IL33198A

IL33198A - Cast nickel base alloy

Info

Publication number: IL33198A
Application number: IL33198A
Authority: IL
Original assignee: Gen Electric
Priority date: 1968-11-01
Filing date: 1969-10-15
Publication date: 1972-07-26
Also published as: DE1952877C3; FR2022356A1; DE1952877A1; DK124893B; FR2022386A1; IL33198A0; GB1256017A; ES372869A1; BE740895A; SE357983B; DE1952877B2; US3615376A; CH533683A

Description

CAST NICKEL BASE ALLOY . 13D-44S2 Advancing technology arid development of improved power producing appafatus such as the gas turbine engine has identified the need for stronger alloys which are stable at relatively high operating temperatures such as up to ISOOST and yet can withstand the corrosive atmospheres in which they are intended to operate. Although a number of alloy systems including those based on the refractory metals have been evaluated for such applications, the nickel base alloy remains the type presently most widely used in such difficult applications.

One high temperature nickel base alloy application of particular intere st is the cast form of the alloy. However, known nickel base alloys in cast forms either are relatively weak or unstable during long time operation or have insufficient resistance in hot corrosive atmospheres particularly in the 1500°F - 1800°F. range.

Briefly stated , it has now been recognized that a cast nickel base alloy having an improved combination of high temperature stability and hot corrosion resistance along with a stress rupture life of at least 25 hour s under stress of 27, 500 psi at 1800°F can.be attained through (1) the control of the type of precipitation of strengthening phases first with carbon, and second with the elements titanium and aluminum in a nickel matrix, (2) the control of the solution strengthening mechanisms as a result of the presence of W and Mo in particular portions to 13D -4452 such as sigma and eta. Broadly, the composition which defines such an alloy consists essentially of, by weight, 0.1 - 0.3% C; greater than 13% but less than 15.6% Cr.; 4 - 6% Ti; 2 - 4% Al; - 0.005 - 0.02% B; 3 - 6% W; 2.5 - 5% Mo; greater than 5% up to 15% Co; up to 0.1% Zr; with the balance nickel and incidental impurities provided that the ratio of Ti to Al is greater than 1 but less than 3, the sum of Ti and Al is 7.5. - 9%, and the sum of all Mo and half the W is 5 - 7%.

In the alloy of the present invention, carbon, preferably in the range of 0.15., - 0.2%, provides for carbide formation which leads to improved strength particularly at high temperatures.

Insufficient carbon, is insufficient for high temperature strength whereas an overabundance of carbon results in lower life and embrittlement at lower temperatures as a result of excessive carbide formation in the grain boundaries.

The element chromium provides oxidation and hot corrosion resistance. However, in amounts of less than 13% there • is insufficient hot corrosion resistance provided in the temperature range of about 1500 - 180.0°F. Cr in amounts greater than 16% leads to the formation of sigma and other deleterious phases without proper phase control. Accordingly, the preferred Cr range is 13.5 - 14.5% to assure such phase control. 13D-4452 As is the case with Cr, Co in excessive amounts can result in sigma phase formation. However, in the proper amounts described herein, Co adds to the gamma prime solubility and affects ductility of the alloy.

Very critical to the alloy of the present invention are the elements W and Mo which generally are identified with the solution strengthening mechanism of a nickel base alloy. However, it has been recognized that a complex control of both sigma phase and of precipitating carbides can be achieved through a careful balance of the amounts of W and Mo. As will be shown in detail in connection with the specific examples, it was unexpectedly recognized that if the total amount of W and Mo were maintained such that the sum of half of the W and all of the Mo were in the range of 5 - 7%, not only could the formation of sigma phase be inhibited but also the more stable MgC carbide could be formed along with the carbide rather than all MggCg . AltnouSn Mo has been included in substantial amounts in certain known nickel base alloys, it has been recognized that Mo on a weight percentage basis is a more potent signia phase former. It has also been previously shown that W additions are beneficial to 1500 - 1800°F stress rupture strength.

Therefore, it is necessary that both Mo and W be present and that the ranges in which the above relationship is maintained is 3 -6% W and 2. 5 - 5% Mo in order to maintain optimum alloy strength and stability.

The elements Ti and Al have been described in connection with their function as the primary precipitation strengthening elements in combination with nickel in forming Nig(Al Ti) . Now it has been <«►·* 13D-4452 ο the 1500 - 1800 F range. This unique combination of Al and Ti along with that just described for Mo and W in their dual control function is one important aspect of the present invention not recognized in known nickel base alloys.

The present invention recognizes that Ti/Al ratio must be greater than 1. to provide such improved hot corrosion resistance but less than 3 to prevent the formation of the weakening eta (Ni^Ti) phase. Al is preferably included in the range of about 2.5 - 3.57c because for one reason it can result in the formation of sigma phase somewhat more readily than does Ti; Al ties up more nickel from the matrix to form the Ni^(Al, Ti), sometimes referred to as gamma prime. This occurs because of the lower atomic weight of aluminum compared with titanium. As the gamma prime content increases, there is less nickel available in the gamma matrix. Therefore, there is a greater tendency for sigma phase formation due to the relatively larger amounts of Cr, Co, Mo and W in the matrix. Accordingly, it is an objective to keep as much nickel as possible in the gamma matrix. .

Hence, keeping close control and lowering the Al content relative to the Ti content will result in less tendency to form the embrittling sigma phase and the higher Ti/Al ratio will improve hot corrosion resistance.

The present invention recognizes the criticality of this balance of aluminum and titanium not only from the standpoint of 13D-4452 no more than 9 weight percent can be tolerated without seriously depleting the nickel matrix. The proper amount of Al stabilizes 13D-4452 formation. With too much Ti, the Ni (Al, Ti) is metastable and breaks down to form the weakening Ni Ti.

Although iron has been included or tolerated in certain relatively large amounts in known nickel base alloys, the present invention recognizes that iron tends to form deleterious phases . Therefore, it is preferred that no iron be present although slight adjustment such as in the solid solution strengthening elements can be made to tolerate small amounts of Fe .

Boron is included within the range of 0. 005 - 0. 02% for its beneficial effect on rupture, strength and ductility. The alloy including boron below that level is weak, whereas too high a boron content results in the formation of excessive borides leading to incipient melting on over temperature exposure.

It has been recognized in evaluation of the present invention that the elements Cb and Ta are not substitutes for W and Mo.

It is believed that about half of the Cb or Ta goes into gamma prime formation, such as Ni (Al, Ti, Cb, Ta), and to carbides . Both deplete 3 the matrix and are undesirable in the balanced alloy defined by the present invention. Both can lead to the formation of sigma phase .

These unusual aspects of the present invention will be more clearly understood from the following detailed examples typical of alloys melted in the evaluation of the alloy of the present invention. The alloys were melted by the commercial vacuum melting techniques widely used in the preparation of nickel base alloys . Heats ranging in size from about 12 pounds to about 1000 pounds have been made, the latter being I3D-4452 into precision casting specimen molds or by remeltxng and casting previously prepared alloy ingots.

The alloy forms representative of those melted within the scope of the present invention are shown in the following Table I.

TABLE I Composition in ' weight percent Alloy includes 0.014 - 0.016% B; 0.03% Zr, balance Ni and incidental impurities. Ti/Al = 8 - 8.1; Ti/Al = 1.6 - 1.7; Mo/W/2 = 5.4 - 6 Alloy C Cr Co Mo W Ti Al 1 .17 14.0 9.9 4.0 3.9 5.0 3.0 2 .16 14.0 9.6 4.0 4.0 5.0 3.0 3 .19 14.2 14.9 4.0 4.0 5.0 3.1 4 .18 14.1 7.5 4.0 4.0 5.0 3.0 .18 14.0 12.3 4,0 4.0 5.0 3.0 6 .19 13.9 10.0 3.0 6.0 5.0 3.0 7 .18 13.5 10.0 4.0 4.0 5.0 3.0 8 .26 14.0 9.8 4.0 4.0 5.0 3.1 9 .19 14.0 10.0 2.9 5.0 "4.9 3.1 .15 14.1 9.4 4.1 4.0 5.0 3.1 Other alloys made and tested during the evaluation of the alloy of the present invention include those shown in the following Table II, outside the scope of the present invention.

TABLE II Composition in weight percent ■ Alloy includes 0.014 - 0.017% B; 0.03 - 0.04% Zr with the balance Ni and incidental impurities Alloy _C_ Cr Co Mo W Ti Al Ti/Al Ti/Al Mo/W/2 11 .14 16.1 10.7 3.2 2.9 3.8 4.0 1.0 7.8 4.7 12 .14 15.5 10.7 3.1 2.7 2.9 5.0 0.6 7.9 4.5 13 .16 15.4 .10,3 4.0 3.8 4.5 2.6 1.7 7.1 5.9 14 .16· 15.6 10.2 4.0 3.9 4.9 2.9 1.7 7.8 6.0 .20 13.0 10.0 4.0 4.0 4.9 3.0 1.6 7.9 6.0 16 .08 14.1 9.9 4.0 4.0 5.0 3.1 1.6 .8.1 ' 6.0 17 .19 14.1 5.0 4.0 4.0 4.9 3.1 1.6 8.0 6.0 18 .18 14.0 0 3.9 4.0 4.9 3.1 1.6 8.0 5.9 13D-4452 The improved characteristics of the present invention are particularly measured by combination of high temperature stress rupture life and stability along with hot corrosion resistance. This improvement as it relates to long time stability is related to the suppression of the formation of such embrittling phases as sigma and eta. These phases are greatly suppressed or are entirely eliminated according to the alloy of the present invention.

When certain known cast alloys are exposed to elevated temperatures, the gamma phase and carbides, which are found in the primary gamma prime phase, agglomerate. At temperatures in the range of about 1300°F - 1800¾\ sigma plates form in matrix areas surrounding the gamma prime. This formation, which is accelerated by stress, appears to relate to excessive chromium in the primary gamma prime and surrounding matrix areas, first reacting with carbon to form grain boundary M C carbides. Then 23 6 when all the available carbon is thus reacted, it appears that excessive chromium in the matrix combines with such elements as Co, Mo, etc., to form a Cr-Co-Mo type sigma. Long time stability testing such as at 1500°F at a stress of 55, 000 psi identifies the strength-reducing nature of sigma phase.

Although sigma phase may be removed by heat treat¬ ment, it will recur when the alley experiences the same time and temperature conditions under which sigma was originally formed.

The alloy of the present invention identifies a different kind of 13D -4452 stability along wit hot corrosion resistance as a result of a different surface reaction product.

In order to understand more fully the present invention and its individual components as they affect the strength and stability of the alloy of the present invention, the following tables have been prepared. These compare the alloy forms both within and outside the scope of the present invention, as shown in full compositions in Tables I and II. The element, content referred to in the tables as well as' throughout this specification is in weight percent and the term "ksi" refers to thousands of pounds per square inch.

TABLE. Ill Element Variation As -Cast Stress Rupture Life(hr Alloy Wt. % 1500 F/55 ksi 1800 F/27.5 ksi Sigma :i8 0 Co *230 too weak none 17 5 Co *481 to test none 6 10.0 Co 744 49 none 12.3 Co 896 50 none 14 15..6 Cr. 586 38 large 6 13.9. Cr 744 49 none 13.0 Cr 588 30 none •16 0.08 C 512 45 small 0.18 C 896 50 none 8 0.26 C 670 43 none *Heat Treated: 2200°F - 2 hrs. ; 2000°F - 4 hrs. , 1550°F - 16 hrs. 1400°F - 16 hrs.

As shown in Table III, in the alloy of the present invention, cobalt below about 15 weight percent does not lead to the formation of sigma phase. However, above 15 weight percent, excessive si ma will form resultin in a different kind of allo of 13D -4452 With respect to the chromium variation shown in Table III, the detrimental effect of the formation of heavy amounts of sigma on long time stability is shown by alloy 14 at 15. 6% Cr. The identification of large amounts of sigma shows alloy 14 to be of a different kind than that of alloy 6 within the scope of the present invention. Alloy 15 at 1 % Cr and only 0. 9% lower than alloy 6, shows a reduction in strength even though all other elements of alloy 15 are within the range of the present invention. Therefore, the alloy of the present invention includes less than 15. 6% but greater than 13% Cr.

With respect to the carbon variation in Table III, it can be noted that at 0. 08% C, insufficient carbon is present to react with Cr within the range of the present invention to prevent Cr from forming sigma platelets. The reduction in long time stability as represented by the 1500°F tests should be noted in this regard. Although amounts of carbon approaching about 0. 3% can be included, it is preferred that carbon at about 0. 2 weight percent be maintained in order to assure the unusually fine properties of the preferred form of the alloy of the present invention.

Although the elements W and Mo have been included in known nickel base alloys singly or interchangeably as solution strengthening elements, the present invention recognizes additional critical roles played by these two elements. Both are involved in the complex control of precipitating carbides and sigma phase formations, although Mo is a more potent sigma phase former. The following Table IV shows the effect and interrelationship of these elements on the alloy of the present invention. 13D-4452 TABLE IV Element Variation As -Cast (weight %) Stress Rupture Life (hrs . ) Alloy Mo _W (Mo/W/ 2) 1500°F/55ksi 1800°F/27. 5ksi . Sigma 19 1 3. 0 7. 6 * 21 * medium 6 0 6. 0 6. 0 744 49 none 9 4. 9 7. 4 439 47 medium 9 9 5. 0 5. 4 746 37 none * based on a 200 hour life 1500°F - 68ksi rupture test In alloy 19, even with Mo as high as 6. 1%, there is insufficient strengthening to provide adequate high temperature stress rupture strength. More importantly, however, is the fact that the total amount of Mo and W is sufficiently high to result in sigma phase formation as measured by the atomic relationship between those elements of (Mo / W/ 2) of as high as the 7. 6% level. The present invention contemplates that relationship to be within the range of 5 - 7% to inhibit sigma phase formation and precipitation of the proper carbides as described before. iUloy 20, a different kind of alloy and outside the scope of the present invention, includes Mo and W within the invention range but with the improper relationship one to the other as shown by the (Mo / W / 2) cf 7. 4%. The formation of medium amounts of sigma resulted in significantly reduced stability as measured by the 1500°F stress rupture test. Alloy forms 6 and 9, within the scope of the present invention, have the proper balance of W and Mo and are a different kind of alloy because of the absence of the sigma structure. This results in improved stability and strength.

In the above Tables III and IV, it should be noted that the alloy forms identified with numbers greater than 10 have compositions 13D-4452 case of alloy 20 is (Mo W / 2) .

The elements Ti and Al contribute to the alloy of the present invention in several ways. This invention recognizes that the proper amount and interrelationship between these elements can control the short time strength, the alloy stability through sigma phase inhibition and, very importantly, provide hot corrosion resistance.

The problem of hot corrosion resistance involves resistance to sulfidation in the range of about 1500 - 1800°F . Above and below, that range, hot corrosion resistance is not as significant a problem in the type of alloys to which the present invention relates because such alloys include the element aluminum. Aluminum oxide which forms on the surface as a reaction product will form a reasonably protective oxidation resistant barrier. The problem of oxidation resistance is different from that of hot corrosion resistance. Normally for oxidation resistance it would be better to have a Ti/Al ratio of less than 1. However, this is contrary to the relationship defined for the alloy of the present invention which requres the Ti /Al ratio of greater than 1. The higher ratio is desirable because TiOg is formed on the surface. The more TiO available, the better is the hot corrosion 2 resistance. However, Ti in amounts which would produce a Ti/Al ratio of about 3: 1 or more, cannot be tolerated in the alloy of the present invention.

The effect of the elements Al and Ti on the alloy of the present invention as it relates to as-cast stress rupture life and stability is shown in the following Table V. 13D-4452 TABLE V As -Cast Wt. % Stress Rupture Life (hrs. ) Alloy Al Ti (Ti/Al) Ti/Al 1500°F/55ksi 1800°F/ 27. 5 ksi Sigma 11 3. 8 4. 0 7. 8 1. 0 . 309 35 medium 12 2. 9 5. 0 7. 9 0. 6 251 25 large 13 4. 5 2. 6 7. 1 . 1. 7 560 19 none 5. 0 3. 0 8. 0 1. 7 896 50 none Although alloys 5, 11 and 12 include about the same amount of the sum of titanium and aluminum, it should be noted that alloy 5 forms no sigma phase whereas alloys 11 and 12 form medium to large amounts of sigma phase. This can be attributed to the improper relationship between the two elements. The fact that different kinds of alloys are formed between alloy 5 and alloys 11 and 12 is further substantiated by the stress rupture life, particularly the stability data represented by the 1500°F tests. Further, it should be noted that alloy 13, although having the proper ratio between Ti and Al, does not have sufficient amounts of these elements to provide the required strength. Therefore, the alloy of the present invention defines the relationship between Ti and Al such that the sum of those elements is in the range of about 7. 5 - 9% and that the Ti /Al ratio is greater than 1 but less than 3: 1.

One important characteristic of the alloy of the present invention which distinguishes it from known alloys presently intended for the same use is its significantly improved hot corrosion resistance. A series of comparison tests to determine the hot corrosion resistance of a variety of alloys was conducted on such known nickel base super alloys as those listed in the following Table VI. 13D-4452 TABLE VI Known Alloys in Wt. % .

Alloys include .01 - .02 B, Balance Ni and incidental impurities Known Alloy C Cr Co Mo W Ti Al Zr Others A ..18 9.5 15.0 3.0 - 4.2 : .5.5 .06 l.V B .07 14.2. 15.0 4.2 - 3.4 4.3 - C 1.3 6.0 7.5 2.0 6.0 1.0 5.5 1..3 9Ta,0.5 Hf 0.5Cb, 0.5Re D .08 15.0 22.0 4.4 - 2.4 4.4 E .14 13.0 - 4.5 - 0.75 6.0 .10 2.3Cb + Ta Because the alloys tested were intended for use in a gas turbine engine, test apparatus simulating conditions in the turbine section of a gas turbine was constructed. The apparatus burned jet fuel, for example JP-5 in a 30 - 1 air-fuel mixture and injected sea water having a composition within the range of ASTM specification D-665-60. The sea water was diluted with distilled water to 5 parts per million. The tests run were cyclic tests over a period of 1, 000 hours including 18 intermittent cooling cycles to room temperature with an air blast. The specimens tested were cast bars ground to a diameter of about 0.130" and about 1.25" long. Results of such a comparative test are shown in the following Table VII, A. 131) -4452 From Table VII, it is easily seen that at all temperatures tested , alloy 2 within the scope of the present invention is remarkedly more resistant to hot corrosion than are all of the other tested known alloys, most of which are presently in use in the hot section of gas turbine engines .

Another measure of hot corrosion resistance involved a study of specimen weight loss rather than surface penetration or thickness loss . Another series of tests resulting in data of which the data of Table VIII is typical were performed on alloys bdth within and outside the scope of the present invention.

TABLE VIII Hot Corrosion Resistance 1750°F for 500 hrs .

Wt. Loss (mils /dia. ) Alloy Gross Max. 0. 3 . 3. 1 1 5. 16. 12 6. 14.

D 3. 10.

E 50. . 53.

Alloy 10 within the scope of the present invention shows significant and remarkable resistance to weight loss after 500 hours at 1700°F as compared with alloys known or outside the scope of the present invention.

. The significantly improved hot corrosion resistance of the alloy of the present invention is based on the fact that it is a different kind of alloy. Hence a different kind of reaction product is formed on the surface of the alloy of the present invention under oxidizing conditions than is formed on the surfaces of certain known nickel base alloys 13D-4452 400 hours at elevated temperatures . The results of one such comparison is shown in the following. Table IX.

TABLE IX X-RAY DIFFRACTION DATA after 400 hrs . exposure Alloy 1700°F ' ' ; 1800°F 2 TiO (S) / Cr90 (M) Matrix (W) TiO (M) /Cr?0„(M) /Spinel (M) B Matrix Matrix (S) /Al2C>3(M) /Ti02(VW) (S) strong, (M) medium, (W) weak, (V) Very Alloy 2 within the scope of the present invention and having a remarkable resistance to hot corrosion had a substantial amount of TiO in its surface, reaction product. Only a small amount of that oxide 2 is found in the reaction product of the alloy B. Thus the two alloys are of a different kind.

Claims

1. ΑρρΙη,Νο. 33198/2 1, A cast nickel base alloy of improved stability, strength and hot corrosion resistance, consisting of, by weight: 0,15 - 0.3 C, said carbon percentage being greater than that required for deoxidation; greater than 13$ but less than 15.6$ Cr; greater than 5 up to 15$ Co; 2,5 - $ Moj 3 r 6$ W; 4 - 6$ Ti; 2 - 4$ Alj 0.p05 - 0.02$ B; ■■■v up to 0.1$ Zr; with the balance nickel and incidental impurities; the ratio Ti/Al being in the range of greater than 1 to less than 3; the sum of Ti and Al being in the range of 7.5 - 9$; and the sum of Mo and half of the V being in the range of 5 - 7$; and further characterized by the substantial absence of sigma phase and a stress rupture life in the as-cast condition of at least 25 under a stress of 27,500 psi at 1800°F, 2, The alloy of claim 1 in which: the C is 0.15 -0.2$; the Cr is 13,5 ^ 14,5$; the Co is 7.5 - 12.5$ the Mo is 3.5 - 4.5$; the W is 3,5 - 4.5$; the Ti is 4i5 - 5.5$; the Al is 2,5 - 3.5$ the B is 0.01 - 0.02$; the Zr is 0.005 -0.1$; and the Ti/Al is 1 - 2$, 3, The alloy of claim 2 in w'kch the Cr is 13,7 -14.3$; the Co is 9 -. 10$; the Mo is 3.7 - 4.3$; the V is 3,7 - 4,3$; the Ti is 4.8 - 5.2$; the Al ie 2.8 - 3.2$; and the Zr is 0,02 - 0.04$. Tel-Aviv, October 14, 1969