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WADC TECHNICAL REPORT 54-278

PART 3

DEVELOPMENT OF TITANIUM-BASE ALLOYSFOR ELEVATED TEMPERATURE APPLICATION

WILLIAM F. CAREW

FRANK A. CROSSLEY

DONALD J. McPHERSON

ARMOUR RESEARCH FOUNDATION

OF ILLINOIS INSTITUTE OF TECHNOLOGY

MAY 1956

WRIGHT AIR DEVELOPMENT CENTER

WADC TECHNICAL REPORT 54-278

PART 3

DEVELOPMENT OF TITANIUM-BASE ALLOYSFOR ELEVATED TEMPERATURE APPLICATION

WILLIAM F. CAREW

FRANK A. CROSSLEY

DONALD 1. McPHERSON

ARMOUR RESEARCH FOUNDATION

OF ILLINOIS INSTITUTE OF TECHNOLOGY

MAY 1956

MATERIALS LABORATORY

CONTRACT No. AF 33(616)-2853

PROJECT No. 7351TASK No. 73510

WRIGHT AIR DEVELOPMENT CENTERAIR RESEARCH AND DEVELOPMENT COMMAND

UNITED STATES AIR FORCE

WRIGHT-PATTERSON AIR FORCE BASE, OHIO

Carpmlntr Litho & Prtg. Co., Springfield, 0.600o- ,inn. l 6

FOREWORD

This report was prepared by the Armour Research Foundation under USAFContract No. A? 33(616)-2853. This contract was initiated under Project No.7351, "Metallic Materials", Task No. 7351O, "Titanium Metal and Alloys",formerly RO No. 615-11, "Titanium Metal and Alloys", and was administeredunder the direction of the Materials laboratory, Directorate of Research,Wright Air Development Center, with Lt. D. A. Wrack acting as projectengineer.

This report covers period of work from January 7 to December 31, 1955.

WADC TR 54-278 Pt 3

ABSTRACT

The principal objective of the work reported herein was a determinationof the effects on mechanical propertia s of complexing the a and 0 phases ofa promising t.A type alloy, Ti-6Al-3Mo. Tin and zirconium were employed asCL complexers and chromium, manganese, and vanadium were employed as P cm-plaxers.

cnmplexing was found to improve creep resistance and rupture strength,while 0 cmplexing reduced these parameters below the leves of the base compo-sition.

In addition, age hardening characteristics of Ti-Al-Ag alloys were deter-mined. Limited creep-rupture data indicated inferiority to a Ti-6AI binarycomposition.

Further studies on the natire of embrittlement in binary Ti-Al alloyswere carried out and results of these studies are reported.

PUBLICATION REVIEW

This report has been reviewed and is approved..

FOR THE COAND:

. R. 14'II,CHE

Technical DirectorMaterials LaboratoWDirectorate of Research

WADC TR 54-278 Pt 3 iii

TABLE OF CON

Page

I. INTRODUCTION 1Objectives of Proam 1

II. EXPERENTAL PROCEDURE . ........ 2

A. Preparation of Alloys .2..... . 2

B. Age Hardening Studies 5...............

C. Tensile Testing 5............... ..

D. Creep and Creep-Rupture Testing ........ . . .7E. Stability Testing. ...... . 7

F. Impact Testing. . . . . . . . . ........ 7

G. Bend Testing . 7

III. RESULTS AND DISCUSSION. o.. . . . .. . 8

A. Age Hardening of Ti-5A1-2Ag and Ti-5A1-7Ag Alloys ...... 8

B. Results for Forging Alloys. .............. 8C. Results for Sheet Alloys ..... 53

D. Properties of Ti-7A-3Mo Alloy. . . . . . . . .60

E. Factors Affecting Ductility of Titanium-Alminum Alloysin the Range from 6 to 10% Alminuin. . . . . . . . .68

IV. SMARY AND CONCLUSIONS ............... .. 86

ERENCES..................... .90

APPENDI ...................... 91

WADC TR 54-278 Pt 3 iv

IST OF TA

1 Materials .0.... . . . . . , .......

2 Master Alloys . . . .o............

3 Chemical AnalysLs of Experimental Al oy. .s ....... . .

4 Aging Characteristics of Ti-Al-Ag Alloys . . . . . , . ,5 Room Temperature Tensile Test Data for Experimental Forging A3leys . .36 Room Temperature Tensile Data for Vacuum Annealed xperixental

Forging Alloys . . . . . . . . . . . .. . .. ..7 Tensile Test Data for Experimental Forging Allos at 620F, .. .28 Tensile Test Data for Experimental Forging Alloys at 8007'.. . , .9 Tensile Test Data for Experimental Forging Alloya at 1020' . . ..

10 Creep-Rupture Data for Experimental Forging Alloys at 80OF . .

U Creep-Rupture Data for Experiental Forging Alleys at 10200 .,

12 Summary of Rupture Strength and Creep Resistance ofExperimental Forging Alloys ..... ..... ,... .

3.3 Stability of Room Temperature Properties of Forging Alloysto Elevated Temperature Exposure . . . . . . . .5

l14 Tensile Data for Experimental Sheet Alloys .... . ... 515 Creep-Rupture Data for Shoot Alloys at 80O and 10206F . , . . .

16 Content of Interstitially Soluble Elements, H, N and 0in Specimens from Forging Study of Ti-7A1-Xo Alloy . . . . . . .6

17 Heat Treatment Comparison for Ti-7Al-3fo (1) Matrial . . . . .618 Heat Treatment Comparison for Ti-7AI-3Ko (IN) Material . . , . .619 Effect of Annealing Temperature on the Tensile Properties

of Ti-6A1 (l1B) Alloy ..

20 Effect of Annealing Temperature n the Tensile Propertiesof Ti-lOA (BB) Alloy... ....... ... . . 9?

21 Effect of Vacutm Annealing on Tensile Properties ofTi-8AI and Ti-lOAl Alloys. .71

22 Temperature Dependence of Tensile Properties of TI-Al Alloys . . . .823 Titaniwl-Almim ALlW*. Debys Lines Other than Aloe Titai . .

WADC TR 54-278 Pt 3 V

LIST OF FIGnUS

Page

1 Age Hardening Characteristics of Ti-5AI-2Ag and Ti-SAl-7Ag Alloys . 10

2 Ti-SAl-7Ag: 1560F-24 hrs-WQ, 930OF-48 hrs-WQ .. ...... 1

3 Ti-SA-7Ag: 1560°F-24 hrs-WQ, 930F-500 hrs-WQ ........ 114 Stress-Rupture and Minimum Creep Rate Curves for

Ti-5AI2Ag Alloy at8O"F 12

% Tensile Properties of Ti-6Al-3Mo Alloy ........... 20

6 Tensile Properties of Ti- -2Mo-lMn Alloy.... .......... 21

7 Tensile Properties of Ti-6A1-2Mo-lCr Alloy ..... 22

8 Tensile Properties of Ti-6A-2Mo-IV Alloy ....... 23

9 Tensile Properties of Ti-6A1-lMo-IMn-IV Alloy . . . 24

10: Tensile Properties of Ti-6A1-IMo-lMn-lCr Alloy .... .... 25

Ii Tensile Properties of Ti-6Ai-3o-iSn Alloy .. . . ... .26

12 Tensile Properties of Ti-6AI-3Mo-lZr Alloy ..... 2713 Tensile Properties of Ti-lOA-3Mo-Zr Alloy . . . . ... . 28

14 Tensile Properties of Ti-6AI-3Mo.2Sn-2Zr Alloy ...... 29

15 Tensile Properties of T-6Al-3Cu Alloy .......... 30

16 Ti-6Al-3Mo, 1470°F-6 hrs-AC, 10207-24 hrs-AC .... . 31

17 Ti-6A-2Mo-iMn, 17O°F-6 hrs-AC, 1020F-24 hrs-AC . . . .... 3118 Ti-6Al-lMo-IMn-lV, 1470F-6 hrs-AC, 1020OF-24 hrs-AC . . . ... 31

19a Ti-6A1-3Cu: 1470eF-6 hrs-AC, l020F-24 hrs-AC ........ 33

19b Ti-6A1-3Cu: VA 1560F-4 hrs-AC, 1020F-24 hrs-AC ..... 33

20 Stress-Rupture Curve for Alloy Ti-6AI-3mo . . ... . . 3521 Stress-Rupture Curves for Alloys Ti-6Al-2Mo, with IMn, ICr or 1V . • 36

22 Stress-Rupture Curves for Alloys Ti-6A1-lMo-In with 1V or 1Cr,and Ti-6Al-3Mo-LSn. . .. ............. 37

23 Stress-Rupture Curves for Alloys Ti-6Ai-3Mo-1Zr, Ti-6A1-3Mo-2Sn-2Zr and Ti6AI-3Cu . 38

24 Minimum Creep Rat* Curves for Alloy Ti-6A-39o . . . . 39

25 MnliafC cV9 ateCurvs forAloyts T6-2N om.3 orCr . .. 14026 Minimum Creep Rate Curves for Alloys Ti-6AI-2Mo-V and

Ti-6A-lMo-IMn-lV ......................... 4127 Minimum Creep Rate Curves for Alloys Ti-6AI-lMo-lMn-lCr and

Ti-6A1-3o-Wjn ... . . ............. 4.. 2

28 Minimum Creep Rate Curves for Alloys Ti-6A1-3(o-WZr andTi-l0Al-3Mo-laZr. . . . . . . . .............. . . 43

WADC TR 54- 2 78 Pt 3 vi

LIST OF FIGURES (continued)

Pal

9 inimum Creep Rata Cu-ves for Alloys Ti-6A1-3Mo-2Sn-22r andTi-6A:L-3Cu 41

30 Tensile Properties of Ti-6Al-2V-lo Alloy . ... ..... .55

31 Tensile Properties of Ti-WA-V-2Mo Alloy . ..... 56 • • • d32 Tensile Properties of Ti-6AI-3V-1o Alloy . . . . . ... . 5733 Stress-Rupture Curves for Alloys Ti-6A1-2V-l and 2M0 and

Ti-6AI-3V-lMo at IO200F. 5934 Correlation of Hydrogen Content with Tensile Ductility for

Ti-7AI-3Mo (4N) Alloy 61................6

35 Influence of Finishing Temperature and Per cent Reduction in Areaon the Tensile Reduction in Area of Ti-7AI-3Mo (JN) Alloy .... 64

36 Ti-7Al-31o (IN), 80% Reduction by Forging at FinishingTemperature of 1950F 67

37 Ti-7AI-3Mo (ON), 75% Reduction by Forging at FinishingTemperature of 18OO°F 67

38 Ti-7AI-3Mor (hN), 75% Reduction by Forging at FinishingTemperature of 1650OF 67................6

39 Proportional Limit and Reduction in Area vs. AnnealingTemperature for Ti-8A1 (2N) Alloy ............ 69

40 Resistivity vs. Annealing Temperature for Ti-8AI (2N) Alloy . 70

41 Reduction in Area vs. Annealing Temperature for Ti-6A1 (IBB) Alloy . 73

42 Fracture Stress vs. Annealing Temperature for Ti-AOAI (IBB) Alloy. . 75

43 Ti-8AI (2N), VA 1470-F-48 hrs-WQ, 750F-64 hrs-AC .. ...... . 77

4 Ti-A1 (lod), VA 1560F-6 hrs-WQ, 1020°F-42 hrs-AC . . .. 77

45 Ti-SAI (Iod), VA 1560"F-6 hrstWQ, 1020"F-500 hrs-AC ..... 7746 Ti-8A1 (Iod), VA 1650F-96 hrs-AC, 1020F-48 brs-AC . . . . . 77

47 Ti-8A1 (2N), VA 2550-F-45 min-FC, 1020eF-48 hrs-WQ. ....... 7948 Ti-lOAl (BB), VA 2550°F-45 min-FC, 1020OF-48 hrs-WQ ...... 7949 Ti-IOAI (IBB), VA 1650°F-96 hrs-AC, 1020OF-48 hrs-AC. ...... 7950 Ti-IOAl (IBB), same as Fig. 49 ............. 79

51 Tensile Reduction in Area vs. Temperature for Ti-6, 8 and 1OAl Alloys 81

52 Notched Impact Strength vs. Test Temperature for Ti-6AI (2) Alloy. . 83

53 Ti-8A1 (lJ), annealed at 1200*F for 1 hour-air cooled, stress-rupture tested at 1020°F under 42,000 psi for 256 hours . . . . . 84

54 Ti-8A1 (2N), same thermal history as Fig. 43 plus l170°F-2 hrs-WQ. . 84

WADC Th 54-278 Pt 3 vii

iST or F!ouR (ontinued)

Pap.

55 Ti-W~ (2Nf), VA 1I70.Jj8 bwuWQ, 10207F-520 hro-WQ ..... .8556 Ti-6Al (2Nf), same, theftal history to Fig. 55 plus lh?70F-2 IreW .f .85

57 Ti-8A1 (Iod), swoi thermal history as rig. 4~5 plus 3470?-) hrm-1Q . .85

58 Ti-.7A1 (lod), oold pressed 70%) 3170OP-l80 hrs4.- .. .. .85

50 Ti-8Al (lad), VA 16&'96hu-C 2F- hrs-WQ . . . . . . .92

61 Ti-8Al (2N), 1560F6 hrosAQ, 1020'1:-500 hro-1% 1470*F-2 hrs-WQ. .92

WADC TR 54278 Pt 3 viii

I. INTRODUCTION

This report summarises the work performed during the period 7 Januarythrough 31 Deoember, 195 on the program entitled ODevelopment of Titanium-Base Alloys for Elevated Temperature Application", under Contract No. AF33(616)-2853.

Objectives of Program

The main objective of this program was to determine the effects on mechani-oal properties of omplexing the a and P phases in a promising a-p type alloy.Previous research encompassing the evaluation of binary, ternary and quaternarytitanium-base alloys demonstrated that Ti-7Al-3Ho has the best combination ofambient and elevated temperature mechanical properties. (1,2,3,4) Because theimmediate objective of the ourrent work was not to develop an alloy with pro-perties superior to the 7Al-3Mo alloy but rather to evaluate the influence ofcomplexing per se, Ti-6A1-3fo was selected as the base material.

Tin and sir onim were selected as a phase eomplexing additives and theresultant alloys, namely, 6A1-3fo-4Sn, 6AI-3(o-lZr and 6A1-3o-2Sn-2Zr, wereinvestigated as potential forging materials.

Additions of P stabilizing elements were made by partial replacement ofthe molybdenum of the base material with chromium, manganese or vanadium. Thetotal amount of P stabilizing additions was maintained at 3%. These alloys,also studied as forging materials, were the following: 6AI-2o-lCr, 6AI-2&o-IMn,6A1-z2o-lV, 6A-lmo-lMn-lCr, and 6A-Imo-IMn-IV.

Because of the excellent elevated temperature performance experienced withalloys based on dispersion-hardening type systms, an a-base alloy having afully developed eutectoid microoonstituent as second phase, Ti-6A1-3Cu, wasexmined for applicability as a forging. alloy.

The 6A1-4V a P alloy developed concurrently under Contract No. AF 33(038)-22806 and Contract No. DA-11-O22-O)-244 (Watertown Arsenal) had shown excep-tional promise in the laboratory as a room and moderately elevated temperaturesheet material. However, its creep-rupture properties at temperature were com-parable only to the binary 6AI a alloy. As a consequence, compositional varia-tion studies were made through substitution of molybdenum for part of thevanadium in hopes of improving creep-rupture characteristics. Modifications ofthe basic 6A1-4V alloy were the following: 6A1-2V-l1o, 6A -2V-2No and 6A.-3V-Mo. Alloys were studied for either forging or sheet application and were

evaluated by tensile tests at room temperature, 6200, 8000 and 1020"F; creepand creep-rupturs tests at 8000 and 10207; and stability tests under stres attetpra os employed n sp testins..

Previous work on the 8A1 a alloy had shown that the ductility of the alloyhinged upon the annealing treatment given it. (5) Heat treating above 14700Ffollowed by water quenching resulted in excellent ductility properties, whereastreating in the temperature range 9300 to 129097 resulted in loss of ductilitywith complete abrittlement occurring upon heat treating at 10200F. In an4

WADC TR 54-278 Pt 3 1

4

effort to increase the aluminua content tolerable in Ti-Al binaiy alloys and toindicate continuity of properties with reference to composition and temperatmurrange where ductility can be expected, alloys of 6 and 1OAl were studied. An81 alloy prepared with iodide titanium was used to gain infomation on theeffect of interstitial content on its embrittlement. Tensile testing, electronmicroscopy and electron or X-ray diffraction techniques were the principal toolsemployed.

Factors influencing ductility in the 7Al-3fo alloy also were investigated.

Aging studies of the 5Al-2Ag and 5Al-7Ag alloys initiated in previousresearch work(5) were continued and some ehanical properties were determined.

II. EXPERD(ENTAL PROCEDURE

A. Preparation of Alloys

1. Materials

With the exception of some Ti-8AI material which was prepared with iodidecrystal bar, test ingots were made with sponge titanium of 120 BHN quality.The iodide titanium was vacuum annealed at pressures of less than 1 micron for4 hours at 1560"F to remove hydrogen prior to the melting operation. The spongems dried for 3 hours at 2509F to remove moisture which may have been picked upduring storage.

The source and major impurities of the high purity elemental additionsused are recorded in Table 1. The high melting additions, chromium, molybdenum,vanadium and zirconium, were charged in master alloys as particles -lA in. +16 mesh in size. The compositions of these master alloys are given in Table 2.

2. Melting Technique

Both non-consumable and constuable electrode arc melting furnaces wereemployed. The non-consmable electrode furnace was used in instances whereonly uall ingots were required. In such cases the ingots were inverted andremelted several times to insure homogeneity. The 6AI(2), iodide 8A1, 5Al-2Agand 5Al-7Ag alloys were processed in this manner, the binary 6 and 8A1 alloysas 200 gram buttons and the ternary alloys as 100 or 200 gram buttons. Allother A...... alloys, following initial mlting In thefurnace and forging to bar stock, were remelted in the consumable electrodefurnace.

3. Forging Practice

An electric furnace was used for heating ingots to temperature. The 6A1(2)alloy ws forged to 1/2 in. roind at 19O0"F. The $Al alloy was forged to

WADc TR 54-278 Pt 3 2N

oq 0

C! 0%.0.0 U 0

-% I.A * 0o 8

43 ON 0

.33 *Cj

"4 C3 0U-

08 03. 1^'0 00

oN O ~ H N Ig s C -

0%0

0 U%

0 0 0%j 0% R00

0) 0

r44T

GH

o co H c

.I0 oil9

f04

0

wA~ R54-78P 3 3

TABLE 2

MASTER ALUJMS

Composition Comments

Ti-6OCr Melting point below that of titanium; readily crushed.

Ti-31Mo Connercially available. Added as chips. Improvement overunalloyed molybdenum. Double melting necessary to ensurehomogeneity.

Mo-22AI Melting point in vicinity of that of titanium; readily crushed.

V-lhAl Commercially available.

Zr-lOA1 Melting point below that of titanium; readily crushed.

5

S WADC TE %I-278 Pt 3 i

1/2 in. diameter bar stock at 2100F. The Ti-Al-Ag alloys were forged to 3/8in. round at 1750*F. The double melted. 1OA alloy was forged initially toelectrode dimensions at 21009F and finish forged following the second meltingoperation at 20500F.

All other experimental alloys were initially forged at 19006F to 1 in.diameter electrodes. Following remelting, the alloys being investigated asforging materials were forged to 1/2 in. diameter rods in the range 1900" to1800F with the lower temperature being employed as the material reached finaldimensions. The three alloys being studied for sheet applicability were forgedin the 1900*-18OF range to 1/2 in. plate and hot rolled to 0.075 in. sheetat 1750*F. Throughout this report forging temperature is defined as the tm-perature of the furnace.

4. Chemical Analysis

Results of chemical analyses on representative samples of each alloy arepresented in Table 3. Hydrogen contents were determined by the hot extractionmethod, except where oxygen content was also determined. In this case thevacuum fusion method was used. In most instances the agreement between analyzedand nominal composition was quite good; as a consequence, each alloy is referredto in the text by its nominal composition.

B. Age Hardening Studies

Age hardening characteristics of the 5Al-2Ag and 5Al-7Ag alloys were evalu-ated in the teuperature range 800* to 1020*F. Samples 318 in. in diameter by1A in. sectioned from forged bar stock were used. All samples were given aninitial 24 hour solution annealing treatment in evacuated Vycor bulbs at 15600F.Aging heat treatments were also carried out in evacuated bulbs. Samples werewater quenched following the solution anneal and the aging heat treatments.

C. Tensile Testing

Forging alloys were tested as standard 3 by 0.252 in. diameter specimens.The sheet alloys were tested as 0.060 in. sheet. Sheet specimens were 7 in.overall with 1/2 x 2 in. reduced section.

A Baldwin-Southwark 60,000 pound capacity tensile machine, equipped withan autographic stress-strain recorder, was used for the tensile testing phaseof this investigation. Tests on the forging and sheet alloys were run at aloading rate of 1000 lbs/min, which amounts to an approxuimate strain rate of0.001 in/in/min at room temperature over the interval in hich the test pieceis deforming elastically. Beyond the proportional limit the strain rate con-tinual37 increases. Tests an the binary Ti-Al alloy. were made at loading ratesof 600 and W0o lbs/b.n.

A Baldwin-Lima-Hamilton separable, microformer-type extensomter wasemployed to give plots of load versus strain, from which 0.2% offset yieldstrength and modulus of elasticity values were calculated. This extensometercan be used for both ambient and elevated temperature testing, and can beadapted for either bar or sheet specimens by a simple exchange of extensionarms.

WADC TR 54-278 Pt 3\\

TABLE 3

CHMICAL ANALSIS OF EXPE RWAL AWlOYS

Nominal Composition Analysed Compoition Hydrogn Content, pm

6A1 (1) 5.9tal 233

6A1 (2) 5.77A.3 89

SAJ. (iodide) 8.239 8.35Al 100

1OA. 9 -096A3. 135

5Al-7Ag 5.32A-6.24, 6.88Ag

6Al-3(o 6.14Al'-.046 126.

6A1-2?4o-IMn 6-37AI-2.14wo-1.134n 161

WA-2N-lMr 6.lOAl-2.O5Mo-0.97Cr 153

6AI-2Mo-1LV 6.23A1-2.gomo-o.90v 189

6A1-214o-Lmn-lv 6.54il-.1cqo-.o2qn-o.9ov 221

6A1-Jmo-imn-lCr 6. 16A1-1 .ooo-1. o7mn- . 00Cr 200

6A1-3(o-IiSn 6-11A1-3.29Ko-3.%6Sn 129

6A1-3Ho-ljZr 6.05Al-3.23Mo-3.9lZr 170

6A1.3(o-2Sn-22r 5. 82A1L-3. 35Mo-1. 98Sn-2. 22Zr 235

10A1-3No-Iiar 10 .01AL-2 9lMo-4..16Zr

6A1-2 v-imo 6. 30A1-2 . b-i. loMo 73

6n -2a 6.19a1-1.q8v-2.O7o 7

6A1-37-INo 6.37A1-3OOV-1.O9Wo 106

WADC TB 5-278 Pt 3 6

For tests in the range -h0" to 570F the lower extension am of the tensiliassembly was equipped with a metal Dewar flask. Baths of various liquid mediawere used to maintain the test piece at the test temperature. The fluid mediaemployed were the following:

-hj07 - dr7 ice and acetone.327 - ice and water.

760 to 2127 - water.2120 to 48Oor - mineral oil.

A 16 in. long by 2-1/ in. I.D. Marshall furnace, regulated by a Wheeleocontroller, was used in the elevated temperature test work. Chromel-alumelthermocouples attached to the shoulders of the test piece indicated test tem-perature, which was held within ±%F over the test section.

D. Creep and Creep-Rupture Testing

Eleven lever-arm creep test stands and five stress aging units were emploin the program. The forging alloy creep specnimb was a modified version of thestandard 0.252 in. diameter by 1.25 in. reduced section tensile specimen. The,modification consisted of machining to much closer tolerances so that the effee-tire reduced section was not more than about h% of the 1.25 in. dimension.Sheet specimens were identical to those employed in tensile testing. Threerecording chromel-alumel thermocouples were attached to each creep specimen bymeans of Inconel clips. Temperature was maintained within ±20F over the testsection. The extensoeter used with the creep test stands for indication ofextension of bar specimens was not readily adaptable for testing of sheet.Consequentl., in the case of the sheet aloys, total extension was obtained onunbroken specimens by accurate optical measurements; extension with time datawere not obtainable.

E. Stability Testing

Unbroken test specimens creep tested at 8000 and 1020OF were tensile testedat room temperature to determine the influence of exposure at temperature understress application. Specimens which maintain ductility values 9Ciparable totheir unexposed counterparts are termed stable.

F. Impact Testing

Standard Charpy V-notch impact specimens were tested in the temperaturerange 32" to 212?F using a Riehle machine.

0. Bend Testing

Guidedbend teits'were made on smples of the three sheet modificationsof the basic 6A1-4V composition. A Di-Acro unit was employed for this testwork.

AD TR 5h-278 Pt 3 7

III. RESULTS AND DISCUSSION

A. Ag Hardening of Ti-5AI-2Ag and Ti-5AI-7Ag Alloys

Following the 24 hour solution anneal at 1%600F, samples of each alloywere aged at 800*, 840*, 930* and 10200F for times of 15 and 30 minutes, 1, 5,24 and 48 hours. The aging results are reorded in Table 4 and diagrmatioalypresented as Figure 1. Over the temperature range investigated the alloysexhibited fairly typical aging response. At 8000F the 5Al-2Ag alloy reachedhigher peak hardness than the 5Al-7Ag alloy. The 5Al-2Ag alloy--although notindicating an overaging tendency at 800F for times up to148 hours--overages atshorter times with rise in temperature. It appears frm the curves at 8000 and840OF that the 5A1-7Ag alloy overages quite rapidly. Disregarding the smallaging peaks for the 5Al-7Ag alloy at the lower temperatures, the resemblancebetween the two alloys is quite marked at 8400F and above.

Microstructures of the two materials wre similar. Representative struc-tures of the 5AI-7Ag alloy are shaor as Figures 2 and 3, which Indicatecopious precipitate occurring within the grains and at the grain boundaries.

Tensile specimens of the 5AI-2Ag alloy were aged for 24 hours at 8006Ffollowing the solution aneal, based on Indications that this was the optimumheat treatment for maximum hardening response. Tensile properties at roomtemperature were the following:

Ultimate YieldxTensile Strength

Specimen Strength (0.2% Offset) Reduction ElongationNo. psi psi in Area, % %

(1) 125,000 105,000 41.5 15.0(2) 132,000 nl,1,5 20.5 12.0

A 5Al-7Ag alloy tensile specimen in the solution annealed condition: 1560"F-2 hours-water quenched, had the following properties: UTS - 126,300 psi,YS - 100,O00 pi, RA - 27%, El -12%. Aging for 48 hours at 9300F followingthe solution annealing treatment gave the following: UTS - 136,600 psi,YS - n7,200 psi, RA - 20%, E1 - 13%.

Creep-rupture tests made on the 5Al-2Ag alloy (H.T.: 1560"F-24 hrs-WQ,800F-24 hrs-AC) and the 5A1-7Ag alloy (H.T.: 1560"F-2 hrs-1Q, 930°F-48hrs-AC) at 800"F indicated similar 100 and 500 hour rupture strengths of68,000 apd 614,000 psi, zespectively; creep stmngths to produce rates of0.2%/hou aWd 0.0106hou were 72,000 and, 63,000 psi,, respectily; see Figmz 14.These alloys, based on the above values, are decidedly inferior to the binary6AI alloy.

B. Results for Forging Alloys

1. Tensile Properties

Tensile characteristics of the alloys under investigation as forgingmaterials were determined at room temperature, 620, 800" and 1020. Duplicate

WADC TR 54-278 Pt 3 8N

TABL3 4i

AGING CEARACTERISTICS OF Ti-Al-Ag ALLOYS

Aging Diamond Pyramid Hardness (20 Kg. Load)Temp.

Alloy* "F 0 15' 30' 1 hr. 5 hrs. 24 hrs. 48 hrs.

Ti-5Al-2Ag 800 275 280 282 290 292 330 328

840 275 277 286 283 290 323 274

930 275 274 282 284 304 289 279

1020 275 302 287 306 308 271 282

Ti-5A1-?Ag 800 295 325 322 305 298 314 298

840 295 315 291 292, 295 306 297

930 295 301 304 295 283 309 283

1020 .295 320 308 302 334 309 33l

• Solution treated: 15600F-24 hrs-WQ.

WADC Tt 5-278 Pt 3 9

N,,

I _____________________________ _____________350____________ _____________________________

0.0

w 300

844

350

A. I30 OF0

5 OAGINGiTIME INI fill

FIG. 0- G 0AD N N C H R CT R S IC F IA I-2%A g

AND TI-5%AI-7%Ag ALLOYS

N WADC TR 54278 Pt 3 10

-* #-40 w

Ar/

W

N eg. No. 1140~8 X 250

Fig. 2

Ti-5Al-7Ag: 1560OF-24~ hrs-WiQ, 930*F-48 hrs4JQ.

J

Neg. No. 11W~9 X 250Fig. 3

Ti-5A1-7Ag: 150 *F-2i4 hre-A, 9307-5W0 hra-WQ.

Etohmnt: 20% Up, 20% HNO3., balanc, glycerine tollowsd by 2% OF in watr.

wADC TR 54s-278 Pt 3 11

TI- SAI-2Ag

CREEP RATE

000

5 0 -RUPTUREcn

10,

110 100 1000RUPTURE TIME IN HOURS

MINIMUM CREEP RATEI IN/IN/HR

FIG. 4- STRESS-RUPTURE AND MINIMUM CREEP RATECURVES FOR Ti-5%At-2%Ag ALLOY AT 800 OF

WADC TR 54-278 Pt 3 12

tests were made at each temperature on spscimsns given the following heat treat-ment: l4701-6 houre-air cool, l020"F-24 hours-air cool. Test results arepresented in Tables 5 to 9 and plotted as tensile property-temperatwe orvesin Figures 5 to 15.

a. Room Temperature Properties

The alloy base material, Ti-6Al-34, showed moderate room temperaturestrength (UTS - 145,000 psi, YS - 127,500 psi) coupled with acceptable ductility(RA- 28.5%, El - 12.5 %). The structure of this alloy, as shomrin Figure 16,was typical Wdmnstitten L-P configuration. This structure was duplicated in •most of the other alloys.

Replacement of 1% of molybdenm with mangaese produced the Ti-A-o-MImodification, which showed a 20,000 psi increase in ultimate strength with a30,000 psi Increase in yield strength. Strengthening we8 obtained at theexpense of ductility, however, which dropped to values of 8.5% RA and 3.5% El.Hydrogen, present to the extent of 161 ppa, could possibly account for thispoor showing, but other of the alloys in this group -with higher b&rogbs contentsshowed no ductility impairment. To check the influence of hydrogen, additionalspecimens were vacuam annealed to reduce hydropn content to About 20 ppm.Results of tensile tests of vameu annealed specimens are Inolvded as Table 6.The 6AI-2Mo-In alloy was highly ductile as vacunm annealed; the other lloyswere not sig ificantly affected ty the treatment. The structure of the 6-fo-INn alloy shom in Figure 17 was significantly different than that of the6AI-YO alloy, consisting of patches of Widmanstitton a-p surrounded by oquazwdgrains of a.;

Substitution of 1% of either chrmium or vanadim for mol7bdenua offeredno Improvement in ultimate strength, although vanadium raised the yield strengthto 135,000 psi. The ductility of the chromim-bearing alloy was similar to thebase material, whereas the vanadium-bearing alloy showed improvement (RAm • 0,El a 36%). Both alloys were structurally similar to the 6A1-3o alloy.

Further P eoplezing to give the alloys Ti-6A-lMn-IV at Ti-6Al-o-1Cr resulted in an increase in ultimate strength to a level in the vicinity of155,000 psi. The former alloy had a much higher yield strength than the basealloy (lJ5,00 psi compared to 127,400 psi), but also had marked*y better due-tility (RA.-lWO, El a 17.5%). The ductility values of the Utter alloy wslightly superior to those of the ternary base material. Figure 18 shows thestructure of the vanadium-bearing alloy which was predominantly Winmst~ttena-P with islands of a. The structure of the 6Al-No-%1fn-lCr aloy was similarto that of the 6Al-3Ko alloy, but coarser.

O09fte a oaqlmd allcu, 6A:L-jg..ir vus only aubwmt2 0"' Volsthan the 6A:L03(o base an the basis Of b h u ,te end yiel p h.Its ductility was nearly identical to the base composition. The A1-3(o-Ifnand 6Al-o-$n-2Zr alloys ndicated an increase in ultimate strength to 155,000psi. The yield strength of the 4Sn bearing alloy at 133,000 psi was slghtlyhigher than that of the 6A -l3o alloy. Th yield strength of the al withaddWons of 2Sn and 2Zr was stil higher it 138,000 pai. Both of these aloysMs campas In du o lities to the alloy bae. All thre alloys wwo: iidw

structurally to the 64.-3Nt alloy.

WADC TR 54-278 Pt 3 13

0) fV '00 c'O co- 0% ON ' ICJ N 0 C&~ V

r4. a; t- O% tC. V Ci -Ci 0! RO le'i Oh

UN, 10u )i0 UA .0 0 Us '0f um it um.0

%(\I I^ If 0 I I 0. 1^ U-c6 a; 1 .- I' c6 4c 4, A I II~ t N I .v I %t

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00

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II 4VJ' I -I HO r I H rI-f H H'

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43

E- a H H H rq H ~ H N4 rq H

C-4

wADC TR 54~-278 Pt 3 15

.1 4C 0"0 C' 1 U0.

00

r) 1* * N % N 0

0

44

oiliVA Vbi V% CO

WADC~~~ 0R5-7 t31

It ~ xm G U-%I C -cu 0 , o ri % rnMe

-4rrr-4-4ON r.4 1ll rorl

VA .4~c

IAV\ I O 00f 0\t' 01JO\ U 00 0001^LI

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Eld 0

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464

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1.4

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m~Q 0% 1A %A -ro O % 00 ,a .g o 9- C

41.

WADC TE 5-278 Pt 3 19

TEST TEMPERATURE, OF10210 390 570 750 930 1110

Ti7 -SAl- 3Mo

140

z020

zI.-0(I)

0-00

0-

wit4

80

60

2 0

a

010 100 200 300 400 500 600

TEST TEMPERATURE, Oc

FIG. 5 - TENSILE PROPERTIES OF TI-6%AI-3%M. ALLOY

WADC TI. 544378 Pt 3 30

TEST TEMPERATURE, OF210 31J jj7T509 1110

140

Ix 12

w I_ _ _ _

S 0

sJo

0

6 0 __ _ _ _ _ _

z

0 Z

P 20

00

0 100 200 300 400 500 600J TEST TEMPERATURE, OC

FIG. 6 - TENSILE PROPERTIES OF Ti-6%AI-2%Mo-I%Mn ALLOY

WADC TR 54-278 Pt 3 21

TEST TEMPERATURE, OF

160 210 390 570 750 930 1110

F AI2Mo-ICr

140-

0-

w

0-

0 RA__ _ __ _

40

60

00

040J

00

0 100 200 300 400 500 600TEST TEMPERATURE, Oc

FIG. 7 - TENSILE PROPERTIES OF Ti-6%AI-2%Mo-I%Cr ALLOY

WADC TR 54-278 Pt 3 22

TEST TEMPERATURE, OF

IS0 210 390 570 750 930 1110

Ti2AMoM-IV_

140

z00

120__ _ _

0%

0-

wIYS

40

z80Z

60

20

00

0 100 200 300 400 500 socTEST TEMPERATURE.'C

FIG. 8 -TENSILE PROPERTIES OF TI-6%AI-2%MO-I%V ALLOY

WADC TR 54-278 Pt 3 23

TEST TEMPERATURE, OF

ISO 210 390 570 750 930 1110

140

I-0

Ilg.. 12C

_j0-

t00 ow .

so-I- 80

-j

60

wRAo0440

0zp oz

02-0I-u 20

00

f 0 O00 200 300 400 500 600

TEST TEMPERATURE, °C

FIG. 9 -TENSILE PROPERTIES OF Ti-6%AI -I%Mo-IMn-I%V ALLOY

WADC TR 54-278 Pt 3 24

TEST TEMPERATURE, OF

IS0 210 390 570 750 930 1111

TI -AI -IMo- I M-Ij

140

zw 120

I-T

0-jU

100o

0

0800-

00

w 40

z0OZ

2- 20 ___

00

00 100 200 300 400 500 60CTEST TEMPERATURE, OC

FIG. 10- TENSILE PROPERTIES OF Ti-6%AI I%MO-I%Mn-I%Cr ALLO)

NWADC TR 54-278 Pt 3 25

TEST TEMPERATURE, OF

Io210 390 570 750 930 1110

TI-3AIo4L

140

wIx

I-120

0 UT

0

0-

YS

0

40

02Z-0

IE-L1D20

00_j W

010 100 200 300 400 500 600

TEST TEMPERATURE, Oc

FIG. 11 - TENSILE PROPERTIES OF T - 6%AI-3%Mo-4%Sn ALLOY

WADC TR 54-278 Pt 3 26

TEST TEMPERATURE, OF

10210 390 570 750 930 1110

Ti -6A - 3Mo -4Zr

z140-

zw

to.

0

0.-

100

80-

60-

~40-

0

z 20 EL0oz

0

w m 01I0 100 200 300 400 500 60(

TEST TEMPERATURE, 0 C

FIG. 12 -TENSILE PROPERTIES OF Ti-6%AI-3%Mo-4%Zr ALLOY

NWADCTR 54-278 Pt3 27

TEST TEMPERATURE. OF

ISO 210 390 570 750 930 1110

Ti - I0A-3Mo-4Zr

z - Sw

I-.S

1-004

Io0-

S 120 _ __ _

a

60 _ _ _ _

0

4cc40

04

z202Z

420

00J w

010 100 200 300 400 500 600

TEST TEMPERATURE, OC

FIG. 13- TENSILE PROPERTIES OF TI-IO%AI-3%/Mo-4%Zr ALLOY

WADC TR 544278 Pt 3 28

TEST TEMPERATURE, OF

10210 390 570 750 930 1110

Ti- AI - 3Mo -25 -2Zrj

14

z

_120

0 10

so.

600-0

w R

0 4

80Z

04 2

-f 0 0 003040-0 oTETTEPRAUE,-

FI.10ESL0RPRISO T-% I3 M-% n2 Z LO

20CT 427 t32

TEST 'TEMPERATURE, OF

60210 390 570 750 930 1110

z

0

G00

0-

w

0-

80

0 0U 2

wcE

0.z-00 003040 0 o

I.--TET EMERTUE

FI.1 ESL POETE FT-6%U3 C LO

00CT 427 t33

AN I~

h w Neg. No. 11055 X 750~ ~ Fig. 16

Ti-6A1-3Mo, 1I470'F-6 hra-AC,Ol- Wj1020OF-24 hra-Ac. a~+

Neg. No. 3U0I4 X 750

Fig. 17Ti-6A-2mo-mn, 1470OF-6 hro-AC,102OF-24 hrs-AC. a + P

Neg. No. 11049 X 75CFig. 18

Ti-6A-lfo-Mn-1V, 110*7-6 bra-AC,1020O7-24 hra-AC. a. +

!tdiut: 60 cc glyerine, 20 cc HNO,, 20 ca HF.

WADC TR 54-278 Pt3 31

A 1OAl-3Ko-IZr alloy, which was unintentionally prepared, was brittle atroom temperature as indicated by failure in the threaded sections of the testpieces. Structurewise, this material was very coarse Widmansttten.

The Ti-6Al-3Cu alloy, which is an a + compound type material, had roomtemperature strength properties very similar to those of the 6Al-3(o alloy, butshowed very poor ductility (one specimen failed in the threads, while the othershowed an RA of 5.5% and an El of 5%). The a + Ti 2 Cu structure for thesespecimens is shown in Figure 19. It was suspected that the low ductility ofthe 6A1-3Cu specimens was due to this WidmanstItten type structure. Much betterductility would be expected from an equiaued structure. Therefore, the 6A1-3Cuspecimens represented in Table 6 were finish forged at a lower temperature thanthat used to produce the structure shown in Figure 19a, i.e., 1750*F comparedto 18006F. However, for unknown reasons only one of these specimens had thedesired structure. This was a ductile specimen and its structure is shown inFigure 19b. The other specimens had WidmanstAtten type structures.

b. Elevated Temperature Properties

At 620OF and above all the alloys had acceptable ductility values. Ingeneral, the ductilities increased with temperature. The is complemed alloyswere somewhat stiffer than the 6A1-3Ko and did not show as great a change withtemperature as did the P compled alloys.

The IOAl-3Mo-4Zr alloy was significantly superior on the basis of strengthconsiderations at 620OF to the other materials, having an ultimate strength atthis temperature of 155,0OK. psi and a yield strength of 130,000 psi. Thesuperior performance of this alloy was maintained at 800* and 1020*F and is areflection of the influence of aluminum in elevated temperature strengthening.

At 620*F on the 6A1 level the 6A1-2Mo-lMn P complezed alloy and two a com-plezed alloys, 6A1-3Mo-Wn and 6A1-3Mo-2Sn-2Zr, were strongest at 122,000 psi.This stress level was approximate3y 10,000 psi higher than that found in the6A1-3Mo alloy. With the exception of the 6A1-2Mo-lV alloy, which had anultimate of 106,000 psi, the other alloys were close to the base at 131,000 psi.

At 8009F the 6A1-2Mo-IMn alloy with an ultimate strength of 117,O00 psiagain showed the highest strength, followed by the a complexed 6AI-3(o-Wsn and6Al-3Ko-2Sn-2Zr alloys at 3-14,000 psi. The remaining a omplexed alloy, 6AI-34o-iiZr, and the other P complexed alloys together with the base material weregrouped between 100,000 and 106,000 psi. The 6AI-3Cu alloy was quite strong atthis temperature with an ultimate strength of 116,000 psi.

The strength of the 6A1-34o alloy dropped precipitously at 1020OF to about82,000 psi. The alloys with chrmim, mananese or mnadim nd th ea And6A-'3o-bZr alloy weregrouped around 90,000 psi. The a oamplexd 6h1-3o- Snand 6A1-3Mo-2Sn-2Zr alloys showed their elevated temperature potential withultimate strengths at 106,000 psi. The copper-bearing alloy was strongest atthis temperature with an ultimate of 109,000 psi.

WADC TR 54-278 Pt 3 32

4

W~~

Neg. No. 11053 X 750a

1I47OWF-6 bra-AC, 1020*F-24 hrs-AC.

Neg. No. 11582 X 75D

b

VA 1560*F-4 bra-Ac, 10209F-24 bra-Ac.

Fig. 19 -Ti-6A-3Cu. a + T12Cu-

WADC TR 5-278 Pt 3 33

2. Creep-Rupture Properties

The forging alloys were oreep tested at 800"F and creep-ruptore tested at10207F. Three creep tests were made at 800*F for each alloy. In most instancesone test was conducted at a stress calculated to cause failure within 300 hours.The remaining tests were run at lower stress levels and were terminated withoutfailure at 300 hours. Creep-rupture tests at 10206F were made to produce suf-ficient data points for reliable extrapolation of the stress vs. rupture timecurves to 500 hour rupture lives and to enable preparation of stress vs. minuucreep rate curves. These curves are presented in Figures 20 to 29. Test dataare contained in Tables 10 and 11. A osary of the rupture strength and creeprate data is also included as Table 12.

On the basis of stress-rupture strength at 10200F weakening resulted fromcomplexing the P phase by substitution of chromum, manganese or vanadium forpart of the molybdenum of the 6A1-3fo base. The 6Al.-.Ko-lCr alloy most closelyapproached the base material, having 100 and 500 hour rupture strengths of32,000 and 22,000 psi, respectively, as ompared to 40,000 and 33,000 psi forthe 6A1-3Mo alloy. The a complexed alloys, 6Al-3mo-hsn, 6Al-3o-hzr, and 6A1-3Ko-2Sn-2Zr, were sipnificantly superior at a rupture life of 100 hours, havingrupture strengths of 48,000, hS,000 and 52,000 psi, respectively. However, theslopes of the stress-rupture curves were steeper than those for the 6Al-3K(oalloy. As a consequence, at 500 hour rupture life -the a complexed alloys withstrengths of 35,500, 34,50o and 36,50O psi, as compared to 33,000 psi for thealloy base, were not as superior as indicated by the shorter time values.Extrapolation of the stress-rupture curves suggests that the 6A1-3Ko alloywould be stronger at 1000 hours rupture and beyond.

Although only limited testing was done on the IOAl-3(o-Zr alloy at 10209Fand only two tests actually were rmn to failure, there vsre definite indicationsthat this alloy has excellent elevated temperature properties. The stress toeffect failure in 500 hours was not accurately determined but should be inexcess of 50,000 pei.

The 6A1-3Cu, + coapound, alloy vas slightly inferior to the 6A1-34o alloyat 10200F.

Comparison of the stress vs. minima creep rate curves showed that at 800"7the 6Al-lo-ln-lCr alloy had a stress of 73,000 psi in comon with the basealloy for a creep rate of 0.001%/hour, but was inferior at higher creep rates.At this temperature none of the other P oomplexed alloys approached the creepresistance of the 6A-3fo alloy. The a completed alloys, 6A1-3Mo-hSn and6A-3(o-2Sn-2Zr, together with the 6A1-30u alloy, had nearly identical creepproperties and were aunaiosatlmo r4e nt, tM, the 61-3NOK alloy (85 ,000psi at a creep rate of 0.001%/hou as compared to 73,000 psi for the base).The 6Al-3(o-IZr alloy very closely approached the other a complexed alloys butwas 2000 psi inferior (83,000 psi at a creep rate of 0.001%/hour). The lOAl-3Ko-hZr alloy had a creep strength of 97,000 psi for a 0.001%/hour rate.

At 1020"F the P complezed alloys were markedly inferior to the 6A1-3foalloy, while the a oamplexzd alloys peve the best performance. Stresses of49,000 and 28,000 psi correspond to creep rates of 0.1%/hour. and 0.0%l/hor,

WADC TR 54-278 Pt 3 34

Ti -6AI-3Mo

IL 1000020-50-

10,1020 0F

1 10 100 1

RUPTURE TIME IN HOURS

FIG. 20 -STRESS-RUPTURE CURVE FOR ALLOYTi- 6%AI-3%Mo

WADC TR 54-278 Pt 3 35

TIS-Z-2m@- I Mn.

50

500

10 1 l lt-I iA.

U)I

-50

Iac

Ti - 6A1-ZMo -1 V

100

8QOOF

50-

1 10 100 10RUPTURE TIME IN HOURS

FIG. 21 -STRESS-RUPTURE CURVES FOR ALLOYSTi-6%A.1-2%Mo, WITH I%Mn, I%Cr OR t%V

WADC T& 54-278 Pt 3 36

I if i I1 1 I I I off V I I I off1

TI-SAI Mo-lMn-IV

100

800 OF

50-

020 OF

10 1 a i i a l m I ~i ll

- TIwSI-IMo-IMn-ICr

0

10

- 50100O

Oo)0

U)A10L

TI-6%- I -% 0t% SIK 1%Vo-OR

too N T-%I3%04S

W ~ ~ ~ ~ ~ ~ 0 AD 05F 28 t33

1000O

50-

10

oo005

cn 1020 OFLi

co

Ti-6A1 - 3Cu

100

50-

1 10 100 1000RUPTURE TIME IN HOURS

FIG. 23- STRESS-RUPTURE CURVES FOR ALLOYSTi-6%AI-3%MO-4% Zr,Ti-6%AI-3%MO-2%Sn-2%Zr ANDTI- 6%AI -3% Cu

WADC TR 54-278 Pt 3 38

00

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WADC R 54278 P 3 3

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100

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iSd 0001 'SS3vJ.s

WADC TB. 54-Z78 Pt 3 40

0

fill1 I I 1 1 I *-

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WADC TR 54-478 Pt 3 42

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as 000 a4~l

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WADC TE 54~-278 Pt 3 4

CJ

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WADC T 54-278 Pt3 4

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WADC'TR 5-278 Pt 3 4

Al00 038 '00 000 00 0300 00000

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wADc TR 54-278 Pt 3 4

respectively, .for the 6A1-30o alloy. At the highe rats the 6A1-3No-afn and6A1-3Ko-laZr alloys wer similar at a stress of 482000 psis but weae -m resistantat the o.01%Aiouir rate at a stress of 37,000 Psi.

The 6A1-3(o-23n-2Zr alloy was most amep resistant, requiring stresses of57.9000 and 400, 000 psi for creep rates of 0. 2%Aaour andO .0l%Aiour,, respectively.It is interesting to note that this alloy at 10207F is vealar than the 7A1-3(o(1N) alloy(3 on a stress-orpture coprsn(2,000O and 360500 psi for 100 and5DQ hour life, respectively, for the 6A1-V31o-'2Sn-2Zr alloy, as compared tostresses for the 7A1-34o alloy of 51.,50o and l6,00 psi) but in identical onthe basis of creep resistance (57,,000 and Wpm00 psi for rates of 0.2%/'iou and0.0i%Aour) respectively).

The 1OA-3qo-IjZr alloy at 10206F had a creep strength for a rate of 0.0l%/hour of approximately 51,000 psi.

The 6A1-3Cu alloy had inferior wreep resistance to the 6A1-3Mo alloy at acreep rate of O.1%waio (1.,000 psi) but was comaal at a rate of 0.01%Aoour(28,000 psi).

It is apparent from the above findings that In oomplein of the a phasewith tin and siroonim lies the hope for aloy imrvmn.The elevated tem-perature performance of the IDAl-3fo-Wjr alloy is noteworthy. Its potentialapplicability is hampered by its poor room temperature properties, however.Because of the potential eiibited by this alloy and the-6A1-3Cu alloy, anattempt to improve their room temperature properties would be amply Justified.

3. Stability Characteristics

The results~of room temperature tensile tests conducted on specimenstested initially in creep for periods of 300 or 1000 hours are soaried InTable 13. The base 6A1-36o alloy vas not sipmificantly affected by the exposurconditions. The 6Al-2!o-In alloy which was low In ductility initially wasalso unaffected. (One specim was brittle as a result of an Imperfection.)Two specimens of the 6A-2H-l0r alloy (800-70,000 pei-305.7 hours and 1020*7-20,000 psi-6W4i.2 hours) were noticeably influenced by stress-aging. The ladkof ductility In the specimen exposed'at 102007 is perhaps explainable by thefact that this test piece extended 13% in creep. Tests on the 6A1..2Mo-lV alloywere variable; exposure at 8007 uinder a stress of 80,000 psi showed smullchange in properties, whereas exposure 'under a stress of 70,000 psi broughtabout a considerable drop in ductility.

The 6A1-lMo-lmn-lV and 6A1-I~o-]Mn-lCr alloys also lost ductility throughexposure. The 1V bearing alloy aged under a stress of 80,000 psi at 800OF wascompletely brittle.

The a oompexd alloys were also seriously affected by stress-aging.Again, the loss in ductility was not systemnatic: and no attempt is made toexplain the behavior.

As in the case of unexposed specimens, 6A1-3Cu alloy specimens which werexposed to stress at O00 and 1020OF had poor or good ductility dpnigUPonthe microstraoture. The two opeclasms ith good ductilitre euiaid; whilethe others were Widmanstitten.

WADC TE 1.-278 Pt' 30

,6t~ClA 0V6%O 0V -4 A 0001fl4 c

00 060,oom~ -uoo u 0 0 0 U% I%vt Cj 0N -O CY 1A N 0 0% HO Ha

0% 4J6% % 6 0.r~4 for rI4N t f

U\ 0 0 0 0 90 00 000 060 0 MI *\\U 00

UNV\ 1U',\O0 0000 000 UA1A1J\ 00 0000c * c6 0 c6 A A A -4 r0 j _0 A. A 0 A

r-: - iRR4% %0k QUU

.- 4-4 gw 60?\C- P ri OH P4 r--lo rr~-

00 000 00 4r6I *864: 6C

%0 *O to r0 t- v %0 0 0 0m m 0 0~ 0i 0%0 0. *q 0 0

00- 0000 000 004000 600c 0 -%A

9 R§ HR §9 R9 §0

10

r-4 er-4H 4C~j CY rI HH r-I e-4r4N -Ir-4 u-lu-H i H -H 4-CY

wADc TR 54-278 Pt3 5

00 0U%^0%^W' AW

00~~ 0V % 0 0V0W4A -4u%*** 00 **

.00 0000 0RRC Rc000

%f 0 0 UN~0l lJVUfit%A0 %A

%k 'O'OC4 cm CYA~

b~~ NN § §~JN

r.,~ 0%' 0 r 84 I4R ~ 00 0082 8o~'e

*q.I~~ * 0 . * * *0Q-' 00 00000 00 10000

~-N~rIl~lO~ 0 N1AO

All-' A rOI 9,:R 9 9I4

§-r r§4Nr *' .' § *

9J NcsE~ 0800\ ,cs 4 C8%'

43-% t N 4% 0

WCD TR 54-278 ft 3 5

C. Results for Sheet Allos

1. Tensile Properties

The tensile characteristics of three sheet alloys, Ti-6A1-2V-lMo, Ti-6A1-2V-2Mo, and Ti-6A1-3V-IMo, were determined at room temperature, 620", 800" and1O20'F. Duplicate tests were run at each test temperature. Specimens weregiven the sans heat treatment given the forging alloys, namely, 1470"F-6 hours-air cool, 1O20"F-24 hours-air cool. The test data are included as Table 14 andplotted in Figures 30 to 32.

Of the three modifications of the 6A1-hV basic composition there was littleto choose between the 6A1-2V-Mo and 6A1-3V-lMo compositions at the lower tem-peratures. The 214o bearing alloy was slight3y strnger at 10200F. The 6A1-2V-lo alloy was about 10,000 psi weaker than the other two modifications overthe entire test temperature range. On a ductility basis all three alloys werequite comparable. It is interesting to note that ductility did not changeappreciably with temperature.

2. Creep-Rupture Properties

Prior to the initiation of the sheet test program it was believed that theextensameters employed in creep testing of bar specimens could be modified forsimilar testing of sheet specimens. It was planned to attach the extensometerto the specimen holders and ompensate for the creep occurring outside the1 in. gage length. Because of excessive elongation of the test pieces at 1020%F,coupled with slippage of the specimen within the holders, the proposed creepmeasuring procedure was abandoned. As a consequence, extension with tim dataare not available for the shaet alloys. At 80007 the test specimens were gagemarked very accurately (1.0000 in. t 0.0002 in.) and in the instances whererupture did not occur the reported total elongatibn values are correct within.0.0002 in. Tests at 1020*F were run to failure and total elongation valuesfor these specimens and all specimens at 8008? which ruptured are accurate towithin +0.005 in. The data are recorded in Table l and curves of stress versusrupture life are presented for the three alloys at 1020"F in Figure 33.

On the basis of rupture strength at 1020"F and the limited total extensiondata at 800"F the 6Ai-2V-21o alloy appears to be the best modification of thethree. The 100 and 500 rupture strengths of this material, 24,000 and 15,000psi, respectively, are well below values for the 6A1-4V alloy for similar rup-ture lives, namely, 33,000 and 26,000 psi. The difference in strength levelsmay be tempered somewhat by the fact that the 6Al-V alloy was tested in barform. Test piece geometry probably plays a significant part in the wide spreadin strength.

3. Stability Characteristics

Specimens for stability studies were taken from unbroken test pieces reeptested at 800F. Only four such specimens were available. The 6Al-2V-lMoalloy tested: 800"?-70,000 psi-306.2 hours had an, ultimate room temperaturestrength of 136,000 psi as compared to an uwexpomsed ultimate of 11.,000 psi.DuctilitV was lowered from 11 to 8%. The 6A-2V-a1o alloy creep tested:

WADC TR 54-278 Pt 3 53

TABLE 14

TENSILE DATA FOR EXfENTAL SHEET ALLOYS

Ultimate YieldTensile Strength Modulus ofStrength (0.2% Offset) Elongation Elasticity DPH

Anloy* psi psi % (six 10-) (20 KgS. lam)

Rom T!m ature

6A1-2V-IMo 140,200 --- 11.0 -- 3301242,000 131,000 11.0 16.8 328

6A1-2V-2Mo 153,700 144000 11.0 16.8 357.51,.6oo) l44,ooo U.0 17.5 347

6A:L-3v-IMo 153,000 150,ooo a 12.9 350151,600 1149,50 3-1.0 15.6 359

620F

6A1-2V-3Mo 98,500 82,800 1.5 15.5102,200 85,500 11.5 11.7

6A1-2V-2Mo 113,000 97,700 12.5 12.1109,600 95,500 12.0 12.1

6A1-3V-Mo 113,5o00 10,ooo 13.0 11.5107,000 92,000 12.0 15.9

800"F

6A1-2V-MO 93,200 80,200 12.0 6.493,800 79,000 11.0 8.6

6A1-2V-N o 106,000 79,300 12.0 --102,000 85,700 12.5 6.8

6Al-3V-imo i.,ooo 76,40 10.0 7.2104,200 77,500 12.0 5.5

1020*F

6h.-2V-Xo 79,500 64, 0o 16.o 8.982,000 67,000 114.0 8.9

W61-2V-mo 90,000 74,14)o 16.5 7.893,200 83,700 13.5 6.6

6A1-3V-lMo 85,100 69,600 20.0 9.087,8oo 75, 400 19.5 --

Alloys beat treated: 2470-6 hru-aC, 102oF-24 hs-AC.a Failure occurred outside gage marks.

WADC TR 54-278 Pt 3 54

TEST TEMPERATURE, OF

210 390 570 750 930 1110

TI- 6A1-2V-IMo

140x

zw

1200-

00=100-

w

80-

-0J

80

6-0

0 E

z 20

0 ,

0o 100 200 i3010 400 5.00 600TEST TEMPERATURE, 0 C

FIG. 30-TENSILE PROPERTIES OF TI-6%AI-2%V-I%Mo ALLOY

WADC TU. 54-278 Pt 3 55

TEST TEMPERATURE, OF

10210 390 570 750 930 1 110

Ti-6AI-2V-2Mo

140

z

120-

00

w

IYS

so-

so-

0-

z 20-0 EL

0-0z

0

0 100 200 300 400 500 600TEST TEM PERATURE, 0C j

FIG. 31 -TENSILE PROPERTIES OF TI-6%AI-2%/V-2%Mo ALLOY

WADC TR 54-278 Pt 3 56

TEST TEMPERATURE, OF'

IS210 390 570 750 930 II

Ti-6A1-3V-lMo

140-

120

0

0

0--

20-

00

0 20 0 3040 0 00ETTMERTRO

FI.3 -ESL4RPRIS FT-% I3 VI M LO

0ACT -7 t3S

TABLE 15

COUP-DPTRE DATA FOR MU AUDO7 AT 8000 AND 1020*7

TotalTest Rupture MeaeurdTamp Stress Time Elongation

A3-oTy* (OF) (psi) (hrs) C(_,)

6Al-2V-IMo 800 80,000 201.6 0.30070,000 306.2a 0.0605

1020 4B,500 2.4 c30,000 33.0 0.59520,000 149 .6b 0.325

6A1-2V-9o 800 80,000 276.3 a60,000 319.a 0.0039

1020 50,00 9.4 0.41040o,000 19.8 0.47030,000 45.8 0.63020,000 248.6 0.800

6A1-3V-lMo 800 80,000 157.6 0.29570,000 3o4.8a 0.046260,000 345.8a 0.0231

1020 50,000 5.6 .5)Wo,ooo 4.9(?) 0.37530,000 27.0 0.82020,000 123.0 1.110

* Alloys heat treated: 14706F-6 his-AC, 10200F-24 bra-AC.

a Test stopped at tme indioated. No failure.

b Test stopped due to excessive elongation.c Failure occurred outside gage marks.

WADCT T 54-278 Pt 3 58

",,

0 6A1-2V-IMOA 6A - 2V -2'Mo

U) 0 6AI -3V- IMO

00

w

102-0 OF

10-1 10 100 1000

RUPTURE TIME IN HOURS

FIG. 33- STRESS-RUPTURE CURVES FOR ALLOYSTi-6%AI-Z%/V- 0 / AND 20/Mo ANDTi-6%AI-3%V-I%Mo AT 1020 OF

WADC TR 54-278 Pt 3 59

8000F-60,000 psi-3l-9.l hours had essentially the same ultimate strength (l%3p000psi) as the unexposed alloy specimen and comparable ductility (12% exposed vers=11% unexposed). Both exposed specimens of the 6A1-3V-lMo alloy failed in Imper-feet areas outside the gage marks. The ultiuates were around 157,00 psi.Because of the position of the fractures the ductilities are not as high asnormally expected. The specimen creep tested at a stress of 70,000 psi showed3% elongation between gage marks, while the specimen exposed under a stress of60,000 psi had 7% elongation between gage marks. Based on the latter value itwould appear that this alloy and the other modifications areireasonably stablematerials.

4. Bend Properti e

Guided bend tests consisting of bending 0.060 in. thick sheet specimensaround mandrels of varying radius were conducted on the alloys in the heattreated condition: 1470OF-6 hra-AC. The 6A1-2V-Mo and MA-2V-2No alloyspassed 3T; each failed trials at 2T at an angle of 700. The 6Al-3V-lHo passed2T.

D. Properties of Ti-7A1-3Ko Alloy

Under Contract No. AF 33(038 )-2286 a forging study was made of the Ti-7A1-3(o alloy (see WADC TR 54-278 Part 2). (4) Wide variations in tensile ductilityas expressed by reduction in area were obtained with differing thermal andworking history. In an effort to isolate the variables influencing ductilityin this alloy, samples representing each of the thermal and working historieswere analyzed for hydrogen. In addition, nitrogen and oxygen determinationswere made on a series of samples representing a study of the effect of theamount of reduction at the initial and finishing temperatures of forging. The'hydrogen analyses were made by the Materials Laboratory, Wright Air DevelopmentCenter. The hydrogen contents are plotted versus tensile reduction in area inFigure 34. There appears to be some correlation between hydrogen content andductility. However, the wide range in ductility for a given hydrogen levelsuggests that there are other factors operative. Both equiaxed and Widmanstittenstructures were observed. This gross structural difference does not appear tobe of significance.

The results of the hydrogen, nitrogen and oxygen analyses on samples fromthe study of working variables is given in Table 16. No single element gives acorrelation with ductility for both the odd numbered group of specimens (finishedat 1800*F) and the even numbered group (finished at 1650"F); nor does the totalatomic percentage of the elements H, N and 0 provide a correlation with ductility.

In a forging study of the 7Al--3No alloy referred to above, it uas foundthat In the two eases In wich it ms applied, the heat trmatbnt, 15OWF-4hours-air cool, 1020"F-24 hows-air cool, resulted in slmiLficantly better duc-tility than the heat treataent: 1610F-24 hours-air cool, 1020OF-24 hours-aircool., For these results refer to Table 25 of WADC TR 54-278 Pt 2. The resultsof another part of this forging study are sumaried in Figuret 35 (frm WADC Ti54-278 Pt 2). The objective of this part was to determine the effects offinishing teaperature and aount of reduction at the finishing temperature ontensile ductility. Finishing temperatures were 18000 and 650"F. Amontate o

WADC TR 54-278 Pt 3 60

0w

zw W

0 w -

o 04o- a w0

J -c

0 0 LhJ0 . -

IEZ z zwiIr

w

Zq

w-

9wC*- o;4 0

0 w

w

N 0

o0 z

0Iw2 0

VWdd IJN31NOD N300MOAH

WADC TR 54-278 Pt 3 61

0

19 § § 0 AH

*0 0 0000

0; 0 0 0 .

oh 0% 0 0

*00 0 0 0 0

0% 0% t O CD te-l e- N ri

00 C m0um 0 4 0z 0 0 O

('1c 41O t40% '0

* 4 04 0 04 04

S 0 '0 1- 4 (A 0

P144 43P1

CO I .. .Og0 i

+).

43 _It X 'sZ 410 CZ

* l I I- V% I r %

ran ~ :1 R. Hg, mU

WADC TR 5-278 Pt 3 62

reduction were 37, 55.5 and 75%. Tenile specimens representing the variousconditions were given the heat treatment: 1610*F-24 hours-air cool, 1020*F-24 houre-air cool before testing. Figure 35 indicates 55.5% reduction to beoptim. The other reductions, 37 and, 7%, resulted in poor ductilities. Thepurpose of the present work was to evaluate the effect of the heat treatment:15600F-4 hours-air cool, 1020"F-24 hours-air cool on these materials. The.results are presented in Table 17.

Only three of the oriinal six conditions are represented. There was notsufficient stock remaining for the other three conditions to make one or moretensile specilmns. Fortunately, the two conditions resulting in the poorestductilities (about 6% tensile reduction in area) are among those represented.These conditions resulted from reducing 75% at the finishing temperatures of18000 and 1650*F, i.e., forging procedures R-1 and R-2, respectively. Thethird condition represented resulted in the best ductility (32% tensile RA) inthe original work. All specimens given the 1560oF for 4 hours coarsening anneahad more than 30% tensile RA. One specimen each for the three onditions wasvacu annealed at 1560*7 for 4 hours to remove hydrogen simultaneously withthe coarsening anneal. The average ductility, as measured by reduction in areafor the three specimens, shoved a 10% relative increase over the average of thifour specimens given the straight 156O°F-4 hours anneal followed by aging. Poithe thermal histories represented in the forging study there appears to belittle doubt that compared to the 1610OF-24 hours anneal, the 150*OF-h hoursanneal results in considerably better and more uniform tensile ductilities.

However, these results are at variance with results of a heat treatmentstudy made on the 7A1-3Ko alloy, see Table 20, WADC TR 5h-278 Pt 2. Thematerial used in this study was not the same heat as was used for the forgingstudy. The ingot was forged from 3 in. round to 1 in. square at 20006F andfinished at 1950T to 1/2 in. round. To check the previous results, sum morespecimens of this material were given the heat treatments under discussion.The results given in Table 18 confirm the previous findings that the 16106F foi24 hours coarsening anneal produced better ductility than the 1560OF for 4hours anneal. Also, two specimens initially annealed at 1610"F for 24 hours,aged and then annealed at 1560F for 4 hours and aged, did not show improvedductility over specimens given one or the other of the double heat treatments.The purpose of the double heat treatment was to superimpose the 15606F a/P raton the structure obtained by the 1610°7 for 24 hours coarsening anneal. Thehigher finishing temperature of this material compared to those used in theforging study resulted in significantly different microstructures, as may beseen by comparing Figures 36, 37'and 38. Structures produced by the 1610OFanneal, 102070 age were similar but more coarse. The material used in the healtreatment study um finished in the P, while the forging study material wasfinished at two temperatures in the a-p field. This may be the significantfactor In the difference in heat treatment res e observed. It may be notedthat microstructures remting from the forging study have a considerablygreater average man free path in the a phase than that of the heat treatmentstudy. Also, the latter structure has a relatively large prior P grain sie,while the prior P structures of the former microstructures have been complete]obliterated.

WADC TR 54-278 Pt 3 63

40

1800* F

30

Izw0

w

z • A-20 20-

w

S1650 F

IO-

HEAT TREATMENT: 1610 0 F-24HRS-AC, 1020OF-24 HRS-AC(EACH POINT AVERAGE FOR TWO TENSILE SPECIMENS)

0 20 40 60 so 100

REDUCTION AT FINISHING TEMPERATURE, PERCENT

FIG. 35-INFLUENCE OF FINISHING TEMPERATUREAND PER CENT REDUCTION IN AREA ONT1j ,TEJNSE f R L -&R.EA OFTW-' 4%A4-8 M 3 A

WAK IM 94 n -24 Pt 3 64

TABLU 17

HEAT TREATM r COMPARISON FOR Ti-7h-3o (ON) MATERTAL

Ultimate YieldTensile Strength Reduction Fractur

Forging Strength (0.2% Offset) in Area Elongation StressProcedure psi psi % % psi

1610 F-24 hrs-AC_, 1020 F-24 hrs-AC*

R-1 172,000 158,000 6.3 3.0 184,00176,000 160,000 5.5 3.0 187,000

R-2 156,000 146,000 6.3 2.0 165,000

R-3 160,000 136,000 314.9 15.0 216,000163,000 136,000 29,3 14.5 208,000

1560*F-4 hrs-AC, 1020OF-24 hrs-AC

R-1 155,000 137,000 44.4 18.0 216,000

R-2 158,000 146,000 31.9 13.5 204,000

R-3 159,000 141,000 32.7 15.0 208,000160,000 142,000 38.1 16.0 213,000

VA 1560F-4 hrs-AC, 10200F-24 hrs-AC

R-1 155,000 137,0o 38.8 15.0 210,ooc

R-2 158,000 146,ooo 43.2 19.0 2114,OOC

R-3 154,000 138,000 41.1 16.o 205,OOC

R-1 Forged from 3 to 1 in. round (89% reduction) at 2000"F, and finished

to 1/2 in. round (75% reduction) at 1800*F.

R-2 Forged from 3 to I in. round (89% reduction) at 2000*F, and finishedto 1/2 in. romd (75% reduction) at 1650*F.

R-3 Forged from 3 to 3/4 in. round (94% reduction) at 2000"F, andfinished to 1/2 in. round (55.5% reduction) at 1800F.

* Data for 1i heat treatment fros WADC TR 54-278 Pt 2,(AF 33(038)-22806).

WADC Tm 54-278 Pt 3 65

TABLZ 18

HEAT TRYATMT COMPARISON FOR Ti-7Al-3o (IN) MATERIAL

Ultimate yieldTensile Strength Reduction FractureStrength (0.2% Offset) in Area Elongation Stress

psi psi % % psi

16100F-24 hrs-AC, 10200F-48 brs-AC

163,000 144,o00 28.0 14.0 210,000*163,000 143,000 25.8 15.5 206,000

156oF-4 hre-AC, 10206F-' 8 bra-AC

160,000 134,000 21.6 14.o 194,000*

162,000. 142,000 16.0 13.0 190,000

164,000 l.4, 000 23.8 13.0 204,000

1610*F-24 hrs-AC, 1020F-48 brs-AC,15 60-4 sra-AC, 1020F-48 hrs-AC

164,ooo 143,000 10.8 8.0 184,ooo162,000 144,000 21.6 15.0 199,O00

* These data from Table 20. WADC TR 54-278 Pt 2. Specimen agedfor 24 hours at 1020F. This is not believed to result inreal difference in properties cmpared to aging for 48 hours.

wADcTR 54-278 Pt 3 66

Neg. No. 3-1212 X 75CFig. 36

Ti-7A1-3Mo (IN), 80% reduction by forgingat finishing temperature of 1950T.Heat tratmnent: 1560F-41 re-AC,1020*F-I,5 hrs-AC. a.+

Neg. No. 11213 X 75(Fig. 37

Ti-7A].-3Ko (4N), 75% reduction by forgingat finishing temper'atuire of 18006F.Heat treatment: VA l560*F-4 bra-AC,102OF-24~ bra-AC. a. +

Neg. No. 11214 ~ X 75CFig. 38

Ti-MA-34o (1OO), 75% reduction by forging&t fildshin te -pre of 1w 47.102 0-F- 2 4s bra-AC. a + P

ftchant: 60 cc glycerine, 20 cc HNO,, 20 cc HF.

wADC TR 54"-278 Pt 3 67

Z. Factors Affecting Ductility of Titanium-Aliumn Alloys in the Range from

6 to 0%Al0inum

1. Introduction

In previous work under Contract No. AF 33(616)-2351(5) the tensile proper-ties, particular3y ductility, of Ti-8A1 were found to be very dependent uponthe temperature from which the specimens were quenched. The relationships foundfor reduction in. area and proportional limit as functions of annealing tempera-ture are shown in Figure 39. In this work annealing time was increased asannealing temperature was lowered in order to compensate somewhat for thedecreasing diffusion rate. With decreasing annealing temperature tensile reduc-tion in area decreased from about 35% to sero for annealing temperatures in therange from I4706 to 1020"F, and increased to the 32-31% level in the range from10200 to 750OF where it was constant down to 390"F. Vaczmn annealing to removehydrogen did not appear to lessen susceptibility to embrittlement when the alloywas annealed at the critical temperature of 1020OF.

Upon annealing resistivity specimens at successively lower temperatures andmeasuring resistivity at room temperature after each anneal, the relatonshipshown in Figure 40 was obtained. Annealing times were the same as used in thetensile study sunarised in Figure 39. The drop in resistivity corresponds tothe loss in tensile ductility and indicates that son. change takes place in thealloy solid solution.

2. Heat Treatments, Room Temperature Properties and-Microstructures

During the present work the above mentioned resistivity specimens werereheated to 14700F, quenched and annealed at 750F for 48 hours. The measuredresistivities are included in Figure W0. The resistivities were at the levelof the specimens as quenched from 1470"F. This means that the state of the.alloy is not changed by annealing at 750°F for 48 hours following a quench fra14700P. Thus, the apparent recovery in ductility by annealing at 750OF orlower is due simply to the fact that the annealing times were not sufficientfor the embrittling reaction to occur.

The present work was primarily concerned with studying the effect of heattreatment on the mechanical properties and microstructures of 6 and IOAI alloysand with further study of the 8Al composition as an iodide titnium-basematerial as well as a sponge base material.

Tensile properties of the 6 and lOAl alloys as quenched from variousannealing temperatures are given in Tables 19 and 20, respectively. These ten-sile data wase obtained gawwerl at two lowiinS riess .60 a4 I~O ",,".! .j$in order to get differnt sapides of strain rate. A few tete of & 20tlalloy were ran at a loading rate of 1000 lbs/zin.

The data for the 6AI alloy showed no significant variation in any tensileproperty with varying annealing temperature. Also, no significant differencein properties was noted for the different loading rates. Tensile reduction inarea values plotted versus annealing temperatures are shown in Fiare Ii,

WADC TR 54-278 Pt 3 68

ANNEALING TEMPERATURE, OF390 570 750 930 1110 1290 1470

LOADING RATE, LBS./MIN.

0 2400 1X 1000 (TENSILE DATA FROM

PART I 6500C-IHOUR -AIR COOL)

__________ ___ ___ __ if

xx

40

~20z

0 x

10 _ _

0L W200 300 400 500 600 700 Boo

ANNEALING TEMPERATURE, C

FIG. 39-PROPORTIONAL LIMIT AND REDUCTION IN AREA VS.ANNEALING TEMPERATURE FOR TI-8%AI (2N) ALLOY.

"N WADO Ti. 54-278 Pt 3 69

ANNEALING TEMPERATURE, OF

194750 930 1110 1290 1470 '1650

19.2

800OC- I HR -WQ400*C-48 HRS-WQ

19.0

1I8.8

0

10

18.4

8.-

18.0

300 400 500 .600 700 800 900ANNEALING TEMPERATURE, 0C

FIG. 40-RESISTIVITY VS. ANNEALING TEMPERATUREFOR Ta-8%AI (2N) ALLOY. SPECIMENSANNEALED AT SUCCESSIVELY LOWERTEMPERATURES.

WADC TR 544278 Pt 3 70

0[ *,V 8§ R o CI

~a H 0 0 0 0 0~ i 0 ~ 0 0

O'ri C- co c-c- coc- Cco 0 c .o 0% 9 L- Gor- co L- UVH 4H H HC -e-i HH HH gvrIr-H f- HH HHI e- - -IHr-

ri

tk 00 0 c'0J 0t 0 0i 00 UN~ OV 0-0 00v

co r- co 00, -4 m--I mA I

H- - - - H r4 HH H HHH

PQ0

A4 c'j- 0.' .JH OH %A~ 0 0c 0IU 0 H - % co m

-4 0 bk %0 0-r HH 0-f % %0 HH ~- H HH HH\ 0 \ H

-4 p -4 4 r-

Nk~~ C'J-t D c yXAr4\ HN %0 r -. 1coI -j P

420 ft~L

'3 \ c' H- 2 0 S-730 - -4E -IrIH H H H

WADC~~~ 0E5428 t37§ 0

0 0cyi~o 1A-i

HO

m H 0 2 N0 h 0IqaH.PX H H H Hc'q HHH HH i r4cvJCJ~

0 ?o-Af.%f\00o00 0000

00 c

0 UN I .~j1 I * .~~~~ 0- 1C~ I 9I 000000

Go I enQpO I0W4j it 64I

00

C.)9 100 -

430 N bO0 - , 0 o C

0\

HA- H H HG~ a g a a

r-4~~~ r- I- -

4, 0 C0 -cc m- P, H

i-ADc TR 54-278 Pt 3 72

ANNEALING TEMPERATURE, OF

80 750 930 1110 1290 1470'

TI-6/oAl ALLOY

LOADING RATE, LBS/MIN.70- 0 600

0 4000SPECIMENS WATER QUENCHED

FROM ANNEALING TEMPERATURE

60- EACH DATA POINT AVERAGE60 FOR TWO SPECIMENS

z

~50-

0

340

2

20

0

0

300 400 500 600 700 800ANNEALING TEMPERATURE, OC

FIG. 41 - REDUCTION IN AREA VS. ANNEALING TEMPERATUREFOR Ti-6/.Al (188) ALLOY

WADC TR 54-278 Pt 3 73

Most of the lOA1 specmens were brittle. Yield strength values wore sur-prisingly low. The average of the 0.2% offset yield values was 120,000 psicompared to 131,000 psi obtained for 8A1 specimens. Fracture stress generallyincreased with increasing annealing temperature as illustrated in Figure 42. Afew specimens, quenched from 16509F or higher, exhibited measurable, but small,ductilities, indicating perhaps that the allay is not inherently brittle.

Two 8AI (sponge) tensile specimens were given the 1020"F-48 hours embritt-ling heat treatment and then annealed at 14706F for 2 hours and water quenched.Tensile properties were as follows:

Ultimate YieldTensile Strength Redction Fracture

Specimen Strength (0.2% Offset) in Area Elong. StressNo. psi psi % % psi

(1) lW..ooo 126,000 33.8 13.0 191,000(2) 146,000 129,500 29.2 13.0 184,000

Comparison of these properties with those for 8A1 specimens given only the1470OF-I hr-WQ heat treatment(5) show that tensile properties including due-tility are completely recovered.

Tensile test results for 8 and lOAI specimens vacuum annealed to removehydrogen are summarized in Table 21. Results of hydrogen and oxygen analysesare included. The hydrogen contents of these specimens were generally around20 ppm, the lowest being 12 ppn. A Ti-8A1 (2N) ape cimen heat treated to aninitially ductile condition (vacuum annealed at l.70F for 48 hours-waterquenched, aged at 750OF for 64 hours-air cooled) was aged at 750"F under astress of 30,000 psi for 1006 hours. The stress-aging treatment causedembrttlment. The microstructure of this specimen shown in Figure 43 indicatescopious precipitate.

A group of iodide titanium-base 8A1 tensile specimens were vacuum annealedat 1560°F for 6 hours and water quenched. One was tested as-quenched, one as-aged at 1020°F for 42 hours, and a third as-aged at 1020F for 500 hours. Theas-quenched specimen was highly ductile and no precipitate was apparent in itsmicrostructure. The specimen aged for 42 hours had a lower yield strength andmuch poorer ductility than the as-quenched specimen. The microstructure at250 X was indicated to contain small particles in the grain boundaries (seeFigure 44). Applying a dynamic vacuum during aging at 10209F for 48 hours alsoresulted in very low ductility. The specimen aged for 500 hours had a stilllower yield strength and was brittle. Its miaotrutare at 250 X was the aeas that of the specimen aged for 42 hours. However, better appreciation of theextent of precipitation and the sise of the particles is realized at 1500 X(see Figure 45). Another specimen of this material vacuua annealed at 16506Ffor 96 hours-air cooled and aged at 1020F for 48 hours had lover yield strengthbut slightly improved ductility compared to the specimen just discussed. Themicrostructure of this specimen is shown in Figure 16. Only about one-half ofthe periphery of the fracture showed signs of plastic defozation.

WADC TR 54-278 Pt 3 74

ANNEALING TEMPERATURE, OF1290 1470 1650 1830 2010

Ti-10/o Al ALLOY

140- LOADING RATE, LBS/Mew.

0 4000

130-

120

000

*110

w 0

w 100-

so-

80

600 700 800 900 1000 1100ANNEALING TEMPERATURE, OC

FIG. 42 - FRACTURE STRESS VS. ANNEALING TEMPERATURE

FOR TJ-1O%AI (188) ALLOY

W ADC TR 54-Z78 Pt 3 75

o0 0 0 IU 0 00

040

0 - -l oV\ 0 (0 0 0'* H4 0 0(Sor0 0

0* NC co CN * 00 % 0 *1*1*1

In 01% co0 IA IIU

~~cm c o 0

'.0 H O N %0 w NO ~ rV

0E- 44 J~H -

w C3

9-,4 S 4

40 41~114 t' C ~~~H~

19% 19

, WADC TR 54-278 Pt3 76

NN

Fig 43Fg.4

"N 'W 'Tt

/ q

'~s* t-r

Neg. No. 1122 X 1000 Neg. No. 31220 X 250Fig. 453 Fig. 46

Ti-8A1 (2N), VA 15670-6 hrs-WQ, Ti-8A1 (lad)., VA 1607-6 hrs-A,10200F500 hrs-AC. Stes agped:pitate. 1020*F-42 hrs-AC. a + precipitate.

30,000t 6s-006c glrine 20 cc precipitate.

WAD R 54-27 Pt37

In an effort to remove as much hydrogen as possible, two bars each ofTi-8A1 and Ti-OAl were taken through essentially the same cycle in a vacmfusion apparatus as was used for analyzing for hydrogsen. The bars were 1/2 in.round by 3 in. long. The specimens, done one at a time, were wrapped in molyb-denum foil to prevent contact with the graphite container and packing. Whereasin a hydrogen determination the sample is held at 2550"F for 15 minutes, thesespecimens were held for 45 minutes. After hydrogen removal, tensile specimenswere machined from the bars. The specimens were aged at 1020°F for 48 hours.Tensile test results for these specimens are included in Table 21. The 8A.specimens exhibited low ductility and low yield strength. A microstructurerepresentative of these two specimens is shown in Figure 47. The specimen with88,000 psi ultimate tensile strength had a fracture similar to that describedfor the Ti-8A1 (iodide) specimen vacuum annealed at 1650°F. It also showedplastic deformation around half of the periphery of the fracture. The fractureof the second specimen was more unusual in that after yielding, the load droppedsteadily. The stress at 0.5% offset was 25,000 pei and fracture occurred at17,000 psi. These fracture characteristics suggest that the yielding obtainedat relatively low stress values was the result of crack formation and propaga-tion occurring before the true yield strength of the matrix is reached. Thecracks propagate accompanied by plastic deformation until conditions for catas-trophic failure obtain, resulting in brittle fracture of the remaining soundmetal.

Specimens of Ti-OAl subjected to the hydrogen removal treatment at 2550*For vacuum annealed at 16500F for 96 hours and subsequently aged at 1020°F for48 hours were all brittle. The microstructures of the specimens vacuum treatedat 2550OF were indicated to have relatively large quantities of precipitate(see Figure 48). The microstructure of a specimen vacuum annealed at 1650°Fwas much cleaner but also appeared to contain a small amount of very fine pre-cipitate (see Figures 49 and 50).

3. Effect of Test Temperature on Mechanical Properties

Tensile properties as a function of test temperature were studied to alimited extent for the 6, 8 and fOAl alloys. The heat treatment, l020oF-48hours-water quench, wich severely embrittled the 8A1 alloy, was applied tosome of the specimens. In addition, the heat treatment, 1470OF-l hour-waterquench, which produced a ductile condition in the 8A1 alloy, was applied tosome 6Al alloy specimens. The tensile data are given in Table 22. Tensilereduction in area values are plotted versus test temperature in Figure 51. The6AI alloy appeared to suffer a transition in ductility between 320 and 850F forboth heat treated conditions. The limited data for the 1020°F for 48 hours heattreated condition indicate that the transition range may be even more narrow,i.e., between 600 and 85°F. At temperatures below room temperature the 1470°F-1 hour-water quench condition vws two to three times more ductile than the1020*F-48 hours-water quench condition in apeement with results for 8l alloyat room temperature. The BAl alloy in the sensitized condition exhibitedincreasing ductility with increasing temperature. However, at 430F the reduc-tion in area was only 2/3 of the room temperature value in the 1470°F-I hour-water quench condition. Specimens of the lOAI alloy in the sensitised condi-tion were completely brittle at temperatures to 300"F. At 480OF a specimenexhibited 10% reduction in area.

WADC TR 54-278 78

Fig 47 Fi.4

Ti8A 2N, A250 -4 mnFC i-OA lB~jV 25*F45mn-C

,1020*F-48 hrs-WQ. + precipitate. 10*F-8hs14. a+peiiae

Ne.N. 111 5 e. o 11 5Fig. 49Fg.5

Ti-lOA1 ~ ~ \\B) VA15*-6hsA, TI-l 1Bsm sFg 9 h

10207-48\ hr-C aeyosr- peiptt saltl oeapralprcpttinganbudre. ahe hihe manfcain

Etchant: 60c\lcrn,2 c N3 0c F

WADC TR5-7 t37

TABLE 22

TDPERATURE DEPENDENCE OF TENSILE PROPERTIES OF Ti-Al ALLOTS

UltimateTest Tensile Reduction FractureTamp Strength in Area Elong. Stress

Alloy Heat Treatnent (oF) psi % % psi ,

6A1 (in) 1020"F-48 hrs-WQ -25 158,000 9.0 3.5 174,00032 150,000 6.5 3.0 160,00060 151,000 9.0 4.0 166,ooo84 141l,000 30.0 14.0 173,000

1470"F-I hr-WQ -4o 167,000 25.5 10.0 222,00032 142,000 19.5 13.0 171,00084 136,ooo 35.o 17.0 176,000

8A1 (2N) 10200F-16 brs-WQ 84 11O,00 0.0 0.0 110,000

10200F-48 hrs-WQ 210 126,000 5.5 3.0 133,000306 131,000 15.5 8.0 155,000396 112,000 20.0 10.5 142,000

lOAl (IBB) 1020°F-48 hrs-WQ 210 73,000 0.0 0.0 73,000302 89,000 0.0 0.0 89,000478 113,000 10.0 3.0 125,000

WADC TR 54-278 Pt 3 80

0 _ _ _ _ _ _ _ _ 0

0 (A U 0I.- f.0 0

x0 re)OD N

0

LL

0 0 0co N I

0 N

0 0

W -w

LU LUJ

Ix-

0~ ~ 0 _____

0~~~ ~ 0__________

VIt

-0 0 0 0 0 'IT in I

1LN3) 83d 'V3NV NI NOI13flG3N 3-IISN3J.

WADC TR 54-278 Pt 3 8

Notch impact as a function of temperature for the 6A1 alloy is given inFigure 52. Impact values were more than twice as high for the 1470"F-1 hour-water quench condition as for the 1020 F-48 hours-water quench condition.Vacuum annealing to remove hydrogen resulted in no significant improvement ofimpact properties of specimens aged at 10200F. The data are not sufficient tosay whether a transition occurs in the temperature range 32* to 212OF coveredby these tests.

The low temperature tensile and impact data for the 6A1 alloy establishthat it embrittles similarly to the 8AI alloy, but not as severely. The lOAlalloy is apparently embrittled in like fashion.

Some interesting microstructural details have been observed in this inves-tigation. The discussion of these observations is included as an appendix,since they are more likely to obscure rather than elucidate the essential find-ings of this investigation.

4. Identification of the Embrittling Agent

Since long times In the temperature range 7500 to 1020"F are apparentlyneeded to develop conspicuous precipitate, a survey was made of the microstruc-tures of creep-rupture specimens of 6A1 (3B) and 8AI (J) alloys. Specimens ofthe latter alloy tested at 8008 and 1020F showed indications of precipitate,with best indications given by specimens tested at 1020°F (see Figure 53).Three specimens of Ti-6A1 (3B) were examined. These specimens had been testedat 800F and had rupture lives of 148, 216 and 319 hours. No good indicationof precipitation was found in the 108-hour rupture specimen; a trace of a pre-cipitate was indicated in the 216-hour specimen; and rather good indication ofa precipitate was found in the 319-hour specimen. However, the evidence ofprecipitation in this specimen was not as good as that shown in Figure 53.

In order to learn more about the precipitate specimens of 6, 8 and fOAl,alloys were vacuum annealed at 1560OF for 6 hours and air cooled. This treat-ment was followed by aging at 1020F for 48 and 500 hours. Initial hydrogencontents were 233, 100 and 135 ppm for the 6, 8 and lOAJ alloys, respectively.Another group of specimens was simultaneously heat treated except that the1560OF anneal was carried out In evacuated Vycor bulbs. These specimens wereexamined metallographicaly. All were indicated to contain precipitate. Themicrostructures were similar to that shown in Figure 43. In general, the vacuumannealed specimens appeared to contain much less of the precipitate than theother specimens. Specimens of 16AI alloy were also examined after aging at1020°F for 48 and 500 hours. These specimens had microstructures similar tothose of the lower aluminum content alloys discussed above. However, the amountof precipitate in the 16AI alloy was indicated to be far less. This alloyanalyzed 56 ppm hydrogen.

Several specimens in which the precipitate was highly developed wereannealed at 1470"F for 2 or 3 hours and water quenched. The results are shownin Figures 54, 56 and 57. Figures 43, 45 and 55 show the corresponding struc-tures before the resolutioning heat treabent. There is no difference in thestructures.

WADC TR 54-278 Pt 3 82

0 '

0 LIA

0r 4ODi

qt 2D U.

0 0 0 0 (')

0 2 )4

*L - -

-0D

00 C 0

w>-

00 1

CI

0 0__ 0 of

CYN

'SS- 14 .L:Vdi~ a3H.LO -A kHVH

WADC R S4278 P 3 6

Neg. No. 10586 X 200Fig. 53

Ti-SAl (1J), annealed at 1200OF for Ihour-air cooled, stress-rupture testedat 1020OF under 42,000 psi for 256hours. a + precipitate.

V

ft

1(.. .' -" ' -.. ""\9

Neg. 11519 X 000

Fig. 53

Ti-8A1 (2N), same thermal history as

Fig. 43 plus 14707F-2 hrs-WQ.

Etchant: 60 cc glycerine, 20 cc HNO3, 20 cc fF.

WADC TM 54-278 Pt 3 84

Fig. 55 Fig 56

IIIf

Ne.N.154 Iw Ne. No 120 x

Fig. 57 Fig. 58

Ti81(NV J40-8hsWTi-8A1 (2N), same thermal history a i7l(o) odpesd7%

02 F2 hr-J. Fig. 45 plus 1470*06-3 hrs-WQ. 17*-8 r-Q

Ethat 60c lcrn,2.c N3 0c F7/DC T5-28Pt38

Several attempts to make identification of the precipitate phase by X-rayand electron diffraction were unsuccessful. Finally, a Debye pattern of apowder sample of the 8Al specimen stress aged at 750*F for 1000 hours wasobtained which gave unequivocal evidence of the presence of a second phase.Exposure conditions under molybdenum Ka radiation were as follows: 50 KV,12 ma and 40 hours. Lines not belonging to a titanim are recorded in Table 23.Included in the table are the lines for the TiAl phase at 38% aluminum. Com-parison of the two sets of lines strongly indicates that the second phase isTiAl (r). Since ductility was recovered by heating embrittled specimens at1)70°F, a powder sample of stress-aged specimen was annealed at 1470F for 1hour and water quenched. The Debye pattern of this powder showed that thesecond phase did not go back into solution. This finding is consistant withthe microstructural observations. The TiAl phase is known to be brittle. (6)However, the facts suggest that the umbrittlement is due to coherency of pre-precipitation particles and that fully precipitated particles in the amountobserved in this work are not harful to ductility.

Further confirmation of the identity of the precipitate was obtained witha sample of 7A1 alloy which was cold pressed 70% and subsequently annealed at1470°F for 180 hours. The microstructure of this sample was clearly two phase,as shown by Figure 58. The Debye pattern of a powder sample of this materialis also included in Table 23. As in the case of the 8A1 specimen, the agree-ment with the TiAl (Y) pattern is rather good.

These findings show that the Ti-Al system as reported in the literature(7)is in error with respect to the extent of aluninum solubility in the a phase.

IV. SMARY AND CONCLUSIONS

The alloys Ti-5AI-2Ag and Ti-5Al-7Ag exhibited fairly typical age hardeningresponse over the temperature range 800* to I120°F. The 5Al-2Ag and 5Al-7Agalloys have poorer creep-rupture properties at 8OF than the 6A1 binary alloy.

Complexing the a phase of the 6A1-3fo alloy with additions of tin and zir-conium increased tensile strength with no sacrifice of room temperature ductility.Complexing the P phase by substitution of chromium, manganese and vanadium alsoeffected an increase in strength with no impairment of room temperature ductilitywith the exception of the 6A1-2Mo-IMn modification. Vacuum annealing this alloyrestored ductility comparable to the base material. The 6A1-3Cu, a + Ti2Cualloy was exceptionally strong at elevated temperatures but had very poor orgood. rom temperature ductility depending upon microstructure. Widanstlttenmicrostructures were brittle or of low ductility, while equiaxed structureswere of relatively high ductility.

On a stress-rupture basis the a cimplexed alloys were slightly stronger at1020"F than the 6A-3Mo alloy. The creep-rupture properties of the P complezedalloys were significantly inferior to the base material at both test temperatures.

WADC TR 54-278 Pt 3 86

ITABKE 23

TITANIUM -AIINU4M ATLOYS

DEBTE LINES OTHER THAN ALPHA TITANIUM

8A1*7A1 P-70* 750"F-I06 hrs TiA1

1470F-180 hrs 30,000 psi 38AI

I d I d 1290°F-330 hro

4 (?) 3.31 3 3.31

2 2.87 4 2.83 2.81

1 2.14 1 2.10 2.08

1 2.02

<<1 (?) 1.93 1.99

1 1.818 1.804

2 1.652 1 1.638 1.648

<< () 1.571 <1 1.536

<<1 (?) 1.519

1 1.398 1.408

1 1.384

1 1.298

1 1.279 2 1.271 1.257

* Mo radiation, 50 KV, 12 ma, 40 hours.

WADC TR 54-278 Pt 3 87

Considering creep resistance, the 6A1-3Cu and the a complexed alloys worecomparable at 800OF and were markedly superior to the 6AI-3Mo alloy. At 1020"Fthe creep resistance of the 6AI-3Cu alloy was comparable to the 6A1-3Mo alloy,whereas the a complexed alloys were still decidedly better.

The 6AI-3Mo alloy was stable after creep exposure in agreement with pre-vious findings for other Ti-Al-Ho alloys. Stability tests of the other experi-mental forging alloys gave variable results. In some instances exposure tostress at elevated temperatures produced severe loss of room temperature duc-tility, while in other cases for the same alloy comparable exposure conditionsproduced little or no property change.

Modification of the 6A1-4V sheet alloy by substitution of molybdenum forpart of the vanadium indicated that the 6AI-2Mo-2V alloy was best of the threemodifications both on a tensile strength and creep-rupture comparison. The6AI-2Mo-2V alloy was appreciably weaker than a 6AI-4V alloy tested in bar for,but the great difference in strength is felt to be due in large measure to speci-men geometry.

All three sheet modifications were indicated to be stable on the basis oflimited stress-aging experiments.

The Ti-7Al-34o alloy had been found to have variable ductility dependingupon hot working history as given the heat treatment: 1610oF-24 hours-aircool, 1020*F-24 hours-air cool. However, as heat treated: 1560*F-4 hours-aircool, 10201F-24 hours-air cool, ductility was unifonmly high. Vacuum annealingto remove hydrogen combined with the 1560°F-4 hours anneal resulted in furthersmall improvement in ductility.

The 8AI alloy made with iodide titanium and vacuum annealed was found toembrittle in the same manner as the sponge-base alloy. A microstructural con-stituent is definitely associated with the embrittlement. The microstructuralconstituent, which was developed by long anneals at 10200 and 750°F, was noteliminated by reheating to 1470*F and quenching. However, electrical resis-tivity, which had been found to decrease when specimens were annealed in thetemperature range of embrittlement, returned to the high value associated withthe ductile state upon reheating to 1470°F and quenching. Also, tensile proper-ties including ductility were completely reversible.

The microconstituent developed in the 8AI alloy was also developed in the6, 10 and 16A1 alloys. The 6A1 alloy was ductile at room temperature (850F)in all heat treated conditions. However, specimens annealed at 102DeF for 48hours showed considerable loss in ductility when tested below room temperature.The 6A1 alloy quenched fran 1470°F also showed ductility loss when tested belowroom temperature, but the loss-was considerably less than for the aged at 10206Fcondition. While tensile properties of the 6A1 alloy at room temperature wereunaffected by aging at 1020F, Charpy V-notch impact strength was decreased bya factor of two.

The lOAf alloy was brittle for all heat treatments. However, fracturestress was higher the higher the temperature from which the specimens were

WADC TR 54-278 Pt 3 88

I quenched. Highest quenching temperature was 2010"F. The lowest test temperture at which lOAI specimens aged at 1020F showed some ductility was jBOeF.

X-ray diffraction evidence indicates that the embrittling precipitate iithe TiAl (Y) phase of the Ti-Al system.

The general results show that the Ti-Al system as reported in the liter4ture is in error.

WADC TR 54-278 Pt 3 89

REFERENCES

1. W. F. Carew, F. A. Crossley, H. D. Kessler and M. Hansen, "Titanium Alloysfor Elevated Temperature Application", Armour Research Foundation SummaryReport, WADC TR 52-245, to the Materials laboratory, WADC, Wright-PattersonAir Force Base, Ohio under Contract No. AF 33(038)-22806, November, 1952.

2. W. F. Carew, F. A. Crossley, H. D. Kessler and M. Hansen, "Titanium Alloysfor Elevated Temperature Application", Armour Research Foundation SummaryReport, WADC TR 53-177, to the Materials Laboratory, WADC, Wright-PattersonAir Force Base, Ohio under Contract No. AF 33(038)-22806. August 1953.

3. W. F. Carew. F. A. Crossley and D. J. McPherson, "Titanium Alloys forElevated Temperature Application", Armour Research Foundation SummaryReport, WADC TR 54-278 Part 1, to the Materials Laboratory, WADC, Wright-Patterson Air Force Base, Ohio under Contract No. AF 33(038)-22806,August 1954.

4. W. F. Carew, F. A. Crossley and D. J. McPherson, "Titanium Alloys forElevated Temperature Application", Armour Research Foundation SummaryReport, WADC TR 54-278 Part 2, to the Materials aboratory, WADC, Wright-Patterson Air Force Base, Ohio under Contract No. AF 33(038)-22806,December 1954.

5. D. J. McPherson and W. Rostoker, "Studies of Phase Relationships andTransformation Processes in Titanium Alloy Systems", Armour ResearchFoundation Summary Report, WADC TR 54-101 Part 2, to the MaterialsLaboratory, WADC, Wright-Patterson Air Force Base, Ohio under Contract No.AF 33(616)-2351, December 1954.

6. J. B. McAndrew and D. J. McPherson, "Investigation of the MetallurgicalCharacteristics of the 36% Aluminum- Titanium-Base Alloys", Armour ResearchFoundation Final Report, WADC TR 53-182 Part 2, to the Materials Laboratory,WADC, Wright-Patterson Air Force Base, Ohio under Contract No. AF 33(616)-196, January 1955.

7. E. B. Bumps, H. D. Kessler and M. Hansen, "Titanium-Aluminum System",Trans. AIME 194 (1952) p. 609.

WADC TR 54-278 Pt 3 90

\S

APPENDIX

SOME OBSERVATIONS ON MICROSTRUCTURE

Five of the 8AI (Iod) tensile specimens for which data are given in Table21 were sectioned longitudinally and examined metallographically. A commonfeature of the microstructures near the fracture surface was the presence ofslip traces. Figure 59 shows slip traces of a specimen quenched from 1560OFand having 30% tensile reduction in area. Figure 60 shows slip traces of aspecimen air cooled from 1650°F and then aged for 48 hours at 1020F. Thisspecimen possessed 3% RA. Slip traces such as these were also found in one ofthree fractured tensile specimens of IOAI alloy. The specimen had been quenchedin iced brine from 1830°F. Slip traces were not observed in 6AI specimensfractured at room temerature.

The occurrence of slip traces was not limited to specimens which wereplastically deformed. They were also observed in 8AI (iodide) specimens havingthe following thermal histories subsequent to forging:

(1) 1470OF-1 hr-WQ, 1290°F-I hr-Q

(2) 1470-F-1 hr-WQ. 111OF-2 hrs-WQ

(3) 1470"F-1 hr-WQ, 1020"F-48 hrs-WQ

These specimens contained approximately I00 ppm hydrogen. Electronmicrographsof the specimen shown in Figure 60 at 11,000 and 22,000 X show the slip tracesto be bands approximately 3 x 10-5 am wide. The matrix was shown to be uniformlypitted with the bands considerably less pitted.

Ordinarily, under circumstances such as these, slip traces are not observed.Their presence indicates that something happens either during slip, or subse-quently, during the preparation for metallographic examination which producesthe etching effect. The slip traces present the possibility that precipitationoccurs during slip.

Another phenomenon observed, but to a more limited extent, was internalcracking. The cracks were transgranular. A specimen of 8A1 alloy containing100 ppm hydrogen was given the following heat treatment: 1560F-6 hrs-WQ,1020°F-500 hrs-WQ. The microstructure of the specimen showed copious precipi-tate in the grain and subgrain boundaries. The specimen was cut in half, andone half was solution heat treated at 1470°F for 2 hours and water quenched.The solution treated half of the specimen was found to contain nunrous cracks(see Figure 61), while the other half was devoid of cracks. The mechanism of"the formation of these cracks is unknown.

WADC TR 54-278 Pt 3 91

Neg. No. 11535 X 1000

ec I Fig. 59AN' Ti-8A1 (Iod), VA 15607-6 hro-WAQ. Fractured

~ ~ ~ tensile specimen. RA -32%.

Neg. No. 11431 X 2000

Fig. 60Ti-8A1 (Iod), VA 16507F-96 Iws-AC,1020 8F-48 hrs-WQ. Fractured tensile

-spe cimen . RA =3%.

~ '~* * .. Neg. No. 11518 F 61X 250

4 Fig.OF.. 61sW- 4, Ti-8A1 (2N), 15607-6 hrs-WQ, 10207'-500

:' 4r-1Q 140F2 r-

Etchant: 60 cc glycerine, 20 cc HNO3, 20 cc HF.

WAC TR54-278 Pt 3 92