of 

forgings are 70 to 90 MPa lower than those o~ VA extrusions, and the ductility 

of the forging 1s sign1fieantly lmver than that of extrusions, showing that 

interparticle bonding was not adequately achieved in the forging (Figure SO). 

Both T6 transverse strengths and ductilities are considerably lower than the 

eorresponding longitudinal properties (Figure 51), reflecting imperfeetions in 

consolidation proeessing for the E~xtrusions as well as for the forging. 

T851 strengths of extrusions are generally superior to those in the 

T6/530 temper, and approximately E!qual to strengths in the T6/560 temper (Figures 51, 52 and 53), comhi-ned with 3 to 5% ductiHty. The 2% streteh introduces defects whieh homogE~nize the precipitation of plate-like phases such as 

493 486 471 

454 

Yield stress (MPa) Ultimate tensile stress 

(MPa) 

V A extrusion 

VA forging 

USGA-Ar 

USGA-He 

RQ 

7075-T73 

13 

Elongation (070) 

GP41-0922-65 

Figure 50. Ambient temp,erahue tensile prol>erties of AI-3Li-l . .3Cu-lMg-0.5Co-0.2Zr extrusions 

ami forgings: T6/530 temper (solution treated S30°C/1 h + aged 177°C/8 h). 

72 
493 VA, longitudinal 

V A, transverse 

RQ, longitudinal 

RQ, transverse 

7075-1'73 

4.5 

13 

Yield stress (MPa) Ultimate tensile stress (MPa) Elongation (070) 

GP4 J ·0922·66 

Figure 51. Longitudinal and transverse ambient temperature properties of 

AI-3Li-1.5Cu-1Mg-O.5Co-O.2Zr extrusions: T6/530 temper. 

VA ~ 

USGA-Ar ISS 

USGA-He ~ t:1iJ 

RQ IE 

7075-T6 0 

11 

II 

Yield stress (MPa) Ultimate tensile stress 

(MPa) 

Elongation (070) 

GP41·0922·67 

Figlm~ 52. Ambient temlleratallre tensile pl'operties of AI-3Li-l.5Cu-1Mg-O.5Co-O.2Zr extrusions: 

T6/560 temper (solution treated 560°CI1 h + aged 177°C/8 h). 

73 
600 

572 573 

Yield stress (MPa) Ultimate tensile stress (MPa) 

V A, extrusion, T851/530 E21 

V A, forging, T851/530 ~ 

VA, extrusion, T851/560 (lEI 

RQ, extrusion, T851/530 rrn 

RQ, extrusion, T851/560 ~ 

4.2 

7075-T6510 

11 

5.1 

Elongation (070) (iP41·092~-6~ 

Figure 53. Ambient temperature tensile properties of AI-3U-1.5Cu-lMg-0.5Co-0.2Zr extrusions 

and forgings: TSS1 temper (solution treated 530 or 560°CIl h + 2070 stretch + aged 

177°GS h). 

Tl(AI 2CuLi). As in the case of AI-Cu-Mg alloys, the increased density of 

precipitation increases the yield stress, but does not affect the ultimate 

tensile stress. The T6/560 and T851 strengths of the Li-containing alloy are 

similar to those of 7075'-T6, but ductilities fall well short of the 11% 

characteristic of thE! l/M alloy. 

The IS00C tensile curves of the Li-containing alloys exhibit flow 

soften:ing beyond the yield point, so that only yield stresses are meaningful 

(Figure 54). The Li-containing alloy retains :its strength at ISOoC (94-100% 

of 2SoC yield strength), with a modest increase in ductility. 

74 
Yield stress (MPa) 

VA, extrusion, T6/530, 25°C ~ 

VA, extrusion, T6/530, 150°C ~ 

VA, extrusion, T8511530, 150°C ~ 

VA, forging, T6/530, 25°C [J 

VA, forging, T6/530, 150°C 0 

Elongation (070) GP41 -0922-69 

Figure 54. Tensile properties at indicated temperatures of AI-3Li-l.5Cu-1Mg-O.5Co-O.2Zr 

extrusions and forgings. 

Fracture surfaces of Al-3Li-l.5Cu-lMg-0.5Co-0.2Zr extrusion tensile 

samples show features on the scale of the subgrains (Figure 55a). At higher 

magnification, the T6/530 fracture surface shows shallow dimples (Figure 55b), 

but dimpling is lE~ss pronounced than in the non-Li-containing alloy. The 

microcrack at the right of Figure S5b is in an Al-Li-Cu-(Mg) constituent particle. Fracture of the RQ extrusion is more explicitly inter-subgranular, 

with less evidence of dimpling, and becomes strongly faceted in the highstrength T851/560 telnper (Figures SSe-d). Fracture surfaces of the forgings 

are similar to those of the extrusions. 

75 
(3) 

L..J 

100 Ilm 

10 Ilm 

(b) 

(d) 

Illm 

LJ 

10 Ilm 

GP41-0922-70 

Figme 55. Fractograpi1ts of AI-3Li-l.5Cu-lMg-O.5Co-O.2Zr extrusions: (a) V A, T6/530, 

longitudinal, 25°C test, genual appearance; (b) V A, T6/530, longitudinal, 25°C test, 

mkros(:opic appearance; (c) lRQ, T6/530, longitudinal, 25°C test; (d) RQ, T8511560, 

longitudinal, 25°C test. 

Fracture toughness values for extrusions and forgings in the T851 temper 

and TL and LT orientations are listed in Table 9. As expected, KQ values for 

the extrusions are larger in the LT than in the TL orientation, but Kq for the 

TL orientation in the forgings is, surprisingly, larger than for the geometrically equivalent LT orientation. Although KQ for the non-Li-containing extrusions and forgings does not meet the plane-strain criterion owing to insufficiently thick samples, the reported values are qualitatively consistent with a 

typical K1C fOlL 2024·-Al of 30 Mpaoml/ 2 • Toughnesses of the Li-containing 

alloys are low as a result of consolidation deficiencies revealed by the 

tensile properties data in Section 6.2.3. 

76 
TABLE 9. FRACTURE TOUGHNESSES OF EXTRUSIONS AND FORGINGS. 

Alloy 

AI-4.4Cu-l.5Mg-lFe-lNi-O.2Zr 

V A powder, T851 temper 

AI-3Li-l.5Cu-lMg-O.5Co-O.2Zr 

V A powder, T851 temper 

6.2.5 Fatigue-Crack-Growth Rate 

LT 

457 

(41 5) 

112 

(10 2) 

Extrusion Forging 

TL LT 

282 31 7 

(256) (288) 

89 106 

(8 I) (9 6) 

TL 

367 

(33 4) 

III 

(10 I) 

GP41-09229 

Fatigue-crack-growth-rate (FCGR) data for extrusions and forgings in the 

T851 temper and TL and LT orientations are shown in Figures 56-59. The 2% 

stretch required for the T851 temp~r also has the effect of relieving compressive stresses introduced during quenching, thus perm1tting a valid FCGR 

determination (23). 

FCGR results on Al-4.4Cu-1.5Mg-1Fe-1Ni-0.2Zr VA extrusions show a 50 to 

60% larger crack-growth rate for the TL than for the LT orientation 

(Figure 56). Th1s result is reasonable in view of the unrecrystallized 

"pancake" grain microstructure and strong extrusion texture. Conversely, the 

FCGR for the LT orientation in the VA forging is 35 to 60% larger than for the 

TL orientation (Figure 57). Since the forging has a symmetrical fiber texture, the LT and TL orientations should be equivalent and the FCGR results 

should be similar; thus the observed behav10r is not well understood. For 

equivalent orientations, the FCGR 1n the forging is 40 to 90% larger than that 

in the extrusion, oW1ng to the difference 1n strength between the two product 

forms. Crack-growth rates 1n the RSP extrusions and forgings are a factor of 

1.7 to 5 larger than in 2124-T851. A similar effect has been observed in comparing the FCGR of P/M 2024-T851 w1th that of l/M 2024-T851 and has been 

attributed to the decreased tortuosity of the crack path and 1ncreased accumulation of plastic zone damage at grain boundaries as the result of the f1ner 

grain size in the P/M material (24). 

77 
2 .-. , 0 

u , >. , u 10- "- 6 , e 8 , '-' 

Z , 

'0 6 , "- , <IS 

'0 4 , 

0' , .... , tU , .... 

..c: , .... 2 , ~ 

0 , .... , 01) 

I , ...10: 

u 10-7 , <IS 

.... , u 8 , 0 

::s 6 , ~ , .... , <IS 

~ 4 , 

5 10 15 202530 

Stress-mtenslty factor range, ~K (MPa rm) GP41 092271 

Figure 56. FCGR results on AI-4.4Cu-l.5Mg-lFe-lNi-O.2Zr VA extrusions, T851 temper: (0) LT 

orientation, (C) TL orientation, and ( - - - )typical results for 2124-T851 (23). 

2 .-. 

0 

u , >. 10-6 u , "- , e 8 '-' I 

Z 6 , 

'0 , "-<IS I 

'0 4 , 

0' , .... I <IS 

.... , 

..c: 2 , .... 

~ , 

0 , .... 

01) I I 

...10: 10-7 , u 

<IS , .... 

u 8 I 

0 I 

::s 6 01) I 

~ 

<IS 

~ 4 

5 10 15 202530 

Stress-mtenslty-factor range, ~K (MPa rm) GP41-092272 

Figure 57. FCGR results on AI-4.4Cu-l.5Mg-lFe-lNi-O.2Zr VA forgings, T851 temper: (0) LT 

orientation, (c) TL orientation, and ( - - - ) typical results for 2124-T851 (23). 

78 
10-6 

8 

~ II 

U 6 / ;., 

u 

"- 4 E , '-' / z 

"0 

"- 2 «S 

"0 

.... 

~ 

«S 

... 10-7 .c: 

~ 8 

0 6 ... 

00 

I 

..>0: 

u 

«S 4 ... 

u 

u 

~ 

!!!l 

~ 2 ~ 

3 5 7 10 15 20 30 

Stress-mtenslty-factor range, ~K (MPa·rm) GP41-092273 

Figure 58. FCGR results on AI-3Li-l.5Cu-lMg-O.5Co-O.2Zr V A extrusions, T8511530 temper: 

( 0 ) LT orientation, (0 ) TL orientation, ( - - - ) typical results for 7075-T651 (24), 

and (-- -) typical results for 7091-T7E69(24). 

10-6 

8 / .-. 

u 6 

U 

;., 

u 4 "-E , '-' I z 

"0 

"- 2 «S 

"0 

.... 

~ 

«S 

... 10-7 

.c: 8 .... 

~ 

0 

... 6 00 

I 

..>0: 

u 4 «S 

... 

u 

u 

~ 

!!!l 

.... 2 «S 

~ 

3 5 7 10 15 202530 

Stress-mtensity-factor range, ~K (MPa ..,J;.) GP41-092274 

Figure 59. FCGR results on AI-3Li-l.5Cu-lMg-O.5Co-0.2Zr VA forgings, T8511530 temper: (0) 

LT orientation, (0) TL orientation, ( - - - ) typical results for 7075-T651 (24), and 

(-- -) typical results for 7091-T7E69(24). 

79 
Sim1lar orientation effects are seen 1n FCGR results on the AI-3L1-1.5CuIMg-0.5Co-0.2Zr extrusions and forg1ngs (F1gures 58 and 59). The stress 

1ntens1ty factor requ1red to produce a crack-growth rate of 10-7 1S about 7 to 

8 MFa.m1/2 for all alloy extrus10ns and forgings, but the slope of the da/dNvs.-~K curve for the Li-conta1ning alloy 1S much larger than that for the nonLi-conta1ning alloy, 1ndicat1ng higher growth rates at larger ~K values. This 

susceptibility of crack-growth rate to increased load for the L1-conta1ning 

alloy may be related to the 1rnperfect consol1dation procesS1ng, wh1ch 1S also 

reflected in degraded transverse tenS1le propert1es. Crack-growth rates for 

the Li-containing alloy extrusions and forgings are smaller than those for 11M 

7075-Al and p/M 7091-AI at ~K values less than about 7 MPa.mI/2 , but 1ncrease 

rapidly, and greatly exceed crack growth rates for the commerc1al alloys at 

higher t.K. 

6.2.6 Creep Resistance 

Results of creep experiments on Al-4.4Cu-l.5Mg-IFe-lNi-0.2Zr T851 

extrusions and forgings at 1500C are shown in Figure 60. On an absolute 

rupture stress basis, the extrusions perform better than the forg1ngs, but the 

performance of the forgings is superior when normal1zed to yield strength. 

'" 08 

'" '" CI) 

... 

.... '" CI) '" .... ... .... CI) '" .... "''0 .... =- .... .... CI) .... 

=> 07 .... 

~ 

06 

10 100 1000 10000 

time (h) GP41-0922-7S 

Figure 60. 150°C creep-test results for T851 temper AI-4.4Cu-1.5Mg-lFe-lNi-0.2Zr: (0) V A extrusions 

and ( 0) VA forgings, normalized to yield stress; (- - -) tYPical results for 2124-T851. 

80 
The rupture time of extrus10ns dnd forgings 1S comparable to that of 2124-AI 

at h1gh loads, but increases slowly with decreas1ng load and 1S a factor of 20 

to 50 smaller than that of 2124-T851 at 60% of the Y1eld stress. 

6.2.7 Stress-Corrosion-Cracking Res1stance 

Results of C-ring tests on Al-3L1-1.5Cu-1Mg-O.5Co-0.2Zr T8S1 extrus10ns 

and forgings are shown in Figures 61 and 62. Because of d1ffer1ng geometries 

of the extruS10ns and forgings, the extrusions were tested 1n the T or1entation while the forgings were tested 1n the more-susceptible S orientation. 

Failure times were relat1vely short for both product forms and particularly so 

for forgings, where failure at 50% of the yield stress occurred in all samples 

in less than 100 h. None of the non-Li-conta1ning alloy C-rings, tested under 

the same relative loads, failed in 1000 h. The relat1ve performances of 000-

Li- and Li-containing alloys and Li-containing alloy extrusions and forg1ngs 

are similar to relative performances of the same alloys and product forms in 

longitudinal and transverse tensile tests, suggesting that the effectivenegs 

of consolidation processing and interparticle bonding largely determines the 

stress-corros1on-cracking (SeC) resistance"of the alloys. 

400 ~ 

.... I I I I .... .... .... .... .... .... .... .... 

300 .... f- .... .... - .... .... 

'2 .... 

0.. ** 

6 .... 

til 

til 

I1J 200 .- .... - til 

"0 

~ 

0. 

0. 

<C 

100 r- *** -

• Old not fat! In 1000 h 

0 I I I I 

0 200 400 600 800 1000 

TIme to faIlure (h) GP41 092276 

Figure 61. Tlme-to-fallure m alternate ImmerSIOn C--rmg tests of AI-3LI-l 5C-u-lMg-O 5C-o-O.2Zr 

V A extrusIOns· T8511530 temper, T orientatIOn 

81 
400~--~~--~----~-----T----~ 

300 

I 

I 

\ 

'2 \ 

0.. \ 

6 2ooE==',- ~ , II) .... 

... .... cI5 .... 

100 

.... .... 

---

"Did not fall In 1000 h 

-------

* 

OL-__ ~ ____ ~ ____ ~ ____ ~ __ ~ 

o 100 200 300 400 500 

Time to failure (h) 

Figure 62. Time-to-failure m alternate immersion Coring tests of AI-3LI-1.5Cu-lMg-O 5LO-O 2Zr 

V A forgmgs: T8511530 temper, S orientation. 

82 
7. TASK 4: CORRELATION OF PROCESS/STRUCTURE/PROPERTIES RELATIONS 

7.1 Effect of Particulate Type on Properties 

Results presented in Figu~es 41, 42, 45, and SO-53 show that the 

add1tional microstructural ref1nement ach1eved by inc~easing the particulate 

so11d1f1cat10n ~ate from 2000-5000 K/s (typical of VA and USGA-Ar powders) to 

SO 000-120 000 K/s (typical of USGA-He powders) or 400 000-500 000 K/s 

(typical of RQ flakes) does not improve the tensile properties of conso11dated 

Eo~ms. The observed variations 1n tensile properties among VA, USGA-Ar, and 

USGA-He extrusions can safely be attributed to differences in Mg concentrat10ns 1n the non-L1-contain1ng alloys, and to differences in Li and Mg concentrations in the L1-contain1ng alloy (Tables 2 and 3). Tensile properties of 

the Li-containing RQ extrusion fall significantly short of those of the VA 

extrus10n, despite the larger solidification rate of the RQ flakes. 

Examination of extrusion micrographs, particularly Figures 32, 33, 35, 

36, 37, 38, and 40, reveals little difference in microstructure among extrusions made from d1fferent particulate types. Subgrain diameters, S1zes and 

spac1ngs of incoherent dispers01ds, and morphologies and distributions of precipitates are similar among all extrusions and virtually identical among 

powder extrusions. The RQ extrusion has fewer oxide particles than the powder 

extrus10ns owing to the larger particulate size (Figures 35d and 36d), and it 

also has a more un1form distribution of Al9Co2 dispersoids (compare Figures 

40a, d, e, and f). The volume fractions of oxides and dispersoids are too 

small to affect tensile properties, however, and the comparable ductilities of 

VA and RQ extrus10ns indicate that Al9Co 2 dispersoids are not spaced closely 

enough to affect ductility. 

Within the range of part1culate solid1fication rates covered 1n this 

study, microstructures of consolidated forms are determined primarily by consolidation details such as degassing condit10ns, presence or absence of hot 

press1ng, extrusion temperature, and extrusion ratio. prov1ded that a m1n1mum 

so11d1ficat10n rate, which must lie below 2000 K/s, is attained, the scale of 

segregation of major alloying elements is suffic1ently fine as to be completely re11eved during consolidation processing. Further increases 1n solid1E1cat10n rate, wh1ch result in finer micro-segregation, decrease the time requ1red 

fo~ segregation to be relieved, but cannot further homogenize the soluble 

83 
alloY1ng elements. T1mes and temperatures 1nvolved 1n convent10nal consol1dation process1ng are suff1ciently long that 1ncoherent d1sperso1ds coarsen to 

about the same Slze and spac1ng, regardless of the 1n1t1al d1str1but1on of the 

d1spersoid-forming elements Fe, N1, or Co. The d1str1butlon of lncoherent 

disperso1ds 1S altered when the original SOlld1f1cat10n 1S non-dendrltic, as 

1n the case of the RQ L1-conta1n1ng alloy. Th1S change 1n SOlld1f1catlon mode 

with increas1ng solldiflcation rate prom1ses a more unlform distrlbut10n of 

d1spersoids and suggests that ch1ll-block-quenched splats are super10r to 

powders for alloys containing large volume fract10ns of d1sperso1ds. 

The substructure of RSP Zr-containing alloys 1S pr1mar1ly f1xed by the 

extrusion process, not by particulate type or SOlldificat10n rate. The coherent, cubic Al3Zr d1sperso1ds responsible for substructure stab1l1zation are 

formed during consol1dation rather than dur1ng SOlld1f1cation, and hence their 

distr1bution is not affected by increases in the SOlld1f1cat10n rate above the 

threshold required to place 0.2 wt% Zr in metastable SOlld solution. Extrusion pole f1gures (Figures 26 and 27) show domination of crystallographic 

texture by the extrusion process and exceptional microstructural stability 

during subsequent solution annealing. The final substructure might be sensitive to particulate type if the substructure refining or stabilizing element 

were weaker in its effect, e.g., Cr or Mn. 

7.2 Effect of Processing Variables on Properties 

7.2.1 Processing Sequence 

To obtain satisfactory properties in forms consolidated from aluminum 

alloy particulates, the 1.5- to 3-nm thick pre-ex1st1ng oX1de layers must be 

effectively d1s1ntegrated to allow the alum1num substrates to bond to one 

another during b1llet formation and extruS10n or forging. OX1de dlslntegratlon depends on effective removal of adsorbed oxygen and water vapor during 

degassing and on the time, temperature, and amount of deformation during 

billet making and final consolidation. Although th1S study was not pr1mar1ly 

directed towards examination of process1ng variables, some conclus10ns can be 

drawn regard1ng the influence of degassing pract1ce and presence or absence of 

a hot pressing step on properties of consol1dated forms. 

84 
The results presented in Section 5.1 indicate that the degass1ng cycle 1S 

much more critical for Li-containing alloy powders than for non-Li-conta1ning 

alloy powders. If degassing is performed pr1marily at h1gh temperatures, 

oxygen-bearing gases react with the powders rather than being removed, and 

oxide layer thicknesses can increase by as much as a factor of 9 (Figures 18b 

and 22b). The relatively slow degassing cycle used during product10n of extrusions or forgings in this study resulted in roughly comparable increases in 

oxygen concentrations in consolidated forms vis-a-vis loose powders in both 

alloys. Oxide layers in the Li-containing alloy extrusions appear to have 

been well d1sintegrated (Figure 40), and it is considered that the poor transverse tensile properties observed in Li-containing extrusions are not the 

result of inadequate degassing practice. Superior longitudinal and transverse 

ductilities at strength levels comparable to those reported in Figures 50 and 

51, for example, have been reported (20,21) in RSP Al-3Li-Mg-Zr and Al-3Li-CuZr extrusions which had oxygen levels comparable to those in this study. 

Oxygen levels <0.1% in RSP Li-containing alloy extrusions are probably acceptable, provided consolidation processing is efficient enough to disintegrate 

the oxide layers. 

Comparison of mechanical properties strongly suggests that the 

consolidation processing sequence which was effective for the non-Li-containing alloys did not result in effective consolidation for the Li-containing 

alloy. For instance, transverse tensile properties for the AI-4.4Cu-l.5MgIFe-lNi-0.2Zr VA extrusions were only slightly lower than the longitudinal 

properties (Figure 43), while transverse properties of the AI-3Li-l.5Cu-lMg0.SCo-0.2Zr extrusions were degraded by 20% or more in strength and by 50% in 

ductility (Figure 51). Ductilities of Li-containing alloy forgings were much 

lower than those of non-Li-containing forgings (compare Figure 43 with 

Figure 50). Fracture toughnesses (Table 9), the variation of FCGR with AK 

(compare Figures 58 and 59 with 56 and 57), and SCC data (Figures 61 and 62) 

were all worse for the Li-containing than for the non-Li-containing alloy. 

An alternative to adducing consolidation deficiencies to explain 

1nadequate properties in the Li-containing alloy is to suppose that RSP Al-3Li 

alloys merely have intrinsically poor properties. However, recent IRAD work 

at MDRL on RSP AI-3Li-Zr, AI-3Li-Mg-Zr, and AI-3Li-Cu-Zr alloy extrusions has 

shown near-equality of longitudinal and transverse properties at a yield 

85 
stress level of 440 MPa (21), and only a 10% loss of Y1eld strength in 

transverse compared w1th longitudinal properties at a yield stress level of 

510 MPa (20). ~o failures of T-orientation SCC C-ring specimens were noted 1n 

tests lasting several months (20). The extrusions were made from VA powders, 

and compositions were s1milar to that of the Li-contain1ng alloy tested in 

th1S study. The major difference between the two mater1als was 1nsert10n of a 

vacuum hot-pressing step pr10r to extrus10n in the MDRL IRAn alloys (20), a 

step wh1ch was omitted in the extrusions for th1s study. 

It is likely that more careful conso11dation process1ng of L1-contain1ng 

extrusions and forgings would improve properties, particularly notch-sensit1ve 

propert1es such as toughness, FCGR, and SCC resistance. In extrus10ns, vacuum 

hot pressing to 100% density polygon1zes powder part1cles and starts the oxide 

disintegration process. Extrusion of a 100% dense compact produces more plastic flow in the neighborhood of interparticle boundaries than does extrusion 

of a 70% dense cold-pressed compact, in which densification and plast1c flow 

must occur simultaneously. The process of disintegrating oX1de layers into 

discrete particles, w1th concomitant bonding of fresh alum1num alloy surfaces, 

can proceed farther in the limited time available for extrus10n 1f the starting form is a 100% dense billet. Kinetics of plastic flow and d1s1ntegrat10n 

of oX1des are more important 1n Li-containing than in non-Li-conta1ning alloys 

owing to the thicker and possibly more adherent oxide layers present 1n the 

former after degassing. 

In forgings, an upset of 55%, sufficient to ensure good properties in the 

Al-4.4Cu-1.5Mg-1Fe-1Ni-0.2Zr alloys, and larger than that successfully used 

for Al-Fe-Ce alloys (2), was 1nsufficient for adequate interpart1cle bond1ng 

1n the Li-containing alloy. In this case, the alloy was hot pressed before 

f1nal forging, but the Li-conta1ning oxide layer was more res1stant to d1s1ntegration. Multiaxial working is probably necessary to produce L1-conta1n1ng 

forgings with good properties. For instance, ABC (multiax1al) forg1ng of an 

RSP Al-3Li-1.5Cu-1Mg-0.5Co-0.2Zr alloy resulted 1n a consolidated form w1th 

properties equivalent to those of an extrusion of 20:1 extrusion rat10 (26). 

86 
7.2.2 Texture Effects 

The difference 1n strength between VA extruS10ns and forg1ngs 1S 80 to 90 

MPa for both alloy compositions (Figures 43 and 50). This d1fference is correlated with the difference between the strong texture of the extrus10ns and 

the weak, axisymmetric texture of the forgings. Since both extrus10ns and 

forgings have a fine subgrain microstructure, the d1fference is pr1mar1ly one 

of subgrain orientation rather than subgrain size. As the test temperature is 

raised, the subgrain boundaries cease to be effective barriers to slip, HallPetch and texture effects disappear, and the strength of extrusions approaches 

that of forgings (F1gures 46 and 47). To increase the strength of forgings, 

the forging strain must be increased to permit development of a more marked 

texture. This can be accomplished by 1ncreasing the upset in a uniax1a1 

forging, or by performing multiaxial forging. 

7.3 Effect of Alloy Chemistry and Microstructure on Properties 

7.3.1 Non-Li-Containing Alloys 

The alloy Al-4.4Cu-I.5Mg-IFe-INi-0.2Zr was designed to 1mprove I500C 

ultimate tensile strength from 405 MPa, typical of 11M 2I24-Al, to 480 MPa, 

which would be required for density-normalized equivalence to Ti-6AI-4V, by a 

combination of dispersoid and substructure strengthening. Examination of 

strength as a function of test temperature (Figure 46) shows that the tensile 

strength was increased to 440 MPa at I500C and the temperature stability of 

strength was improved below I80oC; however, the goal of equivalence to Ti-6AI4V was not reached. 

The yield stress increments resulting from dispersion strengthening and 

substructure refinement can be estimated using the Orowan and Hall-Petch relations, respectively. The Orowan bypassing shear stress in a polycrystal11ne 

aluminum alloy is given by 

Gb 

T = 3.39 -- R,n 

211'D 

x 

4b (1) 

where G = shear modulus of matrix, b = Burgers vector length, D = average 

interparticle spacing, and x = average particle diameter. For Al-4.4Cu-I.5MgIFe-INi-0.2Zr at I50oC, G is assumed to be the same as for 2024-Al, e.g., 26.4 

87 
CPa (22) (corrected for temperature), and b = all = 0.286 rum for Al. From 

F1gure 37a, x = 0.46 x 10-6 m. The VA extrus10n has 0.94 wt% Fe and 0.92 wt% 

N1, equivalent to 4.4 vol% Al9FeN1 d1sperso1ds. Assum1ng a cub1cal array of 

evenly spaced part1cles, the average 1nterpart1cle spacing D 1S 1.04 x 10-6 m. 

Substituting these values 1nto equat10n (1), the calculated Y1eld stress 

1ncrement 1S 23.6 MPa. 

The Hall-petch strengthen1ng is g1ven by 

a = a + kd 1/ 2 

° 

(2) 

where for Al, a = fr1ct1onal stress = 1.6 kg/(mm)2, k = Hall-Petch slope = 

o 

0.22 kg/(mm)3/2 (27), and d = gra1n or subgra1n d1ameter. Assuming a subgrain 

d1ameter of 2 um 1n the Al-4.4Cu-l.SMg-IFe-lNi-0.2Zr VA extrusion 

(Figure 37a), the substructure strengthening is 64 MPa at room temperature, 

and probably somewhat less at IS0oC. 

The contributions of substructure and d1spersoids are not additive, 

particularly in the presence of an array of fine, semi-coherent 8' precipitates, as in the T8S1 temper. The yield stress increment at IS00C resulting 

from 0' precipitation corresponds approximately to the difference in yield 

stress between 2024-T3S1 and -T8S1 at that temperature, or about 90 MPa. 

Assuming that subgrain boundaries, dispersoids, and precipitates are all 

strong barriers, their overall strengthening effect as a first approximation 

is the square root of the sum of squares of the individual contributions, or 

113 MPa. In the presence of precipitates, the additional strengthening effect 

of substructure and disperso1ds is calculated to be about 23 MPa. The 

observed strengthen1ng at IS00C is 22 MPa. At room temperature the effects of 

substructure and disperso1ds are absent, suggest1ng that the precip1tates are 

stronger barriers at 2SoC than at IS0oC. D1fferences 1n texture between I/M 

and P/M materials are not included 1n the above calculat10n. 

This d1scussion suggests that the effect of ref1ned substructure and 

disperso1ds should be greater in the absence of prec1pitates, 1.e., 1n the 

T3S1 (naturally aged) temper. Indeed, the yield strength of VA Al-4.4Cu1.SMg-IFe-lNl-0.2Zr-T3S1 is 417 MPa, compared to 324 MPa for 2024-Al. The 

increment here is close to the sum of the 1ndlvidual contrlbutions, suggesting 

88 
that both substructure refinement and d1spers01ds contribute s1gn1ficantly to 

strength improvement in the T3SI temper. 

At temperatures above ISOoC, the grain or subgra1n boundar1es cease to be 

effective barriers, and Hall-Petch and texture effects cease to contribute to 

strength. This effect has been observed 1n RSP Al-Li alloys by Sastry and 

O'Neal (28). The relatively small effect of dispersion strengthening is 

offset by the lower Cu content of the RSP alloy. The overall analys1s of 

strengthening contributions strongly suggests that the intermediate-to-high 

temperature properties of RSP Al alloys can be improved by dispersion 

strengthening only if the dispers01ds are significantly finer and more closely 

spaced than was the case 1n this study. Th1s can be accomplished by solid1fying at larger rates than used here, or by selecting elemental additives to 

form dispersoids which coarsen at a slower rate than does Al9FeNi. 

7.3.2 Li-Containing Alloy 

The microstructures of RSP Al-3Li-I.SCu-IMg-0.SCo-0.2Zr alloys in the 

solution-treated and aged conditions were generally satisfactory. The 

(Li+Cu+Mg) content of the alloys only slightly exceeded the solubility limit 

at the highest safe solution-treatment temperature of 5600C (Figure 36). 

Thus, few hard constituent particles were available to promote fracture and 

reduce ductility. OX1de layers on the particulates appeared to be well disintegrated, at least 1n the extrusions (for instance, Tigure 40a), suggesting 

that oxygen levels were not excessive. The substructure was thermally stable 

and fine. However, the spacing and distribution of A19C0 2 d1spersoids were 

not sufficient, except perhaps in the RQ extrusion, to significantly increase 

ductility by re11eving planar slip. 

In view of the favorable m1crostructures of the Li-containing alloy 

extrusions and the high strength obtained on a sim1lar alloy 1n another study 

(29), the defic1enc1es 1n properties of the L1-containing alloy are attr1buted 

to problems in consolidation processing rather than deficiencies in alloy 

chemistry or microstructure per see It appears that the addition of Co to 

th1s alloy was superfluous. Also, while the addition of IMg to an Al-3Li1.5Cu-0.2Zr matrix may result in a significant increase 1n strength, the 

reduction in ductility to 2 to 4% is a serious offsetting problem. If alternative solution and aging treatments of Al-L1-Cu-Mg-Zr alloys do not produce 

89 
sign1ficant ductility improvements with compensat1ng strength reduct10ns, 

reversion to Al-L1-Cu-Zr alloys may be required to opt1mize propert1es. 

7.4 Effect of Heat Treatment Variables on Properties 

The effect of heat treatment on properties was not stud1ed extensively 

for the non-Li-conta1ning alloys. Solut10n treating the Al-4.4Cu-l.5Mg-IFeIN1-0.2Zr VA extrUS10n at 5000C rather than 5300C resulted Ln 20 to 40 MPa 

decreases in yield and ultimate tensile strengths oW1ng to less complete solutionizing of Cu and Mg, which resulted in smaller volume fractions of 8~ 

precipitates. As d1scussed 1n Section 7.3.1, the effect of RSP on this alloy 

was more pronounced in the naturally aged than in the peak-aged condition 

because of the larger effects of substructure refinement and dLspers10n 

strengthening when not combined with prec1pitation strengthen1ng. 

The strength increments in the peak-aged Li-containing extrusions caused 

by increasing the prior solution-treatment temperature from 530 to 5600C were 

25 to 30 MPa in yield strength and 40 to 50 MPa 1n ult1mate tensile strength, 

without a noticeable change in ductility. ThLs result is also the effect of 

increased solutionizing of Li, Cu, and Mg. The effect of a 2% stretch between 

solution treating and aging is similar to that of increasing the solutionannealing temperature (Figure 53), and is the result of 1ncreased volume 

fraction and homogeneity of precipitation of the T1(Al 2CuL1) and s1m1lar 

phases (30). The effects of stretch and increasing solution temperature are 

not additive, possibly as a result of deficiencies in consol1dat10n processing. These two effects were found to be additive in Al-3Li-2Cu-0.2Zr, where 

processing had included a vacuum hot pressing step (21). 

7.5 Recommendations for Optimal Composition/Processing/Structure 

Combinations in RSP Aluminum Alloys 

Based on the results of this study, the following recommendat10ns on 

optimal chemistry and processing of RSP Al alloys are: 

1. For Li-conta1ning alloys Lntended primarily for ambient-temperature 

applications, choose a composition such that the combined major alloying 

element addit10ns do not exceed the solubility l1m1t at the solut10nannealing temperature. The alloy should contain Zr for substructure 

90 
ref1nement, but dispersoid-forming element additions to promote duct1lity 

through slip d1spersal are redundant in a fine-gra1ned material. 

2. For alloys 1ntended for intermediate-to-high temperature service, 

d1spers1on strengthening is a valid mechanism for prov1ding a temperature1nsensit1ve strength contribution. However, the spacing of disperso1ds 

must be smaller than the approximately 1 ~m attained 1n th1s study to 

provide a sufficient strength increment. 

3. RSP particulates should be produced by a method which is easy to apply 

and yields clean, easily handled particulates. The absolute so11dif1cat1on 

rate is not 1mportant for Al alloys not containing dispersoids, provided 

that it exceeds - 103 K/s. 

4. Large amounts of hot work during the consolidation process are required 

to effectively disintegrate oX1de layers and promote interparticle bond1ng. 

Li-containing alloys are more sensitive than non-Li-containing alloys 1n 

this respect. Vacuum hot pressing appears to be a necessary prelim1nary to 

extrusion or forging. Forgings of Li-containing alloys require more deformation to achieve a well-bonded consolidated form than forgings of non-L1-

containing alloys. 

5. Suitable variations of solution-treatment temperature, process stretch, 

and aging temperature and time can result in controlled var1ation of properties in RSP aluminum alloys, as in 11M alloys. The variation of properties with heat treatment depends upon the various strength contr1butions 

and may be different for RSP than for 11M alloys. 

91 
8. CONCLUSIONS 

1. RSP part~culates of Al-4.4Cu-I.SMg-lFe-IN~-0.2Zr and Al-3L~-1.SCu-1Mg0.SCo-0.2Zr alloys can be produced by vacuum atom~zat~on, ultrasonic atom~zat~on ~n argon or hel~um, or twin roller-quench~ng at sol~d~f~cat~on rates 

ranging from 2 x 103 to 4 x lOS K/s. 

2. Ox~de layers on as-solid~fied part~culates are 1.S- to 2-nm thick, 

regardless of chemistry or particulate product~on method. 

3. Vacuum degassing of Al-4.4Cu-I.SMg-lFe-lNi-0.2Zr becomes continuously more 

efficient as the temperature is raised. Degassing of Al-3Li-l.SCu-1Mg-O.SCoO.2Zr becomes less efficient above 4S0oC owing to chemical fixing of Hand ° 

by Li. Degassing of RSP Li-containing aluminum alloy powders should be 

carried out at the lowest practical temperature to assure removal rather than 

chemical fixing of Hand 0, and prevent large increases in oxide layer thickness on degassed powder particles. 

4. Consolidation processing by canning, cold compaction, degassing, and hot 

extrusion is sufficient to yield good mechanical properties in non-Li-conta~ning alloys. Li-containing alloys have thicker particulate oxide layers, more 

adherent layers, or both, and require a vacuum hot press~ng step before extrusion to achieve adequate interparticle bonding and good mechanical properties. 

S. Uniaxial forging of non-Li-containing alloys to SO to 60% upset ~s 

sufficient to assure good mechanical properties. The same process is not adequate for Li-containing alloys, and larger upsets or multiaxial forg~ng are 

required to produce good properties in Li-containing alloy powder forgings. 

6. Differences in particulate type have l~ttle influence on properties of 

consolidated forms ~n either Al-4.4Cu-l.SMg-lFe-1N~-0.2Zr or Al-3L~-I.SCu-1MgO.SCo-O.2Zr extrusions or forgings. In both alloys, the final microstructures 

and properties are dominated by the effects of consolidation process~ng. 

7. Deficiencies in consolidation processing of the Li-conta~ning alloy are 

reflected most strongly ~n the transverse tensile propert~es and in notch92 
sensitive properties such as fracture toughness, fatigue-crack-growth rate, 

and stress-corrosion-cracking resistance. 

8. D1fferences in strength between extrusions and forgings are largely the 

result of differences in texture caused by differences in degree of hot work 

during consol1dation. Strengths of the two product forms approach one another 

1n testing above 1500C owing to the ineffectiveness of grain boundaries as 

barriers to slip at high temperatures. 

9. Strength/ductility combinations in both non-Li- and Li-containing alloys 

are strongly affected by combinations of solution treatment, process stretch, 

and aging treatment var1ables. 

10. The dispersoid-forming system Fe+Ni is read1ly added to an Al-4.4Cu-l.SMg 

base composition by RSP, but the resulting dispersoids are too coarse and too 

widely spaced to contribute significantly to intermed1ate-temperature strength 

in conventionally processed consolidated forms. Alternative processing or a 

more effective dispersoid-forming system are required to achieve significant 

dispersion strengthening in this alloy system. 

11. The combined effects of refined substructure and.f1ne, 1ncoherent 

dispers01ds in AI-4.4Cu-l.5Mg-IFe-lNi-0.2Zr-T851 raise the IS00C ultimate 

tensile stress relative to l/M 2124-T851 by about 10%, compared w1th the 20% 

increase in strength which would be required for density-normalized strength 

equivalency to Ti-6Al-4V. The effect of substructural refinement is more 

sign1ficant (110 MPa increase in yield strength) in the T351 temper. 

12. Both substructure refinement by Zr add1tions, and solution or d1spers10n 

strengthening by Mg addition, contr1bute sign1ficantly to the strength of Al3Li-l.SCu-lMg-0.SCo-0.2Zr. However, Mg add1t10ns also decrease the duct1lity 

and fracture toughness, and better strength/ductility combinations may be 

available from Al-3L1-Cu-Zr rather than Al-3L1-Cu-Mg-Zr alloys. 

13. The add1tion of 0.5Co to AI-3Li-l.5Cu-lMg-O.2Zr does not produce a 

sufficient volume fraction of fine, even dispersoids to provide a significant 

duct1lity increase by Orowan bypassing. Dispers01d-forming element additions 

to a fine-grained alloy are superfluous for ductility enhancement. 

93 
14. RSP non-Li- and Li-containing alloys are promising subst1tutes for 

commercial 2XXX and 7XXX alloys in both ambient- and elevated-temperature 

applications. However, further improvements in alloy chem1stry and consol1dat10n processing are required for these alloys to fully real1ze their 

potential. 

94 
9. REFERENCES 

1. S. L. Langenbeck et al., Elevated Temperature Alloy Development, Interim 

Technical Report, Lockheed California Co., Rpt. No. LR29977, AFML 

Contract No. F33615-81-C-5096 (September 1981). 

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Alloy Development, Report No. AFWAL-TR-81-4076 on Air Force Contract 

No. F33615-77-C-5086 (September 1981). 

3. J. P. Durand, R. M. Pelloux, and N. J. Grant, Materials Science and 

Eng1neering~, 247 (1976). 

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Diboride Dispersions by Rapid Solidification, in Rapid Solidif1cation 

Processing: Principles and Technologies II, R. Mehrabian, B. H. Kear, 

and M. Cohen, eds. (Claitor's, Baton Rouge, LA, 1980), p. 154. 

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Missiles and Space Co. Report LMSC-D674504, DARPA/AFML Contract No. 

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Microstructure and Properties of Powder-Processed Alum1num-Lith1um 

Alloys, McDonnell Douglas Report MOC Q0778, AFOSR Contract No. F49620-79-

C-0039 (May 1982). 

8. K. K. Sankaran, J. E. O'Neal, and S. M. L. Sastry, Effects of Secondphase Dispers01ds on Deformation Behavior of Al-Li Alloys, Metallurgical 

Transactions 14A, 2174 (1983). 

95 
9. N. J. Grant, A Review of Various Atomizat10n Processes, in Rap1d 

Solid1fication Processing: Pr1nciples and Technologies, R. Mehrabian, 

B. H. Kear, and M. Cohen, eds. (Claitor's, Baton Rouge, LA, 1978), p. 

230. 

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Metals~, 20 (1983). 

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Mod1fied 7075 Aluminum Alloy by Ultrason1c Atomization, Report No. ADA085141 on Off1ce of Naval Research Contract No. N00014-78-C-0385 (March 

1980). 

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T. H. Sanders, Jr. and E. A. Starke, Jr., eds. (TMS-AIME, Warrendale, PA, 

1981), p. 171. 

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Deformation of Rapidly Solidified AI-3Li Alloys Containing Incoherent 

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E. A. Starke, Jr., eds. (TMS-AIME, Warrendale, PA, 1981), p. 189. 

15. D. P. Ames, R. R. Iran1, and C. F. Callis, The Agreement Between 

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D1vided Powders Includ1ng Phosphates, Journal of Physical Chemistry ~, 

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96 
17. L. F. Mondolfo, Alum1num Alloys: Structure and Properties (Butterworths, 

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the Dendrite Spacing in Splatcooled Aluminum Alloys, Journal of the 

Institute of Metals~, 30 (1968). 

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Compositional Variations on the Microstructures and Properties of Rapidly 

Solidified Al-3Li Alloys, in Proceedings ASTM Symposium on RapidlL 

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and Corrosion-Fatigue Behavior of a Rapidly Solidified Al-Li-Cu-Mn Alloy, 

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eds. (TMS-AIME, Warrendale, PA, 1981), p. 307. 

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Versuchsanstalt fuer Luft- und Raumfahrt, Cologne (1979). 

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97 
26. I. G. Palmer, R. E. Lewis, and D. D. Crooks, The Des1gn and Mechan1cal 

Properties of Rapidly Solidified Al-Li-X Alloys, in Aluminum-L1thium 

Alloys, T. H. Sanders, Jr. and E. A. Starke, Jr., eds. (TMS-AIME, 

Warrendale, PA), p. 241. 

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(Applied Science Publishers, London, 1971), p. 342. 

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and Mechanical properties of Rapidly Solidified Al-Li-Co and Al-Li-Zr 

Alloys, in Aluminum-Lithium Alloys II, T. H. Sanders, Jr. and 

E. A. Starke, Jr., eds. (TMS-AlME, Warrendale, PA, 1984), p. 79. 

29. I. G. Palmer, R. E. Lewis, D. D. Crooks, E. A. Starke, Jr. and 

R. E. Crooks, Effect of processing Variables on Microstructure and 

Properties of Two Al-Li-Cu-Mg-Zr Alloys, in Aluminum-L1thium Alloys II, 

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1984), p. 91. 

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(1981). 

98 
Standard Bibliographic Page 

1 Report No 12 Government AccessIOn No 

NASA CR-177946 

4 Title and Subtitle 

Study on Effects of Powder and Flake Chem1stry and 

Morphology on the Properties of Al-Cu-Mg-X-X-X Powder 

Metallurgy Advanced Aluminum Alloys 

7 Author(s) 

P. J. Meschter, R. J. Lederich, J. E. O'Neal and 

P. S. Pao 

9 Performing OrganizatIOn Name and Address 

McDonnell Douglas Research Laborator1es 

P.O. Box 516 

St. LOU1S, MO 63166 

12 Sponsoring Agency Name and Address 

National Aeronautics and Space Administration 

Langley Research Center 

Hampton, VA 23665 

15 Supplementary Notes 

Langley Technical Monitor: W. Barry Lisagor 

.Final Report 

11\ Abstract 

3 Recipient's Catalog No 

5 Report Date 

November 1985 

6 Performing OrganizatIOn Code 

8 Performing OrganizatIOn Report No 

MDC 01227 

10 Work Unit No 

11 Contract or Grant No 

NASl-17l07 

13 Type of Report and Period Covered 

Contractor Report 

14 Sponsoring Agency Code 

A study was conducted to deterrn1ne the effects of alloy chem1stry and part1culate 

morphology on conso11dation behavior and consolidated product propert1es 1n rap1d 

solidification processed, powder-metallurgical Al-3Li-l.5Cu-lMg-0.5Co-0.2Zr and 

Al-4.4Cu-l.5Mg-Fe-Ni-0.2Zr extrusions and forgings. M1crostructures and mechan1-

cal properties of both alloys were largely unaffected by particulate product10n 

method (vacuum atom1zation, ultrasonic atomization, or twin-roller quenching) and 

I by particulate so11dification rates between 103 and 105 K/s. Conso11dation pro-

! cess1ng by canning, cold compaction, degassing, and hot extrusion was suff1cient 

to yield mechanical properties in the non-Li-containing alloy extrus10ns which 

were similar to those of 7075-Al, but ductilities and fracture toughnesses were 

1nferior owing to poor interparticle bonding caused by lack of a vacuum-hotpress1ng step during consolidation. Mechanical properties of extruS10ns were 

superior to those of forgings owing to the stronger textures produced by the 

more severe hot working dur1ng extrusion. The effects on mechan1cal propert1es 

of d1spersoid size and volume fraction, substructural refinement, so11d solution 

I strengthen1ng by Mg, and precip1tate S1ze and distribution were elucidated for 

I both alloy types. RSP non-Li- and Li-containing alloys were found to be prom1s1ng 

subst~tutes for comrnerclal 2XXX and 7XXX alloys in both amblent-temperature and 

elevated-temperature applicat10ns. Further 1mprovements in alloy chem1stry and 

Lonso11dat10n process1ng were recommended for the RSP alloys to fully rea11ze 

I the1r potential. 

17 Key Words (Suggested by Authors(s)) 

Rapld solidlficat10n 

Aluminum-11thlum alloys 

Consolidation 

Powder metallurgy 

Extrus10n 

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Unclasslf1ed 

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Unclassified - Unlimited 

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