Stress corrosion crack propagation in Ti-13V-11Cr-3Al alloy in methanolic solutions

Corrosion Science, Vol. 19, pp. 799 to 817 Pergamon Press Ltd. 1979. Printed in Great Britain.

STRESS

CORROSION

Ti-13V-11Cr-3A1

ALLOY

CRACK

PROPAGATION

IN METHANOLIC

IN

SOLUTIONS*

R. W. LYCETTand J. C. SCULLY Department of Metallurgy, University of Leeds, Leeds LS2 9JT, England Abstract--Stress corrosion crack propagation in Ti-13V-1 I Cr-3AI alloy in a CHaOH/HCI mixture has been investigated in constant extension rate tests. Transgranular propagation occurs by cleavage on orthogonal planes. Solution additions of As, Hg, Pt and Quinoline indicate that hydrogen absorption causes the cleavage fracture while the effect of such additions upon the corrosion rate is unimportant. Pre-exposure embrittlement experiments on unstressed specimens revealed a reversible embrittlement phenomenon exhibiting the same cleavage fracture which could be prevented by ageing specimens after exposure but prior to fracture, a result that was further evidence for the embrittling role of absorbed hydrogen. INTRODUCTION

STRESS corrosion cracking of titanium alloys in a wide variety of environments has been fully reviewed, a Generally, much emphasis has been placed upon the important commercial ~ q- 13 alloys, Ti-6A1-4V and T i - 8 A l - l M o - l V , less on completely alloys and, perhaps, least of all on 13alloys. The work presented below reports experiments on the 13 alloy Ti-13Cr-lIV-3A! in methanolic solutions. Previous stress corrosion studies on this alloy have been confined mainly to aqueous solutions and a number of phenomenological points have been established. These are now briefly described. Fager and Spurr 2 demonstrated that Ti-13V-l 1Cr-3A! alloy in the single 13phase condition exhibits cleavage on or close to {100} planes when stress corroded in aqueous solutions. Ageing at 590°C reduces susceptibility but also lowers the fracture toughness. a It produces a fine widmanst/itten ~ phase precipitate in the alloy 1 and causes a change in stress corrosion fracture morphology from transgranular to intergranular. 3 Susceptibility in the alloy in the wire form has been reduced by cold work. 4 This was attributed to the creation of a heavily dislocated cellular structure and to the development of a longitudinal < I l0 > texture in the wire which could be expected to raise the {100} cleavage nucleation stress. The discontinuous nature of the crack propagation process was demonstrated by Katz and Gerberich 5 who analyzed their acoustic emission data as indicating a discontinuous crack propagation mode, consisting of a slow 'triggering process' and a rapid cleavage process. In neutral aqueous solutions repassivation appears to exert a considerable influence in the same way as that observed in ~ alloys, e In constant extension rate experiments stress corrosion crack propagation occurs over a narrow range of crosshead speeds. ~ The crack velocity varies with potential in a way similar to that observed in ~ alloys, s In chloride solutions a maximum propagation rate is observed at -- 500 mV(SCE). This value is 0 mV(SCE) in iodide solutions. *Manuscript received April 1979. 799

800

R.W. LvcEa-rand J. C. SCULLY

I n the work presented below experiments using the c o n s t a n t extension rate m e t h o d and reversible embrittlement experiments are interpreted, together with the results o f other workers, as showing that the observed t r a n s g r a n u l a r cleavage is caused by absorbed hydrogen. EXPERIMENTAL METHOD The material was in the form of 0.5 mm thick sheet in the mill annealed, descaled and pickled condition. The tensile stress of specimens annealed after manufacture at 750°C for 1 h was 780 MN m-2, grain size 80 I~m VHN 325 and elongation-to-failure 15-17~0 for a 25 mm gauge length. The principal environment was a solution that has been much used on a alloys. 6 It consisted of AnalaR Grade CHsOH containing 1 vol. ~o AnalaR grade concentrated HCI solution (35.4~0HC1). The mixture contained thereby 0.95 wt. ~.H~O and 0.52 wt. %HCI. Two types of specimens were employed, both cut in the same longitudinal direction from a single sheet of material which contained 13.2V, 10.2Cr, 3.3A1, 0.15Fe, 0.03C. Single edge notched (SEN) specimens were used to study the effect of a number of variables upon crack velocity, fracture morphology, and maximum load and elongation-to-failure. Unnotched tensile specimens with a 25 mm gauge length were used for pre-cxposure reversible embrittlement experiments. Crack velocity was measured on SEN specimens subjected to a range of crosshead speeds on an Instron Tensile Testing Machine. Specimens measured 120 × 16 mm with a notch 2 rnm deep subtending an angle of 60°. A graduated scale with markings 2 mm apart fixed across the specimen in the notched region was examined optically through the wall of a glass-sided beaker with a polyethylene base which was glued around the specimen and then filled with the solution. The specimen was fitted to the testing machine so as to give an 80 mm distance between the grips. The time to traverse the 2 mm distances was measured after the crack had propagated across the first 2 mm distance from the notch tip. Only one crack was initiated. The final part of the fracture occurred by ductile shear in a direction at 30° to the tensile axis. All specimens were examined after fracture by scanning electron microscopy in order to observe fractographic features and also to ensure that the crack front had maintained a front orthogonal to the direction of propagation. The latter point was an important supplement to the visual observations. A few specimens in the form of U-bends were broken in order to augment the fractographic detaiiobtained from the SEN specimens. EXPERIMENTAL RESULTS The experimental results are presented u n d e r two headings: c o n s t a n t extension rate experiments a n d pre-exposure experiments.

Constant extension rate experiments Results obtained for specimens b r o k e n in air and in the methanolic solution u n d e r open circuit conditions are shown in Table 1. Susceptibility can be characterized with respect to four variables: (i) m a x i m u m sustained load, (ii) elongation-to-failure, (iii) crack velocity, a n d (iv) the a m o u n t of cleavage in the fracture surface. The values o f velocity shown i n Table 1 are the arithmetic m e a n velocities calculated across the w i d t h o f the scale 2-12 m m from the notch tip. The times taken to traverse the individual 2 m m intervals were generally n o t c o n s t a n t at any crosshead speed. A typical set o f Values is shown in Table 2. All the values shown in the graphs a n d Tables presented below were calculated in the same way as for Table 1. I n Table 1 it c a n be seen that the m a x i m u m load i n solution sustained was d e p e n d e n t u p o n the crosshead speed. This was further confirmed i n a series o f experim e n t s i n which the load was varied by pre-loading the specimen in air prior to the addition o f the solution to the beaker a n d measuring the crack velocity at two crosshead speeds: 330 and 830 n m s -1. The results are shown in Fig. 1. A linear relationship

Stress corrosion crack propagation in Ti-13V-I 1Cr-3AI alloy

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802 TABLE 2.

TIMES FOR CRACKS TO TRAVERSE 9, n u n INTERVALS AT VARIOUS CROSSHEAD SPEEDS

Crosshead speed

1

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was observed for both crosshead speeds with slopes determined by a least squares method of 2.6 + 0.2 and 2.4 ± 0.3 respectively. From these results it is clear that the relationships between crack velocity and the maximum load is independent of the crosshead speed over the range that has been examined. The response of a propagating crack to changes in crosshead speed was examined by lowering the initial crosshead speed after the propagating crack front had reached the first 2 mm mark, corresponding to the point at which crack velocity measurements were started. The initial range of crosshead speeds was 330 nm s-X-83 v.m s -1. Three lower crosshead speeds were used: 330 and 830 nm s -1 and 1.67 i~m s-L The results are shown in Fig. 2 as the arithmetic averages of the velocities obtained for each series of experiments for any combination ofcrosshead speeds. It can be seen that the crack velocity is strongly dependent upon the initial crosshead speed. Lowering

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the crosshead speed resulted in a fall in crack velocity, as would be expected, but the final velocity was at all times higher than that observed in single crosshead speed experiments employing the second, lower, crosshead speed. Thus the final crack velocity appeared to be jointly dependent upon both the initial and final crosshead speeds. This was borne out by experiments in which the second crosshead speed was reduced to zero, i.e. the crosshead movement was stopped after the crack front had reached the initial 2 mm mark. The results are shown in Fig. 3. The velocities are lower than those shown in Fig. 2, a result that confirms that there is a partial dependence of crack velocity upon the second crosshead speed. These results are reviewed in the Discussion section. In dynamic straining experiments the influence of the machine in causing strain is opposed by the work hardening of the alloy and the balance between the two determines the effective strain-rate at the crack tip. In order to assess this with respect to the experiments employing both constant and altered crosshead speeds the strain hardening coefficient n was measured. A true stress/true strain curve was determined in air. A least squares determinati0n was performed on the region of stable plastic deformation. The value of n was determined to be 0.077 -t- 0.002. Fractographically the alloy exhibited transgranular cleavage, in agreement with previous work. A typical fractograph from a U-bend specimen is shown in Fig. 4. It was similar to fractographs obtained from constant extension rate experiments, an example of which is shown in Fig. 5. The interplanar angles were measured as 90 + 3°, a value that is consistent with previous reports of cleavage on or near to {100} planes. Fracture surfaces from specimens broken in constant extension rate tests showed increasing amounts of tearing as the value of the cross.head speed employed was increased. A ductile fracture is shown in Fig. 6. From Table 1 and fractographic examination, stress corrosion cracking occurred at crosshead speeds ~ 83 Izm s-t. The effect of solution variables was mainly examined at three crosshead speeds: 330 n m s -z and 3.3 and 16.7 i~m s-z. Viscosity. The solution viscosity was increased by adding glycerol to the methanol while ensuring that the concentration of HCI was unaltered. The results of such changes upon cracking behaviour are shown in Fig. 7. At the highest value of crosshead speed

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FIG. 6.

F r a c t o g r a p h of specimen broken at a crosshead speed of 330 a m s -~. T h e ductile fracture is m a d e up of a large n u m b e r of small dimples. S.E.M.

F1o. 8. Fractograph from SEN sper, imen broken in CHaOH/HCI containing 100 mg 1-1 of HgCI2. The surface shows a considerable amount of cleavage. S.E.M. FIG. 9. Fractograph from SEN specimen broken in CHaOH/HCI containing 100 g 1-1 of HgCIz. The fracture surface has been etched by the solution and selective attack, in the form of slots, is evident. S.E.M.

FIG. 12. F r a c t o g r a p h of specimen polarized in C H 3 O H / H C I at 200 t.tA m m -2 for 1200 s a n d subsequently fractured in air at a crosshead speed o f 1.67 ~tm s -~. A considerable a m o u n t of cleavage has occurred.

Stress corrosion crack propagation in Ti-13V-I ICr-3A1 alloy

809

the velocity exhibited very little dependency upon viscosity, the slope being 0.06 -40.04, whereas at the lower values the velocity fell with increasing viscosity, the slopes being - - 0.3 4- 0.05 and -- 0.3 4- 0.I at 3.3 ~tm s -1 and 330 nm s -1 respectively. Table 3 shows the maximum load observed in the experiments. At ihe highest. crosshead speed cracking in glycerol-containing solutions was absent, a result that was confirmed by fractographic observation. At all three values of crosshead speed the alloy exhibited no cracking in glycerol alone. TABLE 3.

T H E EFFECT OF GLYCEROL ADDITIONS UPON MECHANICAL PROPERTIES AT THREE CROSSHEAD SPEEDS

Glycerol (ml) per 100 ml methanol 0 25 50 75 1~ 200

Crosshead speed 330 nms -x Elongation Maximum (ram) load (KN) 0.86 1.8 1.8 2.3 2.0 2.2

2.7 4.5 4.8 5.6 5.0 5.9

Crosshead speed 3.3 ~tm s-a Elongation Maximum (ram) load (KN) 1.3 2.6 2.7 2.7 2.6 2.9

3.1 5.6 6.1 6.0 6.1 6.1

Crosshead speed 16.7 izm s-1 Elongation Maximum (mm) load (KN) 2.7 -2.9 2.9 2.8 2.7

5.1 -5.9 6.2 5.9 6.4

Mercury. The results of additions of HgC12 to the methanol/HC1 solutior~ are shown in Table 4. Only at the lowest crosshead speed was the maximum load reduced as a result of the additions. Except at the highest crosshead speed the proportion of cleavage fracture observed was greatly increased by the additions. An example is shown in Fig. 8 which shows a very large proportion of cleavage fracture. Under identical testing conditions larger proportions of cleavage were seen in fractures obtained in mercury salt-containing methanolic solutions than in mercury salt-free solutions, although the crack velocity was lower. At high levels of mercury salt additions surface corrosion was observed in the form of an unidentified black sludge and yellow colouration of the solution. Cleavage surfaces obtained in such solutions exhibited an etched appearance, an example of which is shown in Fig. 9. A small number of experiments were done in mercury. The failure load was considerably higher than in the methanolic solutions. At 330 n m s -x, for example, the load was increased by a factor of 2. The elongation-to-failure was slightly higher in mercury than in the methanolic solution at the same crosshead speed. The crack velocity was too high to be measured visually although an estimate could be made from the fall in load. Fractographically the cleavage observed was identical to that seen in methanolic solutions. Arsenic. The results of NaAsO2 additions are shown in Table 5. As with low concentrations of HgCI2 the effect of NaAsO2 additions was to increase the proportion of cleavage in the fracture at low values of crosshead speed while also causing a reduction in crack velocity. Platinum. The results of H2PtCI4 additions are shown in Table 6. The proportion of cleavage fracture was much reduced as compared to the methanol/HCl solution with no addition.

810

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Stress corrosion crack propagation in Ti-13V-11Cr-3AI alloy TABLE5.

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MECHANICAL PROPERTIES AND CRACK VELOCITY MEASURED AT THREE CROSSHEAD SPEEDS IN

SOLUTIONSCONTAINING300 ppm OF As ADDEDAS NaAsO= MeOH Crosshead speed 330 n m s -~ 1.67 v.m s-a 8.3 am s-= 16.7 i~m s-~

TABLE6.

HCI

Velocity (l~m s-~)

Elongation

Load (KN)

38 160 360 390

860 ~m !.2 mm 2.4 mm 2.7 mm

2.7 2.7 3.1 5.1

MeOH HC1 + NaAsO= Velocity Elongation Load (~m s-~) (ram) (KN) 14 45 270 520

1.0 1.5 2.3 2.3

1.7 3.1 4.7 5.6

MECHANICAL PROPERTIES AND CRACK VELOCITY MEASURED AT THREE CROSSHEAD SPEEDS IN SOLUTIONS CONTAINING 5 0 0 p p m OF Pt ADDED AS H ~ P t C I e

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Velocity (izm s-1)

MeOH HC! Elongation (mm)

Load (KN)

330 n m s -1 1.67 am s-~ 8.3 I~m s-~

38 160 360

0.86 1.2 2.4

2.7 2.7 3.1

MeOH HCI -t- H=PtCIe Velocity Elongation Load (~m s-a) (ram) (KN) 10 34 151

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The effects of H~PtCIe and NaAsO= additions upon the weight loss/time behaviour of the alloy are shown in Fig. 10. The effect of HgClz additions are shown in Table 8. After 3/4 days the rate of attack fell. Specimens were extensively pitted.

812

R.W.

LYCETT a n d J. C. SCULLY

Quinoline. The results of quinoline additions are shown in Table 9. At crosshead speed values < 16 ~tm s -1 the effect of such additions was to lower the crack velocity. At higher values the crack velocity was increased. The proportion of cleavage in all the fractures was reduced by the quinoline additions. HF. A small number of experiments were done at a crosshead speed of 330 nm s-L The results are shown in Table 7. The solution containing HCI and H F produced a lower maximum load and elongation-to-failure than the solution containing only HCI. In H F alone, however, these measured values were much higher. The HC1/HF mixture produced more cleavage than that obtained in the methanol/HCl solution while in the methanol/HF solution no stress corrosion fracture was observed. No stress corrosion was visible in methanol/H~SO4 experiments. TABLE 7.

MECHANICAL PROPERTIES AND CRACK VELOCITY MEASURED AT A CROSSHEAD SPEED OF 330 n m s -a IN SEVERAL SOLUTIONS

Solution used 100 pts. methanol, plus I pt. HCI

0.5 pts. HCI 0.5 pts. HF I pt. HF l pt. H2SO4

Average elongation 1.03 mm 860 ~tm 2.39 mm 2.93 mm

Average maximum Average crack load (KN) velocity I~m s-1

2.68 2.30 6.1 6.68

38 8 10 7

Pre-exposure experiments Unstressed specimens were exposed under open circuit conditions to the methanol/ HCI mixture for periods of time up to 150 days. Subsequently they were washed, dried and broken in air at a crosshead speed of 1.67 i.tm s -1. All the specimens fractured had elongations-to-failure within the scatter of values of unexposed specimens. There was no evidence of any cleavage in the resulting fractographs. Similar experiments were done in which specimens were exposed for periods of 1200s and 6000s while subjected to anodic current densities. The elongations-tofailure of specimens subsequently broken in air as a function of imposed anodic current density are shown in Fig. 11. Polarization at low current densities caused a reduction in elongation-to-failure to ca. 20 ~o of the air value and less. Higher current densities had little extra effect. Increasing the time of polarization had a marked additional effect. During polarization specimens turned black and the solution became green/yellow in colour. Fracture surfaces o f specimens thus embrittled exhibited extensive amounts of cleavage. A typical example is shown in Fig. 12. One set of specimens polarized at 200 ~tA mm -2 for 1200 s were stored at 22°C and -- 196°C prior to fracture at room temperature. The elongations-to-failure are shown in Table 10. It can be seen that the degree of embrittlement decreased as a function of time, a result that illustrates clearly the reversibility of the embrittling process. This was further emphasized by the decreasing amount of cleavage observed in broken specimens which is also indicated in Table 10. The proportion was obtained by imposing a grid on fractographs and counting the proportion o f squares covering cleaved areas. While such a simple two dimensional analysis of a three dimensional

Stress corrosion crack propagation in Ti-13V-I 1Cr-3Al alloy

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814 TABLE 9.

T ~ WEIGHT LOss OF S P E C I M E N S EXPOSED TO CHaOH/HCi SOLUTION CONTAINING VARIOUS AMOUNTS OF HgCI:

Concentration of HgCI,

Weight loss (%) after exposure time (days) 1 2 3 4

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0.036 0.056 0.124 1.018 7.41

0.059 0.059 0.124 1.127 9.09

0.065 0.068 0.144 1.183 9.49

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FIG. 1I. The relationship between elongation-to-failure and applied anodic current density for unstressed specimens subsequently fractured in air at a crosshead speed of 1.67 ~tm s -1. fracture surface will give only an approximate measurement it is adequate for indicating marked trends. It was observed that the elongation-to-failure did not revert completely to the value obtained with unembrittled specimens. DISCUSSION The results presented above are all consistent with the hypothesis that the cleavage observed in the methanolic solutions is caused by the absorption of hydrogen into the volume of slowly deforming metal at the crack tip. The reversible embrittlement experiments showed clearly that an absorbed mobile species causes the reduction in elongation. Since hydrogen is commonly absorbed when Ti alloys dissolve in water and methanolic solutions there is every reason to suppose that such an occurrence takes place with Ti-13V-11Cr-3A1. To produce an embrittlement effect anodic polarization was required since it was not observed under open circuit conditions. This absence o f an embrittlement may have been the result of a very low hydrogen absorption rate occurring under open circuit conditions, although this point was not demonstrated. It might then be expected that exposure times longer than those employed

Stress corrosion crack propagation in Ti-13V-11Cr-3AI alloy

815

TABL~10. MECHANICALPROPERTIESOFSPECIMENSAGEDATTWOTEMPERATURESAFTERBEINGPOLARIZED AT 2001zA mm-~ FOR 1200 S Ageing temperature Ageing time (days) 0 1 4 8 27

22°C

-- 196°C

Percentage of Elongation-tocleavage on failure the fracture Range Average Range Average 350-580~m 720 i~m-1 mm 850-950izm 540--820~m 2.72-3.1mm

490 ~tm 840 ~tm 910 ~m 700 ixm 3.4 mm

63-74 16--27 20-31 10-17 0

70 21 25 13 0

Percentage of Elongation-tocleavage on failure the fracture Range Average Range Average 350--580t~m 550-790~tm 940 ttm-l.38 mm 1.2-3.9mm 2.9-3.8 mm

490 ~m 640 ~tm 1.1 rnm 2.3 mm 3.3 mm

63-74 19-23 11-15 0-7 0

70 21 12 2 0

(up to 150 days) would have demonstrated embrittlement even under open circuit conditions. Such an effect depends upon the relative rates of hydrogen absorption in comparison to hydrogen diffusion within the lattice. Unless the former is significantly higher than the latter then the necessary development of a volume of metal containing sufficient hydrogen to fail by cleavage when stressed will not occur. The effect of arsenic in increasing the amount of cleavage observed in fractures was evident and explainable by a hydrogen absorption phenomenon. A similar effect was seen in solutions containing Hg ~+. The effect of deposited Hg in delaying the hydrogen recombination stage of the hydrogen evolution reaction could explain the increased amount of cleavage that was observed. This latter effect could also possibly be attributed to the discharged H g atoms causing a liquid metal embrittlement effect. F r o m the results presented above no definite conclusion can be drawn. If, however, such an effect was operative then it might be expected that increasing the HgCla concentration would increase the embritflement effect, since it could be expected that this would be dependent upon the availability of H g z+ ions to be deposited. This did not occur. The effect of H2PtCI6 additions was similar to that observed in ~-Ti alloys, s Much less cleavage was observed. This change can be attributed to the effect if discharged Pt atoms at the crack tip increasing the hydrogen recombination rate and thereby reducing the flux o f absorbed hydrogen. With both Pt and As additions the important effect arises from the change in the flux o f absorbed H atoms. Increasing the corrosion rate does not necessarily increase the amount of cleavage that occurs, as can be seen by comparing the effects of Pt and Hg. Decreasing the corrosion rate does not necessarily increase the amount of cleavage, as can be seen by comparing the effects of As and quinoline. The few experiments done in H F and H2SO4 are consistent with observations made on a ac-Ti alloy, a HCI is necessary for H absorption. Cleavage was not observed in H F in the absence o f HCI, nor in HaSO4. The increased amount of cleavage seen in the H C I / H F mixture has not been reported previously. The important role of the chloride ion in promoting hydrogen entry has not been explained. I t has been suggested 8 that the increased susceptibility observed in solutions containing increased

816

R . W . LYCETT and J. C. SCULLY

amounts of CI-, Br- and I- could be the result in aqueous solutions of the rapid hydrolysis of the adsorbed halide ion monlayer at the crack tip, a process occurring on the metal surface and promoting hydrogen absorption. This process would occur less (if at all) in the presence of F- ion alone since fluoride salts hydrolize less readily. In methanol solutions an equivalent inethanolysis can be hypothesized. The crack propagation process under constant extension rate testing conditions consists of cleavage and ductile rupture processes occurring together. This can be depicted as:

v=vc+v where V = crack velocity, Vc = cleavage velocity and V d = ductile rupture velocity. The longer the time required to reach conditions for V¢ to occur the greater will be the proportion of V d. The cleavage process itself consists of an initiation stage and a propagation stage which together represent the time required for cleavage to occur. The time for cleavage t c can be depicted as: tc = t t + tp where tc = time for cleavage rupture, t~ = time for cleavage initiation, and tp = time for cleavage propagation. There is no evidence that the cleavage propagation rate varies and the main ratedetermining factor controlling cleavage will be ti which will be controlled by the flux of hydrogen atoms entering the alloy at the crack tip. In solutions containing Pt and quinoline this value is low and the more rapid ductile component predominates, resulting in high overall crack velocities. This is also true for the highest crosshead speeds in all solutions. When the flux of hydrogen entering the lattice is increased the time t~ is reduced and cleavage fracture occurs at an earlier stage, corresponding to a lower stress. The ductile component in the fracture will therefore be less and the overall crack velocity falls. This would explain the fall in velocity and the increased proportion of cleavage observed in solutions containing As and Hg. The function of the crack tip strain-rate is clearly of importance, as has been emphasized for ~-Ti alloysa and other alloys. The value of this will depend upon the crosshead movement and the creep-rate resulting from the instantaneous value of the load. The relative importance of these two factors depends upon the strain hardening coefficient of the alloy. The lower this value the lower will be the relative importance of the crosshead movement. In an ~-Ti alloy lowering the crosshead speed during cracking resulted in a fall in crack velocity to a value characteristic of the lower crosshead speed) In the Ti-13V-11Cr-3A1 alloy this was not the result. The initial higher crosshead speed always exerted a lasting influenced upon the value of velocity observed at the lower crosshead speed and also when the crosshead movement was stopped. These results shown in Figs. 2 and 3 confirm the results shown in Fig. 1 in which the overall velocity-determining role of the load was clearly demonstrated. In this alloy strain hardening is relatively unimportant compared to the instantaneous stress in determining the crack tip strain-rate.

Stress corrosion crack propagation in Ti-13V-11Cr-3AI alloy

817

N o attempt has been made to elucidate the mechanism by which hydrogen atoms promote cleavage in the stress corrosion cracking of the BBC lattice. Hydrides occur in [~ alloys only at very high hydrogen levels9 and embrittlement has been observed below these levels? In the [B-rff alloy embrittlement has been observed 1° at levels < 800 ppm in the form of transgranular cleavage. Furthermore, reversible embrittlement was observed. This effect was considered to occur without the formation of hydrides. 1° Ti-13V-llCr-3A1 has been reported 11 to be embrittled by hydrogen without the formation of hydrides. This has also been concluded for a Ti-18Mo [3 alloy) 2 Thus it appears that interstitial hydrogen can cause embrittlement in [~-Ti alloys without hydride formation. Since the detection of very small hydride precipitates in fracture surfaces can be very difficult, and since the possibility of hydride disappearance after fracture by lattice dissolution exists, it may be difficult to establish the mechanism by which absorbed hydrogen promotes transgranular cleavage during stress corrosion crack propagation in the Ti-13V-11Cr-3AI in methanolic solutions. What the above study has shown is that there is much evidence that absorbed hydrogen is responsible for the transgranular cleavage by which stress corrosion cracks propagate in Ti-13V-llCr-3A1 alloy in methanolic solutions. Acknowledgement---One of the authors (R.W.L.) was supported throughout this work by a Science

Research Council Studentship. REFERENCES 1. M. J. BLACKBURN,J. A. FEENEYand T. R. BECK,Advances in Corrosion Science and Technology (Edited by M. G. Fo~rrANAand R. W. STAEHLE),Vol. 3, p. 67. Plenum Press, New York (1973). 2. D. N. FAGERand W. F. SPtmR, Trans. Am. Soc. Metals 61,283 (1968). 3. J. A. FEENEYand M. J. BLACKBURN,The Theory of Stress Corrosion Cracking in Alloys (Edited by J. C. SCtrLLY),p. 355. NATO, Brussels (1971). 4. J. W. MUNrO~ and H. J. RACK, Corrosion 29, 146 (1973). 5. Y. KATZand W. W. GERBE~CH,Int. J. Fract. Mech. 6, 219 (1970). 6. J. C. SCULLYand D. T. POWELL,Corros. Sci. 10, 371 (1970). 7. T. R. BECK, M. J. BLACKnURNand M. O. SPEIDEL,Boeing Quarterly Progress Report No. 11, NASA No. Cr. 103236, March (1969). 8. J. C. SCULLYand T. A. ADEPOJU,Cortes. Sci. 17, 789 (1977). 9. V. A. LIVANOV,B. A. KOLACHOVand A. A. BUHANOVA,The Science, Technology and Application of Titanium (Edited by R. I . . I ~ E and N. E. PROMISEL),p. 561. Pergamon Press, London (1968). 10. J. J. DELuccIA,Final Report, NADC-76207-30 (Juno 1976). 11. W. R. HOLMAN,R. W. CRAWFORDand F. PAREDES,JR., Trans. Am. Inst. Min. metall. Engrs 233, 1836 (1965). 12. N. E. PATON,O. BUCKand J. A. WXLUAMS,Scr. Metall. 9, 687 (1975).