US2867534A - Titanium base alpha dispersoid alloys - Google Patents

Description

Jan. 6, 1959 R.1.JAFFEE l-:TAL

TITANIUM BASE ALPHA DISPERSOID ALLOYS :s sheets-sheet 1 Filed Jan. 23, 1957 WUIU @im QNUEUD OWK MUQQ MMM S 5. mi n TF6. N m w1# ,MN Tu. E wm A EH Jan. 6, 1959 Filed Jan. 23, 1957 @fa EFFECT 0F INCREASING 77N CONTE/vr o/v THE HA fPo/vsss of 77- 7 (U-5,45: ALLOY R. l. JAFFEE ETA'L 2,867,534

TITANIUM BASE ALPHA DISPERSOID ALLOYS 3 Sheets-Sheet 2 2 O157-A NCE FROM QusNcHEa END, /NcHEs assu/7N ssa/voel VH Saz-191A BY Ho/PA cf l?. OGoE/v.

ATTORNEKS.

Jam 5 1959 R. l. JAFFEE ETAL 2,867,534

TITANIUM BASE ALPHA DISPERSOID ALLoYs Filed Jan. 25, 1957 3 Sheets-Sheet 3 oom QBLUML Dmot DU WN 12.. @NIR QLumCbMOnl bbh| b\| R 5. ROBERT BYHORA CE OGDE/v.

A TTONEYS.

ijnited TETANIUM BASE ALPHA DISBERSOID ALLOYS Appiication .ianuary 23, 1957, Serial No. 635,705

26 Claims. (Cl. 7.5-1755) This invention pertains to alloys of titanium containing as essential constituents at least one of the active eutectoid beta promoters, copper, beryllium, nickel, cobalt and silicon, and at least one promoter of alpha titanium selected from the group consisting of aluminum, tin, anti- .fno-ny and zirconium, characterized by good hot Workab"' excellent thermal stability and elevated temperature strength, and in being hardened and strengthened in remarkable degree, without undue embrittlement by quenching or cooling at asufiiciently VVrapid rate from above the critical or eutecto-idtemperature. l

'Copper, nickel, cobalt, silicon'and beryllium all belon to the group of .so-called active eutectoid or compoundv forming beta promoters. That is to say,lthese elements when present as alloy additions to titanium, form compounds on cooling'from above to below the critical temperature, vto produce a microstructure, consistingof alpha titanium plus compounds ,dispersed throughout the alphamatrx. L l

Of the metals of the active eutectoid beta promoter group, copper is effective over-the range of'about 0.5 to nickel and cobalt over the range of about 0.5 to 12% each or. in total amount,l silicon over the range of about 0.25 to 3% and beryllium over the range of about 0.1 to 2%. w

Ot the alpha promoters aforesaid, aluminum mayrange from about 0.5 to1l2%, tin from about 0.5` to 23% and antimony-frcm about 0.5 to 19% and zirconium from about 0.5v to 40%, the preferred range being, however,

The total content of the. alpha proabout 0.5to 20%. meters of this group should be held Within limits of about 0.5 to 40%, the prcportioning being such that when any one is present on the high side of its range, the total content of the others should be held on the low side or sides of their respective ranges.l ln these quench hardenable alloys, tin or antimony may be substituted for aluminum in the proportions of about `3% by weight of tin or antimony for 1% of aluminum. The upperlimits of fabricability for the alloys of the invention are'about 12% for aluminum, 23% for tin and 19% for antimony, l20% for copper, 12% for nickel, 12%for cobalt, 3% for silicon and 2% for beryllium. rhere is no upper limit for fabricability of zirconium alloying additions to titanium, since zirconium is beta isomorphous with titanium. an alpha promoter because alloys of titanium and zirconium always revert to the all-alpha form ony cooling to room temperature fro-m the beta eld.

Of the interstitial alpha promoter additions, carbon is eective over the range of about 0.05 to-0.5%, oxygen over the range of about 0.05 to 0.4% `and nitrogen over the range of about 0.025 to 0.25%.` They total content of the interstitials shouldV not exceed about 0.5%, the lower effective limit being about`0.0?.5%.V

Zirconium vis nevertheless llffiatent O tioned, particularly those containing copper, have out-v promoters in relation to the alpha promoter additions is preferably such that when the beta promoter contentV is on the high side, the alpha promoter content should be on the low side and vice versaof the rangesfor each aforesaid.

Preferred ranges for optimum quench hardenabilit of alloys according to the invention are abo-ut: 5 to 8% copper, 4 to 7% nickel, 5 to 8% cobalt, 0.5 to 1.0% silicon, and 0.5 to 1% beryllium, together with about 2 to 4% aluminum, 4 to 8% tin or 4 to 8% antimony.4

We have found that the series of alloys above menstanding properties in a numberof respects as follows: They have improved hot workability as compared to previo-usly known types of titanium base alloys with high alpha-stabilizer content. They undergo greatly enhanced hardening and strgenthening on quenching or rapidly cooling from the beta and alpha-beta temperature ranges.: They possess excellent elevated temperature strength up' to about 1200a F. They have a high modulus of elasticity and a high ratio of modulus of elasticity to density as compared to previously known types of titanium base alloys. They have excellent thermal stability, particularly asv annealed below the lower critical temperature, the lower critical temperature being that at which vthe beta phase transforms on cooling to the alpha'phase plus the compounds formed by decomposition of the active eutectoid elements present. For the copper-bearing alloys, this compound `is Ti2Cu which corresponds to Fe3C in the iron-carbon phase diagram of steel. vThis thermal stability results from the fact that'the alloy is completely transformed toalpha titanium plus compound. These alloys also possess eX- cellent creepresistance and excellent welding properties.

Thethermally. stable conditions for these alloys may be obtained either by slow cooling through the critical temperature range, or by annealing for about A to 24 hours at a temperature below the critical, usually about 501100". f F. below the critical.` The longer the annealing cycle the softer the material.

At least half of the total content of active eutectoid beta promoters presentin these alloys can be substituted by o-ther'stabilizers of beta titanium such as the beta isomorphous or sluggishly eutectoid beta promoters, molyb-z denum, vanadium, tantalum, columbium, chromium, manganese, ironA and tungsten.l Inl this` substitution such other beta promoters may be added up to about one-half of their total permissible ranges in the absence of the active eutectoid beta promoter addition. Thus, for examplmmolybdenum, vanadium, tantalum and columbium may be added up to about 40% to the titanium-aluminum base alloys of the invention in the absence of copper, nickel, cobalt, etc. On the other hand, if the allo-y lcontains, ffor example, 10% of copper, then metal of the group molbdenum, vanadium,columbium and tantalum may be added up to about 20%, chromium and/or tungsten up to about 12%, manganese up to about 5% and iron up to about 3.5%, as compared to upper limits for these latter additions of up to 25% chromium and/or tungsten, 10% manganese, and 7% iron in the absence of the` active eutectoidaddition. v

These alloys may properly be termed alpha dispersoid alloysv for reasons above stated, and ,are basically ternary alloy systems containing one alloy addition such as aluminum, tin, antimony orzirconium, which strengthens the alpha phase, and one active eutectoid can be produced by adding two or more of the active eu'tecto'id beta promoters'for twoor,lifrlorey 'of the alpha strengthening elements, for these purposes. The lfun'damental two-phase alpha plus 'dispersedcompound"struc-` ture 'remainsunchang'ed however. i n

` These alpha dispersoid alloys are particularly 'adapted for Aelevated`temperaturc service. As stated, they are` usually two-phase alloys,"consisting of a strong stable alpha-phase and a dispersed compound phase. Superior` elevatedY temperature properties lresult from the keying effect of the'h'arddispersedphase.

The-hotwo-rking methods used 1n the fabrication ofV theseV alloysy is of 3considerable'importance, since their propertiesy as heat`tr`eate'd, are controlled to a' large'extentfby v'the prior fabrication history. There are four temperature regions in which these alloys maybe hot worked, yA'rrangedfin theordevr of decreasing tempera'- tu're, `these lregions'are: (la) Ifthe all-beta regie-n, (b) the two-phase,"alpha-beta'region, (c) the three-'phasealpnabeta-cornpoundv region, andg(d) thetwo-phase alpha-v compoundregion, 1 Y y For' 'efxanfple'in the case of Ti-AlfCu alloys of the nvention,"heat` treatment in 'the'rall-betaregion is carriedout inthe range of from about #1650 F. up to usually` knot more thanvl800 Vor l850 F.; the two-phase alpha-beta region is about 1550-165013.; the threephasefalpha-beta-compound regionis about 1475"` F.;

and Athe.two-phase "alpha-compound region'lies below F.)'to place thest'ructure in malleable condition for subsequent heat treatment. p Y v l The critical temperature for these` alloys willQvary n somewhat'with the particular activeeutectoidbeta `promoter 'additionflor the vcopper-bearing alloys it yis about 1475i iF.;forLthenickel-bcaringy about .31420 F.,

. for thecobalt-bearing aboutf1260 F., forthe beryllium-v bearing about l550 FL and for the silicon bearing about 157511?. rAs'stated, for alpha-compound annealing, all of these alloys areann'ealed'atabout 50-100 F. below the'c'ritical. The beta transus temperature of these alloys also varies *somewhat withV 'th'ej alloying content, ybeing vthevlowest temperature,` for'any g'iven'janalysis, above which the alloy is entirely beta.` The optimum temperature for, b eta .quenching these alloys'lv isj to heat the same to:4 aboutf 5 0.-l00 'F.ffV above the v.beta Atransiisl and there# lnponplqu'erfich orair'cool sldscussed'below'. 'Forfthe copper-bearing' alloys theop'timum 'quenching range for the beta 4quenched alloys is about `l500-l8`50 F., Vand for the optimurnquenching 'from the alpha-beta= eld, it is about 1580-1610o F. whenk aluminum ispresent, l,and about 1490-1550 when tinor antimonyis7 substituted'for aluminum.`

`lnthese alloys the 'be't'aphase formed above .the critical .temperature .transforms ,cornpletelyl'Y into rnarten'sit'e tine'eutectoid onA rapid, quenching, ,as yinwater, lor r11211.93. Vyr'rnxture of alpha,titaniumfandjcompounds consistingwofeu'tectoid decompositiwon-"pijoducts on -,Slo-Wer cool- .ing from. above the critical,v "as by' ko'ill quenching or A.cooling in air. vThat is; to sfayfi in-thersevalloysthe beta v v,phase is not retained on quenching fromabove the Acritical,.andV also the beta transformation jis accompanied by marked increase in hardness and'tensile strength.;

'Thema'rten'sitei formed on quenching this type of titanium alloy is quite similar to the martensite formed in "steel on quenching from the austenitizing temperature The martensite thus formed on quenching is subsequent tempering in the alpharange. converted on compound temperature range, into a microstructure c011- Y sisting of a fine dispersion of the beta eutectoid decomposition compound, such as TizCu, or titanium beryllide, etc., in an alpha titanium matrix. This is similar to the tempering of steel, wherein the ymartensite formed on the quenching vis .transforrned into a fine dispersion of Fe3C in a ferrite matrix.

Relatively high Vtemperatures are required .thus to convertfth'efmartensite inthese'titanium alloys, i. e.,"somewhat under 50-1`00 F. belowfthe critical. Thus for the copper-bearing alloys the tempering range is aboutlOOO- 1400515., which is quite high as `'compared tothe tempering of steel, the tempering range of which is about 750- 1000" F. This tempering-of. these titanium alloys conv verts the same to a suitable condition of alpha titanium and dispersed compound, for subsequent service at elev ated temperatures. 'f

The high softening temperature. of these alloys adapts them for hightemperatureuse in such applications fas rotor, blades for 'jet engine compressors, bolts forcom.- pressor casings, linings and housings for jet engine combustio-n chambers, etc. f

,As lan example of the quenched and tempered properties obtainable with these alloys, particularly those containing coppen an analysis comprising about 4% aluminum, 6% copper`,balance titanium of commercial purity has,.as water quenched from 1580-1600 F., a Vickers hardness of 375, a tensile-strength of 190,000 p. s. i., a 0.2% offset yield strength of 130,000 p. s. i., an area reduction of about 30%, a tensile elongation of about 10% and a Vminimum bendv radius of about 4T. This combination .of properties has never hereto-fore been equaled intitanium'alloys, in. so farv as thepresent applicants are aware. VOn tempering this alloy at about 1300 F. for about 1 to 16 hours, the Vickers hardness is reduced .from 375 to 330350,1with an improvement'in bend ductility from 4T as quenched to 3T as tempered. In this condition-y exceptionally good thermal stability and elevated temperature strength Vis obtained for service conditions up to about 60G-1000" F.

The following Table I shows the effects of the various heat treatments aforesaid on 'the mechanical properties of a typical alloy according tothe invention containing 4% aluminum, 6% copper, balance titanium, as compared to the corresponding binary alloy omitting the alpha promoter, i. e., containing 6% copper, balance ti.anium.

TABLE `I The. e'ect of heat treatment on .the tensile propertesvof i two alpha-dspersoz'd alloysk Tensile Properties of 0.040l sheet;

p. s. i.X1,000 Percent Percent Alloy and Heat Treatment Elonga- Reduction in tion ln 0.2% Ultimate 1I Area Offset Strength Yield YTifCu:

Beta furnace cool 110 14 19 Alpha-cmpouncl anneal... 75 100 18 19 Alpha-beta quench 130 170 5 16 Alpha-beta. quench and tamper 120 11 25 Y Beta quench- 135 180 4 8 Beta quench ani temper.-- 130 11 16 Il--LtAl-Cu:A f

Bata; furnace cool 150 19 A 20 Alpha-communi annea1 110 120 15 30 Alpha-beta quench 150 175 f 4' l 1 10 Alpha-beta quench and 1 tempar i l 135 145h 11 `25 l Beta quench .1- 4 ".180 210. T5 8 Beta quench and temper..- 195 3 3 r (1) Beta furnace cool heat treatment.-In this heat treatment, the amount and type of prior work given to the alloy is relatively unimportant. The alloy'is heated to some temperature above the beta transus and held at this temperature for a sufficiently long time to insure transformation to an all-beta structure. Following the annealing period, the alloy is slowlyI cooled through the transformation range, usually by allowing the alloy `to cool with the furnace. yThis treatment results in a structure consisting of a basket-weave alpha pattern surrounded by a. coarse alpha-compound eutectoid structure., The amount of the eutectoid varies with composition, some alloys showing islands of alpha, while in others the various alpha particles are closely in contact. The principal advantage of this heat treatment is its relative independence from prior fabrication history.

(2) Alpha-compound anneal heat treatment.-The prior fabrication controls the structure of alpha-compound annealed material to a considerable extent. Whenever possible,A the alloys are rolled in the alpha-compound region as the final hot-working operation. In some alloys, however, particularly those containing large amounts of aluminum, it is necessary to finish the fabrication in the lower part of the alpha-beta region or in the alpha-beta-compound region due to a compositionalinduced alloy brittleness. The fabrication operations normally serve to break up the compound phase as it forms and to distribute it randomly throughout the matrix. Following the final rolling operations, the alloys are annealed at a temperature just bellow the eutectoid temperature in order to promote compound agglomeration and to remove any residual cold Work from the structure. Following this treatment, the alloys are air cooled.` The structure consists of a uniform dispersion of spheroidal compound articles in a fine-grained alpha matrix.

(3) Alpha-beta quench heat treutrnenz.-'l`he` beta phase of alpha-dispersoid alloys transforms to martensitic alpha or to a very fine eutectoid structure on quenching. Because of the fineness of the structure, it has not been possible to learn the exact nature of the transformation product. The quenched structures have considerably improved strength, but decreased ductility. The purpose of this heat treatment is to develop some of the increased strength resulting from the presence of the martensitic alpha or fine eutectoid while retaining good ductility due to the presence of equiaxed alpha. The final structure is controlled to a considerable extent by the fabrication history. Normally, the alloys are finished in alpha-beta region and aircooled. lf sufficient reduction is given in the alpha-beta region, a uniformly fine equiaxed-alpha structure either surrounded by the transformation product or containing islands of the transformation product, depending on composition, is formed. On heat treatment, the alloys are heated to some point in the alpha-beta region and allowed to come to equilibrium. At the annealing temperature, the alloys consist of equiaxed alpha in a matrix of beta. The size and shape of the alpha grains is controlled by the prior fabrication history, while the amount of alpha is controlled by the annealing temperature. At the end of the equilibration period, the alloys are quenched rapidly to room temperature, transforming the beta either to martensitic alpha or to a very fine alpha-compound eutectoid.

(4) Alpha-beta quench and temper heut treatment.- Tempering of an alloy quenched from the alpha-beta field promotes decomposition of the martensitic alpha into compound and alpha and spheroidization of the fine alpha-compound eutectoid. By controlling the tempering conditions, the type of compound structure can be controlled. Thus, it is` possible to develop either a very fine structure or one which is comparable to in coarseness that which resultsfrom the normal alpha-compound anneal heat treatment. out at some-temperature below the eutectoid temperature. yBoth the amount anddistribution of-the compound phase are'controlled by the `quenched structure.

y (5) Bem quench lheat treatmentln this heat treatment, the alloy is first heated for a suicient time at sometemperature above the beta transus lto insure complete transformation to beta and then quenched to room temperature. As in the case of` the beta furnace cool heatv treatment, the final structure, with the possible exception of grain size, is independentof the fabrication history. The structure of the quenched alloy yconsists almost entirely of martensitic alpha. If the quench is delayed slightly or if ythe section size is large, regions containing very fine eutectoid are formed. This heat treatment has the advantages of producing high strengths in most alpha-dispersoid alloys and of being relatively indestructures thus far obtained have all contained'a-rela-l tively fine compound dispersion.

Examination of the data given in Table I above shows that the maximum ductility was obtained when the alloys were given the alpha-compound anneal heat treatment. The beta furnace cooled heat treatment results in a stronger, less ductile structure than the alpha-compound anneal. In the Ti-6Cu alloy, copper functions both as an alpha strengthener and as a compound former. In the Ti-4Al-6Cu alloy, aluminum serves as an alpha strengthener while copper is present chiefiy because of its compound-forming ability. Quenching these alloys to form martensitic alpha results in an increase in strength and a decrease in ductility. As would be expected, the increase in strength is greater on quenching from the all-beta region than from the alpha-beta region. l

Because of the rapid decomposition of the beta phase of these alloys during cooling, appreciable improvements in properties may also be obtained by heating sumciently high to form some beta phase, and air cooling. Thus, it is possible to carry out extensive forming operations on the alloy in the soft condition and then harden the formed product by an elevated-temperature,anneal followed by air cooling. Such a heat treatment results in only a minimum amount of distortion. The amount of hardening obtained in a given alloy is controlled bothlby the temperature to which the alloy is heated and by the velocity of the air cool. The effects of this heat treatment on several alpha-dispersoid alloys are shown in the following Table I-A.

Tempering is normally carried TABLE I-A Increase in hardness of severalialph'a-dspersod alloys resulting from the air-cool heat treatment Vickers Hardness Number Anneal (l0-Kg. Lead) Composition (Wt. Percent), ing Tem- Bal. Titanium perature,

F. Soft After Air Increase Condi- Cooling in Hardtion 1 ness 4G11 1, 600y 220 253 33 60u.. 1, 600 218 306 88 80u 1, 600 226 401 75 2Al-6Sn-4Cu-0-1 Be 1, 600 292 400 108 6Sn-4Cu- 1, 600 271 333 62 9Sn- 4Cu 1, 500 376 387 11 l Normally as alpha-compound annealed.

` The following Table I-B shows the effects of various of` the heat treatmentsabove discussed on the room temperature mechanical properties of a typical forging alloy according to the invention, having the composition T1--7A1-3Cu TABLE' I-B Room temperature tenszle propertzes of a T1-7A1-3Cu alloy [Rolled from to )4 plate at 1700 F. prior to heat treating as indicated i below. Tensile tests in transverse direction] Tensile Properties, p. s. i. 1,000 Percent Percent Heat Treatment Elonga- Reduction tion in 0.2% Ultimate in 1" VArea Offset Strength Yield 7 hrs. at 1,400 F., air cool 122 135 11 11 7 hrs. at 1,400 F., air cool plus age 145 8 8 7 hrs. at 1,400 F., furnace cool 132 144 10 8 7 hrs. at 1,400 F., furnace cool plus age s 140 148 10 9 6 hrs. at 1,400 F., plus 24 hrs.

at 1,200o F. and air cool to room temp 123 137 14 23 6 hrs. at 1,400" F., plus 24 hrs.

at 1,200 F. and air vcool to room temp. plus ase 133 146 16 16 l/ hr. at 1,700 F., Water quench 132 164 6 25 hr. at 1,700 F., water quench plus age 154 175 4 8 %:1,700 F., 1VQ+7:1,400 F FG b 135 147 10 8 %:1,700 F., WQ+7:1,400 F.,

FC plus age b 149 8 5 %hr. at 1,700 F., air cool 133 150 16 14 hr. at 1,700 F., air cool plus age 147 152 6 8 B hr. at 1,700 F., furnace coo1 137 148 10 8 hr. at 1,700 F., furnace cool plus age 139 148 6 6 1g hr. at 1,8i0 F., Water quench 150 168 9 9 .1g hr. at 1,840 F., water quench plus age A 178 2 10 %:l,840 F., W'Q-.L7:1,4O0 F.,

FC c 145 155 3 8 .I/:1,840 F., WQ+7:1,400 F.,

FC plus age 1, 158 4 5 y hr. at 1,840 F., air cool 138 150 6 8 hr. at 1,840 F., air cool plus age a 140 149 2 2 ,1% hr. at 1,840 F., furnace cool.- 124 126 4 6 '1g hr. at 1,840" F., furnace cool plus age 127 132 2 6 24 hrs. at 1,200 F., air cool 134 148 10 11 24 hrs. at 1,200 F., oir cool plus age 140 151 8 5 Agingr heat treatment: 100 hours at 800 F., and air cool.

b Dual heat treatment: f hour at 1,700J F. and Water quench followed `by hours at 1,400a F. and furnace cooling'.

Dual heat treatment: 1/5 hour at 1,840D F. and Water quench followed 'by 7 hours at 1,400o F. and furnace cooling.

d Low temperature anneal.

8 responding data .for these alloys in the `beta :quenched condition. v

'TABLE II Average mechanical properties1 of Ti-Al-,Cu alloys las alpha-compound annealed Tensile Properties, p. s. i. X 1,000 Percent Percent Composition longa- Reduc- (Wt. Percent), tion in tion in VHN MBR 1 Bal. Titanium 0.2% Ultimate 1" Area Offset Strength Yield 113 119 15 122 129 18 125 134 12 146 158 18 111 .119 14 Y 124 127 10 6.5Al-3Cu 135 140 10 7Al-2Gu V 128 128 11 7Al3C11 132 144 10 BA1-5G11 158 169 8 5Ai-6Sn-3Cu. 139 153 14 3Al-12Sn-3Cu 153 163 14 12811-3011 108 124 14 l Results from both sheet and bar tensile bars included in average. Number of tests vary from one to three.

2 Minimum bend radius T.

TABLE lIn Average `mechanical properties of Ti-Al--Cu alloys-as vbeta quenched Tensile Properties, p. s. i. X 1,000 Percent Percent Composition Elonga- Reduc- (Wt. Percent), tion in tion in VHN MBR Bal. Titanium 0.2% Ultimate 1l Area Offset Strength Yield Broke in grips before yield.

`tion as given in Table II,-reveals that only the'most highly alloyed samples hadA strength-elongation properties as beta quenched which were comparable to those of the alpha-compound annealed material. vg'l'-he, ,dii"t`er ence between -the two ,heat `treated conditions is greater for these alloys `lhavingpa smaller.ltotal.'.alloy-.contnt It can thus be concluded that quenching or quenching and tempering heat treatments will be advantageous only for the most highly alloyed systems of this series.

The excellent elevated temperature properties, i. e., at 1000 F., of the Ti--Al-Cu alloys, as alpha-compound annealed, are shown by the test results given in l the following Table IV.

TABLE IV Elevated temperature (1000 F.)1 tenszle properties of Tr-(Al, Sb, Sn) Cu alloys as alpha-compound annealed Tensile Properties, Y p. S. l. X 1,000

Composition (Wt. Percent), Percent Percent Bal. Titanium Elonga- Rednc- Est. Ultimate tionin 1" tion in Yield' Strength Area 15 20 37 40 94 60 46 82 55r 72 se 9o 57 83 35 74 56 77 45 87 62v 82 A50 68 68 84 52 84 82 103 40 74 69 99 45 49 77 96 55 83 84 109 45 70 96 1.2 40 66 88 125 20 87 126 30 53 93 118 50 75 98 28 25 64 107 5 15 73 33 6l. 86 19 55 112 3 3 63 24 61 23 64 l Both bar and sheet test specimens were used.

10 From the foregoing dataof Table IV, alloyingis seen to result in a marked increase in elevated tempera `ture tensile strength. lThe beneficial effectsfof copper are shown to be quite dependent on the alpha promoter or aluminum content. vAt low aluminum contents the addition of copper results in'a large improvement inl elevated temperature strength. As the amount of aluminum is increased, however, tbe addition of copper is less benecial. For example, an alloy containing about 2 to 4% by weight of aluminum, isimpr'oved with copper additions up to 10%, while in an alloy containing 10% alutniriurn,v more than4% copper results in a reduction Ain elevated temperaturev strength. Since binary titaniumaluminum alloys containing moderate amounts of alumi num are dicultto fabricate, the beneflcial effects of copper additions thereto Von elevated temperature strength `will be ,particularly advantageous in wrought alloys of this series. Y

In order to obtain information on the thermal stability of TiAlCu alpha dispersoid alloys after various heat "treatments, they were exposed to 800 F. for 24 hours and the hardness and bend properties re-determined. The elect of this elevated temperature exposure is shown bythe data in the following Table V. From this data it will be seen -thatthe 'quenched' alloys were unstable, as wouldbe expected.,v The alpha-compound annealed and beta quenchedand tempered conditions were found to be they most stable. These results also indicate that a considerable amount of compound precipitation occurred during the thermal exposure in the alpha-beta quenched and tempered alloys.

TABLE V Effects of aging 24 hours at 800 F. on hardness and bend properties of Ti--Al-Cu alloys As Heat Change on Treated aging 24 hrs.

at 800 F. Alloy Heat Treatment VHN, MBR, VHN, MBR, 10kg. T 10-kg. T

load long. load long.

Ti-2Al-4Cu .Alpha-compound anneal 297 2.6 +15 +0. 2 Ti-4A1-4Cudn 347 3.6 -5 -0. 2 Ti-4Al-2Cu. ---.-dn 330 B. 6 +10 0. 6

Average. .--.-n 325 3. 3 +7 -0. 2

Ti-2Al-4Cu Alpha-beta quench 353 4.6 +11 +0. 7 Ti4Al-4Cu rin 385 4. 4 +50 +5. 6 Ti-4Al-2Cu dn 361 2. 8 +3 0. 2

Averagedo 366 3. 9 +21 +2.()

Ti2Al4Cu Alpha ,beta quench+temper 328 4. 2 +47 +1. O Ti-4Al-4CLL dn 385 8. 3 +18 +1. 3 Ti4Alf2Cu dn 340 3. 6 +44. +0. 4

Average. dn 351 5. 4 +36 +0. 9

Ti-2Al-4Cu Beta quench 360 5.0 +12 +5. 0 T-4Al 4Cu d0 351 3. 4 +55 +0. 6 Ti-4A1-2Cu dn 317 3.0 +37 +1. 6

Average. .--.-dn 343 3. 8 +35 +4. 4

Ti2Al4Cu Beta quench temper..- 385 9.4 -11 f l+0. 4 T14A14Cu dn 369 10. 0 +18 0. 0 Ti-4A12Cudn 362 3. l +22 0. 0

Average. n 372 7. 5 +10 +0. 1

.TABLE yVI ,Results of stress-ruptured testson Ti-Al-Cualloys as alpha-compound annealed Ccmposition Test l Applied Rupture'liimef Percent (Wt. Percent), Temper- Stress, -Hrs. Elonga- Bal. Titanium ature, F. p. s. i. tion 6Cu 800 '80,000 Broke on load- 15.

4in-sou son i4. 4A1-6Cu 800' 23. 4A1-8Cu. 800 5. -4Al-10Cu 800 l 6.

l Test discontinued after 362 hours.

The effect of substituting the alpha promoters tin and .antimony for yall or part of aluminum in the copper- :bearing alpha dispersoid alloys is shown by the test results in the followingTables VII, VIII and IX.-

TABLE VII Average mechanical properties of Ti-,Cuplus A1,Sn,Sb alloys ASA ALPHA-COMPOUND ANNEALED Tensile properties, p. s. i. 1,000 Per- Per- Composition (Wt. cent cent Percent), Bal. Tita- Elon- Reduc- VHN MBR nium 0.2% Ultimate gation tionin Offset Strength in 1" Area Yield 40u-GS11 87 98 20 46 277; 2.1 6Cu-6Sn- 90 112 14 25 6Cu- 6Sb 97 119 10 16 4Cu- 4Al-4Sn 134 146 13Y 25 `366` 5.6 4.5Cu 2Al-3Sn... 101 113 6 17 6. 6 4C`u- 2A1-3Sn- 0.1Be- 97 113 12 19 1. 5 6Sn-4Al-4Cu 131 143 l0 14 6Sb-4Al-4Cl1.. 131 138 11 A11 9Sn3Al-4Cu 124 136 13 26 9Sb-3AI 107 111 17 33 9Sb-3Al-4Cu. 136 142 7 v 15 12Sn-2A1 95 108 20 38 12S11Y 2Al--4Cu. 132 142 12 -18 12Sb- 2A1... 106 110 18 30 12Sb-2Al-4 132 140 14 23 15S11-1A1 99 107 l17 39 15S111Al 135 144 l17 26 l5Sb-1Al... 109 112 16 "29 AS BETA QUENCHED 4Cu-6Sr1 121 146 2 3 370 ytSCu-6Sn 151 195 6 13 .430.

6Cu6Sb 140 184 4 12v 388 7. 1 4Cu-4Al-4Sn 151 170 5 13.

ELEVATED' TEMPERATURE (1,000l vF.) TENSILE PROPER- TIES AS ALPHA-COMPOUND ANNEALED 4Cu-4Al-4Sn 63 -82v 40- 87- CuVlOSb 50 66 48 67 ,TABLE 'VIII l`fEyeet ollagz'ng 24 hours at 800 F.,on the hardness and bend properties of Ti-Cuplus Al,Sn alloys as alphacompound annealed -Resalts of stress-rapture tests on a -Ti--Cu--plus Al,Sn alloys as alpha-compound annealed 'Test Applied Rupture 'Percent Composition (Wt. Percent), Tem- Stress, Time, Elonga- Bal. Titanium pegature p. s. i. hrs. tion 4Cu-4A14Sn .800. 80, 000 1 407. 8 '(1) 800 85, 000 l 282. 5 0.167

Y800 v95, 000. l 343 4. 6

l These results from one specimen. A iter 407.8 hours at 80.000 p.=s. i. the stress was raised t0 85,000 p. s. i. After 282.5 hours at 85,000 pas. i., the stress was raised to 90,000 p. s. i. The test was discontinued after 343 hours at 95,000 p. s. i.

These data show that tin or antimony may be substituted for aluminum on a replacement basis of about 3% Vof the former for 1% of the latter, on a generally comparable basis as regards properties of the resulting alloy. A definite increase in density results from such substitution of tin or antimony for aluminum. However, such substitution does improve fabricability somewhat, tin being superior to antimony in this respect.

The alpha-compound annealed properties of these Sn, Sb4 alpha substituted alloys, are shown to be reasonably analogous to the Ti-Al-Cu alloys in the same condition of heat treatment. Here again tin appears to be slightly superior .to antimony in its elect. Such alpha substituted alloys are shown to have in the beta quenched condition approximately the same mechanical property levels as the equivalent Ti-Al--Cu alloys, and this sameV equivalency is shown to hold at elevated temperatures also. The stress-rupture tests show that even further awr improvement in stress-rupture properties is obtained by increasing the aluminum equivalent, alpha strengthener content above 4%. The aging tests show that these alpha substituted alloys have good thermal stability in the alpha-compound annealed condition.

The effect of replacing all or part of the copper in the alpha dispersoid copper-bearing alloys above discussed by the other compound formers beryllium, nickel, cobalt and silicon, is shown in the following Tables X to XIII,

Tensile Properties, p. s. i. X Per- Per- Cornposition (Wt. ,000 cent cent Percent),Bal. Tita- Elon- -Re- VHN MB- nlum gaton duction R, 0.2% Ultimate in 1" in T Onset Strength Area Yield 68 74 26 51V 163 1. 2. 58 80 19 3l 158 1.1 52 72 29 46Y 188 0.25 65 83 20 39 218 1.4 76 95 17 31 226 .2.1 70 86 28 44 264 2.1 84 96 17 27 256 V2.3 84 96 20 42 297 2.6 84 98 20 32 254Y 2.2 88 104 16 24 258 vr2.3 117 136 10 10 835Y 5.4 127 146 9 12Y 356 3.8 96 107 6 16 316 6.8 108 122 4 12 331 .8.7 103 116 14 25 95 106 16 38 317 V2.4 2A13.9Cu-0.15Be 97 113 10 18 V322v Y.4.6. .4A15Cu-1Mn 128 135 15 37 372 2. 8

TABLE XI I Average mechanical properties of Ti--Alplus Cu, Be,

inc., which also show the effect of replacing part of th copper with the relatively sluggishly eutectoid element manganese.

TABLE X Average mechanical properties of Ti-Alplus Cu, Be, Ni, Co, Si alloys as alpha-compound annealed C0, Ni', Si alloys as beta quenched Tensile Properties, p. s. i. X Per- Per- Compositlon (Wt. ,000 cont cent Percent),Bal.Tita Elon- Re- VHN MB- nlum gation duction R,

0.2% Ultimate in 1" in T Offset Strength Area Yield 67 88 17 21 191 0.8 80 99 8 8 273 4.8 80 96 8 10 272 2.5 84 108 6 12 260 1.4 99 120 6 15 268 3.0 96 111 8 10 330 5.0 91 114 11 24 255 2.2 105 116 11' 15 311 3.3 113 133 4 8 303 4.0 4Al-0.5Be 132 152 2 7 382 )10.0 0.5Sl 193 0 1 GSi 215 1.0

5CH-0.5131n 473 10.0

2A15Cu0.5Be 473 10.0 4.41-5Cu-1Mn 198 204 1 3 469 10.0

30u-1 nSi 291 3.1

5Cu-0.5Si... 359 6.0

2A1-3Cu-1Si 354 5. 5

377 5.3 275 4.8 382 y3.2 400 4.4 433 8.8 446 9.2 433 5.6 381 2.6 236 1.8 2A11.0Si* 326 2. 4

These samples quenched from the alpha-betalteld. v i i 14 TABLE X11 Elevated temperature (1000 F.) tensile properties of Ti-`-Alplus Cu, Be, Ni, Co, Si alloys as alpha-compound annealed Tensile Properties,

p. s. i.X1,000 -Percent Percent Composition (Wt. Percent) Elon- Reduc- Bal. Titanium gation tion in Est. Ultimate in 1 Area Yield Strength Strength Be 17 22 55 83 2Al-0.5Be 22 36 46 74 4Al-4Co- 55 44 90 4Al-6Co 57 44 87 10A1-5C0 43 8l 5 5 4Al-4N1 47 42 64 4Al-6Nl.-. 53 53 84 10A1-5 N'l 59 79 15 30 4Al-0.5S 60 26 57 4Al-5Cu-1Mn 60 76 46 97 For comparison:

4Al-4Cu..- 50 60 46 82 4Al-6Cu. 55 72 36 90 BA1-40u- 96 30 83 10M-40u 126 30 53 TABLE XIII Eect of aging 24 hours at 800 F. on the hardness and bend properties of alpha-compound annealed 'I1- Alplus Cu, Be, N1, Co, S1 alloys Before Aging Increase Due Composition (Wt. Percent), to Aging Bal. Titanium MBR I VHN MBR l VHN 2Al-0 5Be 2. 5 241 -0. 5 +20 2. 4 272 +0. 1 -78 *May have been annealed in alpha-betacompound eld instead 0f .alpha-compound held.

Referring to the mechanical properties for these alloys in the alpha-compound annealed condition, as given in Table X, it will be seen that all of these alloys except Ti-Al-Ni and Ti-Al-Cu-Be are found to be equivalent to the Ti-Al-Cu alloys. The Ti-Al-Be and Ti-Al-Si alloys are shown to be equivalent to the low alloy Ti-Al-Cu alloys, while the Ti-Al-Co and Ti--Al-Cu--Mn alloys are shown to'be equivalent to `the high alloy Ti-Al--Cu alloys.

Referring to the beta-quenched properties of Table XI, it will be seen that all of these alloys undergo quench hardening, the Ti-Al-Be and Ti-Al-Si alloys less so than the alloys containing nickel, cobalt and manganese. The combination of high hardness and good bendproperties exhibited by the Ti-fAl-Co alloys as quenched, appears particularly promising.

With respect to the elevated temperature tensile properties of Table XII, it will be seen that on an equal weight basis, both cobalt and nickel are somewhat inferior to copper in their effect on elevated temperature properties. The addition of a small amount of manganese is benecial in this regard. The Ti-4Al-0.5Si alloy compares favorably with the Ti-4Al-4Cu alloy in strength but not in ductility. l The-.aging tests of Table XIII show that the Ti--Al--Be, Ti-fAl-Si and Ti-Al--Ni alloys possess reasonably goodthermal stability. The Ti-Al--Co and Ti-Al- .Cui-M11 alloys. however, show denite indications of VV15A instability. These resultsindi'cate that manganese promotes thelreten'tioniof vthe beta yphase due toits beta sfavbrlrzmg characteristics.

The foregoing test results indicate that neither silicon nor beryllium canrcompletelyv replace copper in alpha dispersoid alloys, since'thetotal-amount which can be added 1s small. The-'replacement of a small part of the copper 5 according to the invention, as rolled and annealed at about 1550 F., are given in the following Table XIV:

TABLE XIV Me'hanical properties of TIL-(Al, Sb, Sn)-(Cu, Be)

' alloys as annealed (COMMERCIAL PURITY 'Pi-BASE) Tensile lroperties,

p. s. 1. 1,000 Percent Y Y Elon- Min. Vickers A1 Cu Other gation Bend, T Hardness:-- A 0.2%. AUltimate in ll Long.-

. YOffset: Strength Yield '5 97 101 22 2. 5 334 -5 0. 5 110Y V117 20 2.5 316A f5 1 107 115 20 2. 5 305 5 1. 5 114' 122 18 2. 7 318 5 2 110 119 16 2.' 5 312 5 3 123 133 12 2.4 354 5 5 Y 141 152 12 5.1 347 5 0.5. 0. 127 130 14 '1.7 341 5 1 0. 126 132 17 2.`4 `325 5 1. 5 0. 129 135 2.3 343 k 5 2 0.10 130 135 15 4. 9 350 5 3 0.10 140 147 13 4. 9 378 5 5 0.10 147 156 11 Br 375` 5 0. 5 0.20 127 136 19 2. 5 363 l.5 1 0.20.. 132 137 16 2.3 357 `5 1.5' 0.20 120 132 19 1.7 371 5 2 0.20- 142 149 12 5. 2 v364 5 3 0.20- 153 153 10 4. 9 353. l 5 5 0.20. 153 160 a Br 372 l5 0.5 0.25 14o 141 12 1. 7 .345

5 1 0.250-- 136 139 15 2. 5 334 5 1. 5 0.250-- 139 145 14 2 5 343 5 2 0.250-- 138 143 10 5.1 376 5 3 0.250-- 139 149 9 Br 388 5 0. 5 0.10--- 113 118 14 2. 4 350 A5 1 0.10 14 2.3 343 5 1. 5 0.10 17 2. 7 369 Y A5j 2 0.10 14 4.8 349 5 g 0.5 0.20 17 2.4 354 5 1 0.2 16 2. 5 385 5 1. 5 0.2 14 3.6 380 e 5 2 0.2 11 4.9 364 `5 0.5 13 2. 5 346 5 l 1.5 12 2.6 375 5 0. 5 13 2. 3 370 Y'5 1 11 5. 2 383 "5""1.5 1 6.1 390 4 2. 5 9 3. 4 349 4 2. 5 5 6. 7 347 4 2. 5 12 2.8 344 4 2. 5 8 3. 7 361 4 2. 5 8 3. 5 333 4 2. 5 5 3. 9 343 4 '2. 5 l1 4. 1 350 4 2. 5 1 3. 6 346` 4 2. 5 9 3.6 344 5 11 311 5 7 349 A.l Sn

5 2. 5 137 139 10 3. 8 400 5 55 72 19 3. 1 269 i 10 75 94 15v 2. 0 276 15 100 112 9 2. 9 322 10 98 108 22 1.9 340 10 10u-'0.10: 94 108 14 3.4 347 10 10u-01N: V114 128 11 6.3 375 (COMMERCIAL PURITY Ti-BASE) 10 2.50121.` 96 106 18 2. 6 325 10 50u.-- 122 131 13 3. 6 A 342 10Sb2.50u 129 136 8 2. 3 344 (IODIDE Ti-BASE) 10 2.5Cu-1'.'5Fe.^- 117 132 11 2. 0 357 10 250g-725ML--- 129 136 7 1.9 331 10 2.5Cu4.5Mz1 139 '149 10 1.5 351A (IODIDE 'Il-BASE) Tensile Properties,

p. S. l.X1,000 Pcent in V. k

on- 1c ers u Co other 02% gation Bleggg'rl Hardness Oset Ultlmate 1n 1 Yield Strength 53 77 22 0. 7 202 4 88 101 19 1. 7 304 4 2 2.5 Or- 135 145 9 2. 6 353 4 2 2.5 Cu 113 138 8 7. 2 341 4 2 1.5 Fe- 118 142 5 Br 368 4 2 2.5 Mn-- 117 138 12 1. 9 348v 4 2 2.5 M0 128 142 8 1. 7 372 4 2 2.5 N1.-- 96 117 5 6. 2 296 4 2 51a- 108 127 4 7. 1 335 4 2 5V. 154 160 5 2. 3 368 4 2 2.5 117 135 8 6.0 345 4 2 5Cb 139 147 6 3. 8 353 1 10Sn 90 106 19 2. 7 289 2. 5 10Sn 108 125 11 2. 6 317 5 10511-.-- 111 124 3 Br 357 2. 5 10Sb 108 119 11 5. 4 309 2 10Sn2Ni 97 118 12 4. 9 325 2 10Sn-2.5Cu 114 136 5 Br 339 2 10Sn-2.5Mn-. 147 '153 9 4. 1 354 (IODIDE Ti-BASE) (COMMERCIAL PURITY T BASE) Al Ni Our researches have established that Ti--Al-Cu alloys according to the invention containing about 4 to 8% aluminum and 2 to 4% copper have excellent properties as forging alloys.

A typical alloy of this type having the composition Ti-7Al--3Cu was forged at 1800 F.

beta field (1700 F.), or from the beta eld (1840u FA.)l v

As is also shown, the greatest response to heat. treatment is obtained by quenching from the beta field, which produces an increase in strength of about 35,000 p. s. i.

as compared to furnace cooling from the beta field. A better combination of strength and ductility is obtained by quenching from the alpha-beta field, where the strength is almost as high as that obtained by quenching from the beta field, and the reduction in area is above 20%.Y Air cooling from the alpha-beta or beta eld causesv a small increase in strength (about 15,000 p. s. i.), while furnace cooling from the alpha-beta field has little effect on strength. Furnace cooling from the beta eld produces some loss in strength as' compared with the alpha-compound, furnace-cooled condition. quenching from the beta eld and tempering to produce a random distribution of very small compound particles in an acicular.4

alpha matrix results in an increase in strength of about '10,000 p. s. i. as compared' to the alpha-compound any nealed condition. The tempered alpha-beta quenched specimen does not show this strength increase, probably ature tensile properties of the Ti-7Al--3Cu alloy b oth prior and subsequent to aging for 100 hours at 800 F. and also after subjection to creep exposure for 100 hours at 750 F. at a stress of 65,000 p. s. i.

TABLE XV 20 one employing the highest final annealing temperature (1400 F.) is the leasteiective from the stability standpoint. Since the other three anneals finishv at lower temperatures, it appears. that a gain in stability is obtained by equilibrating the alpha and compound phases at a temperature closer to the creep exposure temperature. It is to b e noted from the.. data of this table, however, that although excellent thermal stability is obtained with the second annealingcycle listed, the stability after creepv exposure is impaired with; this as Well as. Vwith. all of the other annealing cycles.

The elect on stability of substitutingV Various other alpha promoters, such as one orV more of'tin, antimony and zirconium for part or all of the aluminum in the Creep data and subsequent room-temperature tensile-test1 data of a Ti-7A1-3Cu alloy Creep Data. Room-.Temperature Tensile Properties e.

. Tensile Properties, Annealing Heat Treatment Residual Min. p. s. l. X 1,000 Percent Time, Temp., Stress, Strain, i Creep E1onga- Percent. hours F. 1,000 Percent Rate tion Reduc? p. s. i. in 1 per hr. 0.2% Ult. in 1" tion in 1 Onset Str. Area Yield 6 hr. at 1,400 F., plus 2 hr. at 800 F., plus 24 hr. at 1,200 F., and air cool to room temp None 125 138` 10' 16 Do 100 132 147y f 14 16 D0 100 126, 131 3 5. 6 hr. at 1,400 F., plus 24 hr. at 1,200 F., and air e001 to room None 123 137 14 23 100 133 146 16 16 Do 100 126 132 6 S. 2 hr. at 1,400 F., furnace cool to 1,20 F., hr. at 1,200 F. furnace e001 to 1,000 F hold 16 hr., and air cool None 137 147 6 8 D 100 800 None 141 150 8 10.

None 132 144 10 8 100 800 None 140 148 10a 9 100 800 0.247 0. 0009 (b) 100, 2 4.

Transverse bar specimens. bBrittle fracture occurred before yield.

` It will be seen from the foregoing data, that of the Ti-7Al-3Cu alloy is; four alpha'compound stabilizing anneals employed, the following Table XVI:

TABLE XVI Rham-temperature and 1,0001"w tensile properties of alphgzgdspersoid alloys [As annealed 7 hours at 1,400 F., furnace cooledto 1,100" F., held 16 hours at,1,100 F., and air-cooledto room temperature] shown by the test results in the Exposure Data. Tensile Test Data Tensile Properties, Comp. Percent, Bal. Ti Min. p. s. i.X1,000 Percent Stress, Percent Creep Test. Elonga- Percent Time, Te1IE1p., 1,000 Plastic Rate Temp., tionin Red. in MBR, hours4 pgs. i. Strain per hr. F. 0.2% Ult. 1 Area T, L

Y Oset Str.

Yield None 75 12s 137 14 Mii-3G11 Bar (T) 1--, 100 800 None 75 133 146 16 100 750 65 0. 22 0. 0009 75 126. 132 6 one TABLE XvI.-Con1inued Exposure Data Tensile Test Data w Tensile Properties, Comp. Percent, Bal. Tl Mln. p. s. i. 1,000 Percent; v

Stress, Percent Creep Test. E long'a- Percent Time, Temp., 1,000 Plastic Rate Temp., twg'm Red-m MBR, hours F. p. s. i. Strain per hr. F- 02% Ult. 1 Area T, L

Offset Str. Yield 75 133 '154 s 9 75 151 164 4 7 75 147 159 13 23 75 152 166 7 11 75 144 157 5 10 75 134 .145 10 3 300 None V75 139 140 2 3 14 300 65 0.10 0.0004 75 132 0 0 A15Sn5Zr30u Bar (T).--- z142 300 65 0.13 0.0012 75 147 143 2 1 120 1,000 None 75 143 0 0 Nene 1,000 34 109 13 22 None 75 123` 139 14 17 113 300 Nene y 75- I 137 v 143 4 3 117 300 65 0. 153 0. 0003 75 136 0 0 m-ezr-sou Bar (T) 120 1, 000 None 75 139 149 12 16 99 300 65 0.167 0.0014 75 123 1 3 93 1, 000 15 0. 130 0.0012 75 141 154 14 12 Nene .75 123 123 12 13 118 800 None 75 129 135 12 7 139 300 0. 13 0. 0005 75 127 130 9 3 4.41-12zr-3ou Ber (T) 115 1, 000 15 0. 25 0. 0016 75 123 130 14 11 97 1, 000 15 o. 20 0. 0003 75 115 126 2 4 Nene 1, 000 73 91 34 69 Nene 75 134 142 14 20 104 300 70 0.06 0. 0003 75 137 146 7 15 93 300 65 0. 09 0.0006 75 137 145 10 13 4111-12zr-3ou 4 Bar (T) 104 1, 000 10 0. 10 0. 0005 75 135 141 9 13 99 1,000 15 o. 19 0.0012 75 134 20 S22 13 11 5 None 113 300 Nene 75 137 143 4 8 117 300 0. 15 0. 0003 75 136 0 0 6111-6zr-3ou13ar (T) 2144 300 65 0.16 N11 75 141 150 5 7 120 1, 000 None 75 139 149 12 16 Nene 1, 000 32 105 26 64 Nene 75 91 102 13 40 Sh L 142 300 50 3g. 3. 75 111 114 10 21 iAl-SCu eet 25 800 65 74 1,000 10 1.23 0.0163 75 30 92 11 36 Nene 75 74 34 21 35 2.5 0.1 300 17.10 2. 251-3011 sheer L 0.1 300 20.00 2.3 1,000 10 2.11 0.0253 75 03 73 13 44 2.6

N4@ 72 13 13 3 3 300 50 0. 33 0.0076 12HR-30u Sheet (L) i 99 1, 006 10 0. 63 0. 0043 75 11o 125 11 32 2,1

N... e .3 19 3 5 102 300 50 0. 13 0. 0017 12%3011 Sheet (L) 136 1,000 15 0.56 0.0052 75 100 112 10 23 3.2 N2 3 3 3 3 e 35 96 300 3o 1.44 0.0 3C Sheet (L) i 96 1,000 5 0. 60 0. 0031 12; 12g; 1g 3g 1. 9 N 5A112Zf--6C11B* (L) Nggg 75 150 157 12 21 y 300 65 0. 9 N11 75 151 162 7 9 93 1,000 25 0.436 0.0033 75 150 159 12 10 Bar-(Thu. 69 1,000 15 l 0.056 0.0016 75 145 156 10 23 99 300 75 0 0.0060 75 A151 v 161 10 10 300 96 121 24 57. ne; (L) i 1,000 90 103 2s 60 1 Type Specimen. 2 120 hrs. at 1, 000 F. stabilization anneal prior to creep exposure. 3 Average properties of 7, 1b. mgots.

In the data shown by the above tests, three levels of Over, the Ti-l2sl-3Cu S almost Completely Stablestronger than the Ti-4Al-3Cu composition.

Ialpha stabilizer additions were made to the Ti--3Cu alloy. Tin and antimony were substituted for part or all of the aluminum in the 3:1 ratio above mentioned. In other alloys shown by the data, zirconium is added in varying amounts as an alpha promoter, while in still Aother data, molybdenum was added as a beta-isomer- Vphous promoter.

The data given in the above table shows that the substitution of tin for part of the aluminum at the 7% (equivalent) aluminum level does not impart greater 70 stability. Higher strength is obtainable with higher tinto-aluminum ratio, but the stability remains about the same. At the 4% (equivalent) aluminum level, the Ti-12Sn-3Cu and the Ti-12Sb-3Cu alloys are More- Also both the Ti-4Al-3Cu and Ti--12Sb-3Cu alloys are more stable than the alloys containing higher alpha stabilizer additions. The two alloys with no or low alpha stabilizer additions (Ti-3Cu and Ti-2Al-3Cu) also show high degrees of stability, as shown particularly by the bend duetility Values. The addition of 2% molybdenum imparts high strength with adequate ductility, but the resulting alloys are unstable. v

A` comparison of the elongation data given in the above table suggests that two different types of embrittlement are encountered in these alloys. The rst, due to copper precipitation, is not too serious, resulting in only slight loss in ductility. Embrittlement of this type is seen quite clearly infcompositions such as Ti-3Cu,

ity is about equally great at 800 and 1000" F. 'Ihe secmolybdenum evidently increases the tendency toward em.-V brittlement, as shown bythe tendency for Ti-A'l--f 2Mo- 3Cu to be embrittled both at-800 and 1000 F.

Presumably the detrimental effect .of molybdenum is .con-.V

nected with precipitation of TizCu.

The replacement of aluminum plus ltin and/or ant-irnony-i I by zirconium is effective in removing the tendency toward ordering. For

aluminum equivalent) `is extremely unstable,VV Ti-6Al-6Zr-3Cu (6% aluminum equivalent) 'is un-Y aluminum equivalent) is stable. Also, since 'l`i 7tAlf3;CuV

was embrittled only at 800 F., it is apparent that zirconiurn -additions lower the permissible alum-inum-plus-tin and/or antimony content slightly, possibly as much as 1% aluminum equivalent.

Alloys according to the invention containing about 5..to.f l

8% copper and 8 to 14% tin and preferably about 6 to 7% copper and 10% tin have excellent properties for use in sheet form, i. e., in the airplane industry forv'ving coverings and the like. As alpha-compound annealed. at about 1475 F., these alloys are relatively soft, have relatively low yield strength and excellent bend ductility and thus can be easily formed and bent to shape. In addition, they undergo a marked strengthening and hardening when subjected to relatively slow cooling from temperature being equivalent to a cooling rate of about 10-15 seconds Yto cool from 1540 F. down to 1000 F. (3G-50 F. per

second), (3) even with very gslow cooling rates, as depicted at greater distances from the quenched end, appreciable hardening occurs.

The low as-quenched hardness is indicative of either soft martensite formation or soft retained beta, or both. Although thestructure at the quenched end appears to be chiefly martensite, it is probable that some retained beta is present also. The structure at the peak hardness is vpredominantly martensite with about 5% eutectoid structure,

(some omega-hardened beta also may be present which cannot be seen vin the structure). A complete eutectoidy structure is not obtained until about 2" from the quenched end. l

To check the stability and the effect of itempering on the various structures above referred to, additional speci- --mens were endquenchedpand were re-heated for 16 hours at 800, 1000, 1100 and `1200 F., respectively, and tested for hardness with resultsalso as shown `in Figure 1 lby the graphs, respectively, so designated. It is evident from these additional Vtests that the martensitic structure near the quenched end is not stable. Thus, as shown in this drawing, aging at 800 F. causes an increase in hardness even above the peak hardness to a distance of about l" from the quenched end. Beyond this distance, where the structure is predominantly eutectoid, there is no r,significant change in hardness. Tempering for 16 hours at 1.100051?. Vcauses a Vdecrease in hardness throughout the wpeak hardness range, .giving a fairly uniform hardness along the length of the specimen. Similarly, tempering at f 12.00 F.7causes a softening over the entire length of the above the beta-transus, i. e., from about 1500-1550 .F-

or above. The fact that these alloys can be strengthened by this relatively simple and short-time heat treatment consisting in slow air cooling from above the criticalgis of great practical advantage, because it permits the use of jigs to hold the fabricated parts in position, such as a `Wingfoil and `air frame, during the strengthening `heat This relatively short-time heat treatment retreatment. quired for strengthening these alloys, standsin marked contrast .to the long .periods of heat treatment required for alloys which are hardened and strengthened by the agehardening procedure. These Ti-Sn-Cu alloys also possess the further advantage that after being strengthened by the air cooling heat treatment aforesaid, they may be thermally stabilized by a relatively simple tempering treatment lin air at about l000-1100 F. with considerable leeway permitted in both time and temperature of vsuch treatment.

The air hardening characteristics .of these alloys is well illustrated by resort to the Jominy test according to which a bar of the alloy is heated above critical, for example, at 1540 F., and thereupon end-quenched in water. The results of such a test carried out on a Ti--10Sn-7Cu alloy are shown in Figure 1 `of thea-nnexed drawings by the graph marked `as quenched. In this drawing,xthe

ydistance from the quenched end in inchesis plotted as abscissae while the Vickers hardness of the quenched specimen at any given distance from the quenched end as plotted as ordinates. The important features of this endquenched graph as thus depicted are: (l) the alloy is soft in the quenched condition, (2) 'the peakl hardness occurs at a distance of about 1%," from'the quenched end, thus specirnento a substantially uniform hardness. These data show that in order to obtain avstable structure, alow enough cooling rate should be used to secure a eutectoid structure; and also that it is possible to temper the highvhardness martensitic structure to a hardness about the sameas the eutectoid structure.

It has heretofore been shown in U. S. Patent 2,769,707

-to M. B. Vordahl, that tin slows down the beta decomposition reaction Irates in alpha-beta, Ti-base alloys. Our

' investigations indicate that tin acts similarly in titaniumcopper alloys, las evidenced by the lominy end-quench .tests shown in Figures 2 and 3 ofi-the accompanying drawings, Figure 2 illustrating the effect on hardenability of varying the tin'content at the 2Ti-7Cu level, and Figure 3 illustrating theettect .of varying the Copper Content 0n hardenability of the-Ti-lOSn base-alloy. It will be seen from these graphs that maximum depth of hardening occurs at the .7% copper (eutectoid) rlevel. The shallower depth of hardening which occurs with both hypoand hyper-eutectoid alloys is believed to be caused by proeutectoid alpha or compound particles acting las nucleation sites for the eutectoid reaction. It is apparent from these graphs that to obtain maximum depth of hardening, the copper content should .beof the `eutectoid v("7 composition and the ktin content should be kept Vas high as possible.

Referring to Figure 2, it will be seen that as the tin content is decreased, the hardness peak shifts `progressively toward the quenched end, so that these Ti-Sn-Cu alloys containing about 6% and under of tin, obtain their maximum hardness on quenching from above the critical. Conversely, the fact that the hardness peak shifts progressively further from the quenched end Vas the tin contentis increased, evidences the fact above mentioned that thetin .addition is slowing down the beta Vdecomposition.reaction rate.

The Ti-lOSn-7Cu alloy is relatively soft in 'the alphacompound annealed condition and is capable of being 25; hardened by air-hardening treatment to useful strength levels. Typical tensile properties of this alloy in the alpha-compound annealed condition are:

The effect on tensile properties of heat treating this alloy from above the critical is dependent upon the solutionannealing temperature and upon the cooling rate, as shown by the data in the following Table XVII:

The beta-transus for this alloy occurs at about 1500 F. andl it will be seen from the data in the above table that the maximum strength is imparted to the alloy when it is cooled from the heat treating temperature to below 1000 F. in about ten seconds. It was indicated above by the end-quench graphs shown in Figure 1, that tempering for 16 hours at 1000 F. tends to overcome the hard- -ness variations caused by the diiferent cooling rates ncorporated in the end-quench tests. Also variations in tensile properties tend to be smoothed out by tempering treatments.

In the alpha-compound annealed condition, the stability of the Ti-10Sn--7Cu alloy is good in both unstressed and stressed conditions as shown by the data in Table XVIII:

TABLE XVIII TTSS Exposure Tensile Properties,

' p. s. i. X 1,000

Percent Percent Elonga- Reduc- Annealing Heat Treatment Stress, Residual 0.2 tion in tion in Time, Temp, 1,000 Strain, Otiset Ult. 1" Area hours F. p. s. i. Percent Yield Str.

6 hr. at 1,400 F. plus 2 hr. at 800 F., plus 24 hr. at 1,200 F., and air cool None 100 111 14 23 Do 100 800 None 99 106 8 37 100 750 l 1. 6l 110 116 11 20 None 100 110 12 30 100 800 None 102 110 11 25 t 1 100 750 b 3.04 110 113 4 17 a hr. at; 1,200 F., furnace cool to 1 16 hr. at 1,000 F., and air cool.. None 100 109 10 29 Do 100 800 None 106 116 13 21 Do 100 750 1. 64 111 116 9 19 16 hr. at 1,400 F. and air cool-- None 109 126 10 24 Do 100 800 None 111 124 12 30 Do 100 750 d 0.55 120 124 4 0.00357 percent per hour minimum creep rate during TTSS exposure.

b0.0142 percent per hour minimum creep rate during TTSS exposure.

TABLE XVII Tensile properties of T110Sn7Cu alloy in several heattreaied conditions Heat Treatment Tensile Properties,

p. S. l. X 1,000

Percent Percent Elonga- Reduc- Coollng 0.2% tion in tion in Solution Anneal- Rate, Oiset Ultimate 1" Area lng Temp., F. see. Yield Strength Strength l Time to cool from heat-treating temperature to below 1,000 F.

percent per hour minimum creep rate during TTSS exposure. percent per hour minimum creep rate during TTSS exposure.

Althoughthe uniform elongationof this alloy in the.;4

annealed condition is only about 2% at room temperature, it is increased to about 10% on heating to the relatively low temperature of 200 F., thereby to permit an excellent stretch forming.

As pointed out in U. S. Patent 2,669,514 to vFinlay et al., tin in excess of about 5%, imparts to Ti-base alloys, excellent contamination resistance to atmospheric gases at elevated temperatures. This is true of the Ti-lOSn-7Cu alloy, which, in addition, has excellent oxidation resistance at the solution-annealing temperature of 1525 F. and at the annealing temperature of 1400 F. Less than 0.025 mil of the metal is lost in two hours at 1525 P. and only about 0.03 mil is lost in sixteen hours at 1400 F. In terms of metal loss per year, this amounts to about 0.015 per year at 1400 F. and 0.075 per year at 1525 F.

The effect on room temperature and tensile properties of varying the tin and copper contents of these Ti-Sn-Cu sheet alloys, in both the alpha-compound annealed and heat treated conditions, is shown lin the following Table XIX. The succeeding Table XX shows Ti-12Sn--5 Cu--2Mo alloys.

zcfai Z? y TABLE' XX l R'oomemperatzgkrje fensz'le test Yprodpenz'esv of various heartreatable alpha-dispersod sheetalloysf [A'lphe-compound annealed plus solution heat treat a's indicated] Ti-8Sn-7C`u rIensil'e Properties, f

p. s. i. X 1,000- Percent Heat Treatment (20 Minutes at f VHN Percent Reduc- Temperture) (Surface) Elongation 1n 0.2% Ultimate Ition in 1" A rea Olset Strength Yield Nimes: 265 104Y i 114g 1s 31 None 275 111 1v19` 10 26 1,500 F., slow air 0001..-- 328 139 152 5 15 1,525 F., water quench 320 91 152 6 12 1,525 F., water quench plus tem.

per b 370 145 159 1 7 1,525 F., fast air cool 445 180 1 6 1,525 F., fast air cool plus temper b 390 163i, 175'; 1 10 1,525a F Slow air C001 360 162 173 1 10 525 F slow air cool plus temper b 385 170 177 1 8 1,525 F., very slow air cool.- 330 147 162 3 10 1,525 F., Very slow air co Y temper b 330A Y 142 y 156A 3 11 1,550 F., slow air cool 395 184 195 1 8 T1-10S11-7Ou 1,550 F., slow air cool 324 129 142 7 27 Du'plee spe'imfen 3 130 142 6 13 1,525 F., slow air cool... 336 142 155 5 17 Duplicate specimen 339 137 153 7 18 1,550 F., slow air cool. 386 168 175 3 14 Duplicate specimen 396 167 V185 3 7 Ti-llSn-Gu None 290 112 122 11 32 1-,525 F., slow air cool 416 106 1,525 F., slow air cool plus temper b, 403 149 Ti-12Sn-5Cu 285 106 118 11 295 105 119 11 29 1,500 F., slow air cool. 340 143 154 4 22 1,525 F., Water quench 345 108 159" 7 32 1,525 F., water quench plus temper b 400 162 168 1y 4 1,525 F., fast air C001 390 169 186 f 2 13 1,525 F., fast air cool plus temper b.-. 370 153 165 3 5 1,525 F., slow air cool 365 157 168 3 11 1,525 F., slow air cool plus temperb 370 146 157 3 11 340 135 147 6 15 Ti-12Sn-4Cu Ti-12Sn'-3Cu N 011e 295 94 110 11 28 1,300 F., furnace 0 F.,

hold 1 hr. at 1,000 F. and air cool.- 310 95 100 14 30 Ti-14Sn-5Cu None d 362 Ti-14Sn-3Cu None 337 124 136 3 14 (i) Alpha-plus-compound annealed 16 hrs. at 1,400 F. and air cooled. `(b) Tempered 2 hrs. at 1,100 F. andair cooled. Brittle fracture occurring in shoulder area. d) Longitudmal MBR, T= 10.

e) Longitudinal MBR, T=2.9/6.5.

YTABLEXX -f Room temperature tensile test" properties ofTi.-Sn-4Cu2Fe and Ti12Sn5Cu2Mo alloys 'ri-iosn-wu-zre Heat Treatment Tensile 1roperfies,

p. 5.1.' X 1,000 Percen Percent Elon- Reduc- VHN gation tion in T 0.2% Ultimate in 1" Area Time, hours S?" Cooling Rate Offset Strength Yield 7 1, 200 Fast air 0001..--- 116 125 7 29 315 1,280 Slow air cool... v 124 136 10 15 312 1, 360 rln A 114 128 4 19 317 1,400 do 129 161 5 12 362 1,460 do 184l 185 1 6 445 T1-12Sn-5Cu-2Mo 1, 40o Fast air cool s1 136 1o 19 28o 1,400 -,do 82 138 10 26 280 1 500 1294 132 1 1 354 60 157 13 27 318 58 165 13 24 285 127 147 3 8 361 y 167 170 1 6 454 143 168 1 6 400 1, 525 Quench 177 l 3 468 1, 525 Fast air cool 172 1 5 475 1, 525 Slow air cooll 186 1 3 470 1, 525 Very slow air cooll... 176 178 1 6 448 Tempered 2 hrs. at 1,100 F. after heat treatment indicated.

The above tables show that the substitution of 2% iron for part of the copper lowers the permissible solution annealing temperature. However, our further tests have shown that substitution of 4% ironfor part of the copper produces an alloy which is not responsive to heat treatment on slow air cooling. Substitution of 2% iron lowers the effective solution-annealing temperature by about 75-100 F., and lowers the eutectoid temperature by about G-250 F. The above data alsorshows that substitution of molybdenum for part of thecopper as a sluggish beta stabilizer to slow down the beta decomposition reaction, is quite effective in this respect. .As' shown, molybdenum also stabilizes the betal phase sufciently so that, in the quenched and fast air cooled conditions, the alloy is mechanically unstable as evidenced by low yield strength and by the high ultimate to yield strength ratio. f

The alloys of the invention may be produced by arc melting in an inert atmosphere, such as argon, in a cold mold furnace, for example, one employing a water-cooled crucible, or by equivalent procedures which substantially prevent atmospheric contamination or contamination by other elements productive of undue embrittlement or impairment of properties.

These alloys are useful in sheet form with minimum bend ductlities as high as 20T, and may be forged or rolled with tensile elongations as low as 1 or 2%.

Even in the fully annealed condition, the majority of these alloys have tensile strengths at room temperature of at least 100,000 p. s. i.; while those that do not may be heat treated as by quenching from the beta or alpha-beta fields to impart thereto a tensile strength of at least 100,000 p. s. i.

This application is a 4continuation-in-part of our copending applications as follows: Serial No. 515,539, filed June 14, 1955 (now abandoned), and Serial No. 421,980, tiled April 8, 1954 (now abandoned), which in turn'is a continuation-impart of our now abandoned case Serial No. 270,896, filed February 9, 1952; yand Serial Nos. 393,578, filed November 23, 1953 (now U. S. Patent 2,796,347, granted June 18, 1957); 396,756, led December 7, 1953 (now U". S. Patent 2,797,996, granted July 2, 1957); 384,326, filed October 5, 1953 (now abanantimony, 0.5 to 20% of at least one active eutectoidv beta promoter selected from the group consisting ofcopper, nickel, cobalt, silicon and beryllium, but not toexceed 12% of nickel andl cobalt, 3% of silicon'and 2% of beryllium, balance substantially titanium, characterized in being materially hardened and strengthened on heating above the critical temperature, until a betacontaining microstructure is obtained and thereuponl rapidly cooling to a temperature substantially below thecritical.

2. An alloy consisting essentially of about: 0.5 to l 23% of at least one alpha promoter selected from the,

group consisting of aluminum, tin, antimony and zirconium, but not to exceed 12% aluminum and 19% antimony, 0.5 to 20% of at least one active eutectoid beta promoter selected from the group consisting of copper,

nickel, cobalt, silicon and beryllium, but not to exceedl 12%of nickel and cobalt, 3% of siliconv and 2% of beryllium, up to 0.5% of elements of the group carbon, oxygen and nitrogen, but not to exceed 0.4% oxygen and 0.25% nitrogen, balance substantially titanium, characterized in having in the annealed condition a tensile strength of at least 100,000 p. s. i. and a minimum tensile elongation ofabout 2% and in being materially hardened and strengthened on heating above the ycritical temperature, until a beta-containing vmicrostructure is obtained and thereupon rapidly cooling to a.v temperature substantially below the critical.

, 3. An alloy consisting essentially of about: 0.5` to 23%l of at least one 'alpha promoter selectedl frornthey group consisting of aluminum, tin, antimony and zir-v conium, but not to exceed 12% aluminum and 19%.anti-I mony, 0.5 to 20% of copper, and the balance substantially titanium, characterized in having in the annealed condition a tensile strength of at least 100,000 p. s. i., a minimum tensile elongation of about 2%, and in beingmaascissa El terlally hardened and strengthened on heating above the critical temperature until a beta-containing microstructure is obtained and thereupon rapidly cooling to, a

temperature substantially below the critical.

4. An alloy consisting essentially of about: 0.5 to 23% of at least one alpha promoter selected from the group conslsting of aluminum, tin, antimony and zirconium, but- 32? characterized by high stability after exposure in both stressed jand: unstressed conditions to elevated temperai tures up to about 1000 F.

9.. An alloy'consisting essentially of about: 0.5 to 12% aluminum, 0.5V to'20% copper, balance substantially titanium, characterized-in being forgeable and in being materially hardenedand strengthened on heating above the not to exceed 12% aluminum and 19% antimony, 0.5 to

12% of at least one active eutectoid betaA promoter selected from the group consisting of nickel and cobalt,

and the balance substantially titanium, characterized in rapidly cooling to a temperature substantially below the A' critical.

5. An alloy consisting essentially of about: 0.5 to 23% of at least one alpha promoter selected from the group consisting of aluminum, tin, antimony and zirconium, but i' not to exceed 12% aluminum and 19% antimony, 0.5i to 3% of the active eutectoid beta promoter silicon, and 1 the balance substantially titanium, characterized in having 1n the annealed condition a tensile strength of at least 100,000 p: s. i. a minimum tensile elongation of about 2%, and in being materially hardened and strengthened' on heating above the critical temperature, untill a betacontaining microstructure is obtained and thereupon rap-l idly cooling to a temperature substantially below the critical.

6. An alloy consisting essentially of about: 0.5 to 23 of atV least one aloha promoter selected from the group consisting of aluminum. tin. antimony and zirconium, but not to exceed 12% aluminum and l19% antimony. 0.1 to 2% of the active eutectoid beta promoter beryllium,r and the balance substantially titanium, characterized in having in the annealed condition a tensile strength of-v atA least 100,000 p. s. i., a minimum tensile elongation of about 2%. and in being materially hardened and strengthened on heating above the critical temperature. until a beta-containing microstructure is obtained. and thereupon rapidly cooling to a temperature substantially below the critical.

7. An alloy consisting essentially of about: 0.5 to 20% of at least one active eutectoid beta promoter selectedj from the group consisting of copper. nickel.` cobalt kand beryllium. but not to exceed 12% nickel and cobalt, 3% silicon and 2% beryllium, said alloy also containing at least one alpha promoter selected from the group consisting of aluminum, tin and antimony in total amountsuch that the sum of aluminum plus one-third the tin plus one-third the antimony is under 6% by weight of tlie total alloy. said alloy also containing about l` to 20% zirconium, balance substantially titanium, characterized in being materially'hardened and strengthened on heating above the critical temperature until a beta-containing microstructure is obtained, and thereupon rapidly cooling to a temperature substantially above the critical, said alloy being also characterized by high stability after exposure in both stressed and unstressed conditions to elevatedtemperatures up to about 1000 F.

8. Analloy consisting essentially of about: 0.5 to 20% of the active eutectoid beta promoter copper, and containing at least one aloha promoter selected from the group consisting of aluminum, tin and antimony in total amount such that the sum of the aluminumplus onetbird the tin plus one-third the antimony is less than 6% by Weight of the total alloy, said alloy also containing about 1 to 20% of the alpha promoter zirconium, balance substantially titanium, characterized in being materially hardened and strengthened on heating above the critical temperature until a beta-containing microstructure is obtained, and thereupon rapidly cooling to a temperature substantially below the critical, said alloy being also critical temperature until a beta-containing microstructureis obtained and thereupon rapidly cooling to a temperature substantially' below the" critical.

10. An alloy consisting essentially of about: 4 to 8% aluminum, 2 to 4% copper, balance substantially titanium, characterized inhaving, as annealed, a rtensile strength of atleast 100,000 p.r s. i., a minimum tensile elongation-of -about 2%, andi in being materially -har-V dened and' strengthened onpheatingf above the critical temperaturev untilV ahem-containing microstructure is Vobtained and thereuponV rapidly cooling to a temperature substantially below' the critical.

11'. An alloy consisting essentially of about: 0:5 to under 6% aluminum, about 11 to 12% zirconium, about 0.5'` to 20% copper, and' the balance substantially titanium, characterized in having, as; annealed, ya tensilek strength off atfleastfl'QOOO p.V s. i., a minimum tensile elongation of about 2%, and inbeing materially hardened and strengthened onrheating, above the critical temperature untilv a beta-containing vmicrostructure is obtained and thereupon rapidly'cooling; to a temperature substantially below the critical.

12. Analloytconsistingessentiallyof about: 0.5 to 23% of at least one alpha promoter selected from the group consisting ,of` aluminum, tin` and antimony, about l to 20% of the alpha promoter zirconium, 0.5 to 20% of at least one active eutectoid beta promoter selected from the group consisting Iof copper, nickel,` cobalt, silicon and beryllium, but not to exceed'l2% of nickel andY cobalt, 3% silicon and 2%A beryllium, balance substantially titanium, characterized in having in the annealed condition a tensile strength of at least 100,000 p. s. i., a minimum tensile elongation of about 2%, and in being materially hardenedv and strengthenedv on heating above the critical temperature until a beta-containing microstructure is obtained and thereupon rapidly cooling to a temperature below the critical.

An alloy consisting essentially of about: 8 to 14% tin, 5 to, 8% copper, balance. substantially titanium, characterized` in being materially hardenedY and` strengthened on heating above the critical temperature until a betacontaining microstructure is obtained and thereupon air cooling to a temperature substantially below the critical.

14. An alloy consisting essentially of about: 0.5 to 23% of at least one element of the group aluminum, tin and antimony, but not to exceed 12% of aluminum and 19% 5'5: antimony, 0.5,to 20% copper, balance titanium.

l5. An alloy consisting of about: 0.5 to 23% of at least one element selected from the group consisting of aluminum, tin and antimony, but not to exceed'12% aluminum and 19% antimony, 0.5 to 12% of at least one element selected from the group consisting of nickel and cobalt, balance titanium.

1'6. An alloy consisting of about: 0.5 to 23% of at least one element selected from the group consisting of aluminum, tin and antimony, but not to exceed 12% aluminum and 19% antimony, 0.25 to 3% silicon and the balance titanium.

17. An alloy consistingv of about: 0.5 to723% of at least one Yelement selected from thegroup consisting of aluminum, tin and antimony, but not to exceed 12%v aluminum and 19% antimony, 0.1 to 2% beryllium, balance titanium.

18. The method of annealing au alloyaccording to.

claim 1 which comprises, heating above the beta transus until transformation to the beta phase is substantially complete, and thereupon slowly cooling at the rate of not more than -10 temperature range.

19. The method of annealing an alloy according to claim 1 which comprises, heating at temperature of about 50-100 F. below the eutectoid temperature to remove strain and promote compound agglomeration and thereupon cooling substantially to room temperature.

20. The method of hardening an alloy according to claim 1 which comprises, heating above the eutectoid temperature until equilibrium of the alloy and transformation to the beta phase is substantially complete, and thereupon cooling rapidly through the eutectoid temperature at a rate suicient to produce a fine eutectoid structure, said cooling being at least as rapid as an air cool.

21. The method of hardening an alloy according to claim 1 which comprises, heating the alloy above the beta transus until transformation to the beta phase is substantially complete, and thereupon cooling through the eutectoid temperature with suicient rapidity to produce martensite, said cooling being at least as rapid as an air cool.

22. The method of hardening an alloy according to claim 1 which comprises, heating the alloy in the alphabeta temperature range until equilibrium is established, and thereupon rapidly cooling through the eutectoid temperature, said cooling being at least as rapid as an air cool.

23. The method of hardening and tempering an alloy according to claim l which comprises, heating above the eutectoidtemperature until equilibrium of the alloy and transformation to the beta phase is substantially complete, thereupon cooling at least as rapidly as an air cool F. per minute through the eutectoidl through the eutectoid temperature and at a rate sufficient to produce a line eutectoid structure, and thereupon aging at temperature below the eutectoid temperature.

24. The method of hardening and tempering an alloy according to claim 1 which comprises, heating the alloy above the beta transus until transformation to the beta phase is substantially complete, thereupon cooling through the eutectoid temperature with suicient rapidity to produce martensite, said cooling being at least as rapid as an air cool, and thereupon aging at temperature below the eutectoid temperature.

25. The method of hardening and tempering an alloy according to claim 1 which comprises, heating the alloy in the alpha-beta temperature range until equilibrium is established, thereupon rapidly cooling through the eutectoid temperature, said cooling being at least as rapid as an air cool, and thereupon aging at temperature below the eutectoid temperature.

26. The method of hardening an alloy according to claim 1 which comprises heating the alloy above the critical or eutectoid temperature and then cooling with sufficient rapidity to cause eutectoid decomposition of the beta formed on heating, said cooling being at least as rapid as an air cool, but slow enough to prevent undue distortion of the heat-treated alloy.

References Cited in the tile of this patent UNITED STATES PATENTS Frazier Dec. 16, 1952 2,669,513 Jaee Feb. 16, 1954

Claims (1)

1. AN ALLOY CONSISTING ESSENTIALLY OF ABOUT: 0.5 TO 23% OF AT LEAST ONE ALPHA PROMOTER SELECTED FROM THE GROUP CONSISTING OF ALUMINUM, TIN, ANTIMONY AND ZIRCONIUM, BUT NOT TO EXCEED 12% ALUMINUM AND 19% ANTIMONY, 0.5 TO 20% OF AT LEAST ONE ACTIVE EUTECTOID BETA PROMOTER SELECTED FROM THE GROUP CONSISTING OF COPPER, NICKEL, COBALT, SILICON AND BERYLLIUM, BUT NOT TO EXCEED 12% OF NICKEL AND COBALT, 3% OF SILICON AND 2% OF BERYLLIUM, BALANCE SUBSTANTIALLY TITANIUM, CHARACTERIZED IN BEING MATERIALLY HARDENED AND STRENGTHENED ON HEATING ABOVE THE CRITICAL TEMPERATURE, UNTIL A BETACONTAINING MICROSTRUCTURE IS OBTAINED AND THEREUPON RAPIDLY COOLING TO A TEMPERATURE SUBSTANTIALLY BELOW THE CRITICAL.

Images