THE TEMPERATURE DEPENDENCE of SUSTAINED LOAD CRACKING in Ti-6211 C.L

2557

THE TEMPERATURE DEPENDENCE OF SUSTAINED LOAD CRACKING IN Ti-6211 C.L. Hoffmann, Combustion Engineering, Windsor, CT 06095 USA J.E. Cox, R.W. Judy, Jr., and B.B. Rath, Washington, DC 20375 USA

In recent years, the design of structures and components has appropriately e1nphasized fail-safe principles. The basis for these principles includes the recognition that cracks and voids that either are initially in the components or are formed early during the service 1 i fe can continue to grow during service. The durabi 1ity or useful 1ife. of a component is directly related to the rate of degradation of its load-bearing capacity by subcritical crack growth and is, therefore, a function of the material's resistance to subcritical crack growth [1-3]. Sustained load cracking (SLC) due to internal hydrogen in titanium alloys is very complex and depends on alloy composition, microstructure, hydrogen content, and applied stress intensity. One area of interest in investigations of SLC in the near-alpha alloy Ti-6Al-2Cb-1Ta-0.8Mo (Ti-6211) is the temperature dependence of crack growth; this aspect can provide information concerning t herma 1 ly activated rate control 1i ng processes which may determine SLC crack growth rates [4].

In work on the titanium alloy Ti-6Al-6V-2S·n, Moody and- Gerberich [5,6] calculated an activation energy for SLC crack growth in air of approximately 30 kJ/mol. The value was related to the diffusion of nydrogen through the beta phase. Threshold stress intensities for SLC crack growth decreased slightly as temperature was decreased below ambient temperature. A significant increase in threshold values was observed at temperatures greater than ambient temperature. The behavior at 1ow te1nµeratures was explained by the temperature dependence of beta phase hydrogen· diffusivity. Difficulty in nucleating hydrides at higher tenperatures was suggested as an explanation for the observed increase in threshold stress intensity. Another study of Ti-6Al-6V-2Sn [7] obtained an activation energy of 14.4 kJ/mol and attributed the temperature dependence of SLC to stress induced hydride precipitation. The proposed model consisted of stress-enhanced transport of hydrogen and nucleation of a cloud of hydrides at the crack tip. Cracking proceeded by a rupture process that linked up cleaved regions resulting fr~n the presence of hydride precipitates. In addition, activation energies have ·been calculated for stress corrosion and corrosion fatigue in T(-6Al-4V [8,9] and cracking of Ti-5Al-2.5Sn and Ti-5Al-4V in gaseous hydrogen [10,11]. All of these studies agree that the cracking process is related to the transport of hydrogen and hydride formation in titanium alloys [8-13].

In the present study, the tenperature dependence of Ti-6Al-2Cb-1Ta-0.8Mo (Ti-5211) alloy as a functio~ of hydrogen content is examined. SLC crack growth rates and threshold s~ress intensities were determined in an effort to define the rate controlling parameter for SLC in this alloy.

Material and Test Methods Beta processed, 25 mm (one inch) thick plate material of Ti-6211 was used 2558 for this segment of the experimental work. A typical mi crostructure for this material is shown in Fig. 1. Crack growth rates as a function of applied stress intensity (K) were determined using side-grooved WOL specimens. Standard pin and clevis loadiny in static load frames was used for determinations of crack growtl1 rates at ambient temperatures [14].

I 40 Plllj

Fig. 1: Microstructure of beta processed Ti-6211. To facilitate determinations of crack growth rates at temperatures other than ambient, a bolt-loaded WOL specimen was used. A constant displacement, decreasing load (and hence, a decreasing stress intensity) SLC test is performed by this method; in addition, the loading bolt is instrumented, allowing the loads to be monitored and the crack growth rates to be computed. A threshold stress intensity for SLC is also obtained since the stress intensity and the crack growth rate both decrease with increased crack length. When the crack growth rate approaches arrest conditions, the stress intensity approaches the threshold for SLC. Good agreement between crack growth rate data obtained by pin-loaded and bolt-loaded methods was demonstrated for the Ti-6211 material. SLC crack growth tests were perfonned for a temperature range of -7S°C to +120°C. Hydrogen was charged into the WOL specimens. using a modified Sieverts apparatus. The charging temperature was 760°C (1400°F), which was not high enough to cause modifications to the microstructure.

Results and Discussion SLC crack growth rates as a function of temperature are compared in Figure 2 for a hydrogen content of 1000 ppm and in Figure 3 for a .hydrogen content of 2SO ppm. Increasing hydrogen content caused increased crack growth rates and decreased threshold stress intensity values for SLC. Similar effects of hydrogen have been observed for SLC crack growth and fatigue crack growth in other titanium alloys [4,7,12,16]. The 1000 ppm material exhibits the highest SLC crack growth rates and lowest threshold stress intensity near ambient temperature (2S°C) (Figure 2). Crack growth rate decreases and threshold stress intensity level increases when the temperature is lowered to -1S°C. Further decrease in the temperature to -7S°C raises the SLC threshold still higher. There is a crossover in crack growth rates at -7S°C, which is probably related to a decrease in fracture toughness at this temperature. The same trends are observed at temperatures higher than ambient. Raising the temperature .to 70°C increases the SLC threshold and decreases the growth rate;: At a tanperature of 120°C, resistance to sustained load cracking increases considerably. 2559

K (1111 ..ll'Fi:I K (ksi ,/Iii.) 20 30 40 1ci' 20 30 40 "° 60 70 Tl-6211 10° Ti-6211 1rf' 1000 ppm H 1 1 10- j 10- ~

10-2 .§'. 10-2.::. 10-1 120-C 15 20'C 10-3 ~ 10-3 g 10-2 10-4 .~~:"_~. 10-4 10-3 10-• ~ .. 10 30 40 50 20 30 40 50 60 JD 80 K (MPo,/iii) K (MPo ./ml Fig. 2: SLC crack growth rate Fig. 3: SLC crack growth rate vs stress intensity as a func- vs stress intensity as a function of tanperature with 1000 tion of temperature with 250 PJlll hydrogen. PJlll hydrogen. The same trend for SLC threshold behavior is observed for material containin\:) 250 ppm hydrogen (Fig. 3). The threshold reaches a minimum at 25°C. No detectable SLC crack growth occurred at temperatures above room temperature at the 250 ppm hydrogen level. Specimens were loaded to high stress intensities at 70°C for long times with no apparent crack growth takiny place. The crack growth rates for 250 PJlll hydrogen material do not exhibit any clear pattern because of the crossover in the data. At low stress intensities, maximum SLC growth rates are observed at 25°C. At higher stress intensity levels, maximum growth rates occur at -15°C. This behavior may be related to greater driving force for hydride formation at -15°C than at +25°C. At -75°C, the rate of diffusion of hydrogen may be slowed enough to reduce the rate of crack growth again.

Curves of the type shown in Figure 4 are obtained when standard Arrhenius type plots are constructed for the SLC crack growth data. The SLC crack growth rates versus reciprocal absolute temperature are plotted at four stress intensity levels for 1000 PJlll hydrogen material (left) and at five stress intensity levels for 250 piJ11 hydrogen lnaterial (right). The curves show very strong dependence of SLC crack growth rates on stress intensity. Crack ~rowth rates should be relatively independent of temperature in order to determine a thermal activation energy for the.process. A weaker stress intensity dependence of activation energy was observed during stress corrosion cracking of Ti-6Al-4V [8]. The decrease in activation ener~ with increased stress' intensity was attributed to lattice dilation ahead of the· crack tip alrowing for faster diffusion rates. Increased lattice di lat ion also lowers the resistance to hydride nucleation, since the volume expansion of the hydride phase will be more easily accommodated [17-21]. The strong stress ·intensity dependence of SLC crack growth rates in Ti-6211 as a function of. temperature precludes the determination of a unique thermal activation erieryy for the SLC crack growth process.

The dependence of SLC thresholds on temperature for hydrogen levels of 250 p'pm and 1000 ppm is shown in: Fig. 5. A minimum threshold stress intensity is observed near ambient temperature at both hydrogen levels. At higher temperatures the threshold .increases rapidly to the point where no SLC crack \:)rowth will occur, as in the case of the material containing 250 PiJTI hydrogen tested at 70°C. This observation is consistent with SLC results 2560

100------~ 1rfl SLC IN Ti-6211 AIR SLCN~H 250 ppm H ~~

K=26MPovm 164 '--~---'---~---'--~-.l-----' 2.0 30 4.0 1 1/T (1000x"K- l Fig. 4: Arrhenius plot of SLC crack growth rate vs reciprocal temperature as a function of stress intensity for beta processed Ti -6211 with 1000 PJJTI hydrogen (left) and 250 pJJTI hydrogen (right). on other titanium alloys; other investigators have concluded that difficulty in nucleation of hydrides above ambient temperature is responsible for increased thresholds or for the need for higher stress intensities to produce a given crack growth rate at higher temperatures [5,7]. At temperatures below ambient, the SLC threshold increases for both hydrogen levels. However, as the temperature is decreased further, the threshold appears to become independent of temperature. The temperature independence of the SLC threshold at low temperatures is likely a result of sluggish hydrogen diffusion rates and the existence of titanium hydrides formed on cool i ng.

The tenperature dependent SLC behavior of Ti-6211 in the beta processed condition indicates that a hydride precipitation mechanism controls the SLC threshold. Formation, growth, and fracture of hydrides in the crack tip region determines the SLC threshold and crack growth rates. The thrt:!shold stress intensity is in large part determined from considerations of initial hydrogen concentration and hydrogen diffusivity and solubility under stress [5-7,11,16,22]. These factors explain the observed tenperature dependence of the threshold for SLC shown in Fig. 5. The driving force for hydride nucleation is reduced as the temperature is incrt:!ased because the hydrogen solubility is very sensitive to temperature. At low temperatures, although the driving force is large, the growth kinetics are slow because the diffusion rate decreases with temperature [21]. In addition, the effect of increased strength at low tenperatures may be to inhibit hydride precipitation [11]. These two opposing trends result in a minimum threshold near ambient temperature whert:! the driving force for nucleation and the diffusion rate for growth result in the 1naximum rate of formation of critical sized hydrides or a critical distribution of hydrides, which produce fracture. The apparent ~lateau in threshold behavior at low temperatures may be attributed to slow hydrogen diffusion rates and the existence of hydrides formed on cooling that are sufficient to cause SLC crack growth. 2561

50 SLC IN Ti-6211 AIR

10 -1~00~--=-s~o~~~o___._____,s~o,--'--rl10=0~~,~so TEMPERATURE , °C Fig. 5: variation of threshold stress intensity (KThl for SLC for beta processed Ti-6211. Aluminum additions to titanium cause an apparent increase in hydrogen solubility. The apparent increase in solubility is actually formation of a supersaturated solution of hydrogen in the lattice, which can result in precipitation of hydrides when stresses are applied. The supersaturated solution forms because hydride formation and its associated volume expansion is inhibited by the mechanical constraint of the stronger matrix of the titanium-aluminum alloy [17,19-21]. Studies on hydrogen solubility of titanium alloys have shown that the equilibrium hydrogen solubility in Ti -6211 is simi 1ar to the sol ubi 1 i ty in the Ti -6A l bi nary [23]. Currently, no information is available regarding hydride precipitation as a function of hydrogen content and temperature for alloys such as Ti-6211. Further evidence in support of the hydride precipitation mechanism is shown in Figure 6. The solubility of hydrogen in a Ti-6Al binary is plotted as a function of reciprocal temperature [17]. Also shown on Figure 6 is the extrapolated value for hydrogen diffusivity in beta titanium [24]. Reasonably good agreement has been obtained between diffusivities measured at low temperatures and extrapolated values of diffusivities obtained at high temperature· in titanium alloys [25]. The direct influence of temperature on the rates of hydride formation and growth is shown by the hydrogen solubility and diffusivity curves in Figure 6 [11].

Data points which represent conditions of SLC test temperature and hydrogen content are. plotted against the hydrogen solubility and diffusivity curves. A clear trend is represented by the data points; they indicate that no SLC crack growth will occur when the conditions of test temperature and hydrogen content are more than 100 ppm below the sofobi 1ity curve. At a temperature of -76°C and hydrogen content .of 40 ppm no SLC crack growth was observed although the test conditions are located slightly above the solubility curve {Figure 6). In this case, hydrides may form on cooling as indicated by the solubility curve. 2562

Assuming hydrides are fonned, the low initial hydrogen content and slow rate of diffusion of hydrogen at -76°C could prevent the hydrides from being larye enough in size and/or distribution to cause fracture •

., ' x x x ' 'x, ' 10LL..L='+..L...L...L...L...L...L.-'--':'"=!--'-'-''-'-''-'-'.....,...='-""'"'~~..L...L."'=":,...... ' 20 30 4.0 1 50 1000/T (°K- )

Fig. 6: Arrhenius plot of solubility limit vs reciprocal temperature for binary alpha titanium alloy Ti-6Al [after Paton et al, ref. 17]. Combinations of hydrogen content and temperature are shown for SLC crack growth in Ti-6211. Extrapolated value of hydrogen diffusivity in beta titanium after Wasilewski and Kehl (24]. It should be noted that the hydrogen solubility curve shown in Figure 6 will be affected by the presence of beta phase in the Ti-6211 alloy and by the stress intensity at the crack tip. For a given bulk hydrogen content the two-phase alpha/beta alloy Ti-6211 will have a lower hydrogen content in the alpha phase than the Ti-6Al binary due to partitioning of hydrogen to the beta phase, which has a higher hydrogen solubility (20,26]. The effect of hydrogen partitioning is to shift the points for SLC test conditions to lower values than are shown in Figure 6. However, the applied stress intensity will tend to lower the hydrogen solubility curve due to fonnation of strain induced hydrides and diffusion of hydrogen out of the beta phase into the alpha phase. The changes in solubility caused by presence of beta phase in the alloy and the applied. stress intensity tend to cancel out. Therefore, the comparison of SLC test conditions with hydrogen solubility in Ti-6Al provides a good first approximation of the SLC susceptibility of Ti-6211 until additional infonnation is obtained concerning hydrogen solubility and hydride precipitation in two phase titanium alloys such as Ti-6211.

Summary The infonnation obtained in the present study on the temperature 2563 dependence of SLC supports the contention that fonnation and growth of hydrides is the mechanism responsible for crack growth in alloys such as Ti-6211. SLC threshold and crack growth behavior were qualitatively described in tenns of hydrogen solubility and hydride precipitation for a beta processed microstructure of Ti-6211. The minimum threshold for SLC observed at ambient temperature results from the temperature and stress dependence of hydrogen solubility and diffusivity. However, SLC crack yrowth in Ti-6211 cannot be attributed to a unique thermal activation energy that is representative of one single kinetic process involved. SLC crack yrowth behavior is detennined by complex interactions of the effects of initial hydrogen content, hydrogen diffusion, hydrogen solubility, hydride nucleation, applied stress intensity, and microstructure.

Acknowledgment The authors wish to acknowledge the able assistance of Mr. W.E. King, Jr. in performing much of the experimental work involved in this.program.

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(25) A.H. Dexter, R.G. Derrick and M.R. Louthan, Jr.: Corrosion 27 (1971) 466. (26) N.A. Tiner, T.L. MacKay, S.K. Asunmaa and R.G. Ingersoll: Trans ASM 61 (1968) 195.