failures of welded titanium aircraft ducts

P e r g a m o n Engineering Failure Analysis, Vol 2, No. 4 pp. 257-273, 1995

Elsevier Science Ltd Printed in Great Britain

1350-6307/95 $9.50 + 0.00

1350-6307(95)00025-9

F A I L U R E S O F W E L D E D T I T A N I U M A I R C R A F T D U C T S

S. P. LYNCH,* B. HOLE * and TIMOTIUS PASANG*

*Aeronautical & Maritime Research Laboratory, Defence Science and Technology Organisation, Department of Defence, Melbourne, Australia; *Materials Evaluation Facility, Directorate of

Aviation Safety Regulation, Civil Aviation Authority, Canberra, Australia; *Nusantara Aircraft Industries Ltd (PT. IPTN), Bandung, Indonesia

(Received 2 August 1995)

Abstract--An in-service failure of a thin-walled titanium bleed-air duct from a wide-bodied commercial aircraft has been investigated. Cracking had occurred in the heat-affected zone (HAZ) adjacent to a circumferential weld joining two sections of the duct which was manufactured from commercially pure (grade 3) titanium sheet. Specimens were cut from the duct to include an intact weld and tested under known conditions (overload, fatigue, sustained loading) for comparison with the failed duct. Metallographic observations showed that cracking occurred through an acicular tr HAZ, and fractographic observations revealed brittle, cleavage-like cracking with occasional areas of fatigue striations for both the failed duct and fatigued specimens. These observations, and the absence of sustained-load cracking in test specimens, suggested that the in-service failure had occurred primarily by fatigue. Observa- tions also indicated that small, cleavage-like cracks had been present in the ducts prior to service, although whether these cracks were caused by overload tearing, stress-corrosion cracking, hot-salt cracking or fatigue was not clear. Other possible causes of cleavage-like cracking, e.g. the presence of hydrides as proposed for previous failures, contamination of HAZs by oxygen/nitrogen, are discussed. Possible ways of preventing further failures are then outlined.

1. INTRODUCTION

There have been a number of in-flight failures of pneumatic titanium ducts in wide-bodied, commercial aircraft, and all have been associated with cracking adjacent to welds. In some cases, failure of the ducts has compromised aircraft safety by blowing access panels from the aircraft, causing loss of cabin pressure, and disabling de-icing systems. In other cases, failures have resulted in loud bangs, then entry of hot air plus particles of insulation into the cabin compartment, thereby causing some concern and discomfort to passengers!

The ducts are typically 150-180 mm in diameter with wall thicknesses (t) from 0.5 to 1 mm, and are generally manufactured from grade 3 titanium sheet (composition:wt% maximum 0.350, 0.30Fe, 0.10C, 0.05N, 0.015H, remainder Ti: typical H levels are 0.003-0.006 [1]). Ducts are welded by a tungsten inert-gas process using commercial purity welding wire, and are not always stress-relieved after welding. A proof pressure test (2.65 MPa) and inspection by a dye-penetrant technique are subsequently carried out. Portions of some ducts are then sprayed with a gold paint for protection from hydraulic fluid that can cause environmentally induced cracking at elevated temperatures. The gold-painting process requires subsequent heating to - 4 5 0 °C in air for 10 min, producing a blue heat-tint on the uncoated areas of the duct. In service, the ducts carry pressurised air (~310 kPa with occasional surges to higher pressures) initially at 20 °C and then increasing to temperatures in the range 120-175 °C.

Examination of previous failures by the aircraft manufacturers [1] showed that cleavage-like cracking had occurred in the heat-affected zone (HAZ) adjacent to welds. Metallographic observations showed that the weld and HAZ were characterised by an acicular tr-phase separated by thin strips of fl-phase, with extensive hydride

Copyright (~) 1995 Commonwealth of Australia

257

258 s.P. LYNCH et al.

precipitation extending from the inside surface to the mid-thickness position. The base metal consisted of an equiaxed o~-phase with small amounts of globular /3, and isolated hydride needles. It was suggested [1] that hydrogen diffused from the base metal (30-60 ppmH) to the HAZ (near the inside surfaces) at elevated temperatures due to the presence of high residual tensile stresses at this position. (Hydrogen could also diffuse towards the HAZ under the influence of a thermal gradient during welding [2].) Crack initiation and growth was envisaged as occurring primarily by fracture of hydrides during pressurisations of the ducts at 20 °C, with hydrides subsequently dissolving as the operating temperature increased and then reforming at crack tips on cooling, so that further increments of cracking occurred on subsequent pressurisations [1].

In the present paper, a detailed examination of a recent (1993) duct failure is described. The characteristics of the failure are then compared with those of specimens (cut from intact areas of the duct) fractured under known conditions. It is proposed that crack growth in service occurs mainly by fatigue from pre-existing cracks, with hydrogen-assisted cracking probably playing only a minor role in this case.

2. EXAMINATION OF THE FAILED DUCT

For the failure described in detail in the following, the duct had been in service for about 26,600 h and had experienced 19,275 pressurisation cycles. The duct had fractured through the toe of a weld in the HAZ, where a curved 0.8 mm thick section had been welded to a straight 0.5 mm thick section (Figs 1 and 2). The duct surfaces were a blotchy blue/purple/brown, although colours on the inside surface were obscured by a grey deposit in many areas. The weld metal and areas up to 15 mm

LOCATION OF ~ ' ~ THE FAILURE ~ /

Fig. 1. Schematic diagram showing part of the duct system in an aircraft with the failure site arrowed.

Failures of welded titanium aircraft ducts 259

Fig. 2. Macroscopic views of failed duct showing fracture in the HAZ adjacent to the weld. The part of the fracture S-S was covered by deposits.

from the weld exhibited darker, more uniform blue/purple/brown/yellow tints than the rest of the duct, with the inside surface more yellowish than the outside. Polishing the surfaces of the ducts with alumina paste (Brasso ®) removed the colours from the main body of the duct before those around the weld.

The microstructure of the parent metal and HAZ material adjacent to cracks was similar to that described in the introduction for earlier failures, but there were fewer hydrides in the HAZ. Thus, the parent sheets exhibited equiaxed tr, whereas the thin-section H A Z exhibited coarse acicular tr, with the fracture path often following the interfaces of the acicular tr (Fig. 3). There was no significant difference in the hardness (VH2.5 ~ 266) of the parent sheets and HAZs (which extended 3-4 mm from the weld). Fracture surfaces exhibited no signs of macroscopic ductility and there was no obvious demarcation between subcritical cracking and overload regions. However, the fracture surface around about one-fifth of the circumference (Fig. 2) was covered by deposits, indicating that a crack had been present for some time. Fracture surfaces were cleaned by repeated stripping of cellulose acetate replicas and examined with a binocular microscope (up to x 60). Fractures had a cleavage-like appearance, and small areas adjacent to the inner surface were blue/straw-coloured. River lines and steps on the fracture surface suggested that cracks had initiated at multiple sites around the inside surface.

Examination of fracture surfaces by SEM showed that the detailed appearance of cleavage-like fracture surfaces varied widely from area to area. The fracture topography was complex (and hence stereographic pairs of micrographs were taken routinely) and related to the underlying acicular a~ microstructure (Fig. 4). At high magnifications, small, shallow dimples were observed on some facets [Fig. 5(a)]. Other areas were stepped, with some step faces exhibiting fine-scale terraces and tear ridges, and other step faces exhibiting a relatively featureless, more classical, cleavage-like appearance [Fig. 5(b)]. Fatigue striations, with a fairly regular spacing (1-2/~m), were also apparent in a few areas [Fig. 5(c) and (d)]. Secondary cracking, probably along particular crystallographic planes, was also observed in some areas [Fig. 5(c)]. Isolated regions were relatively ductile, exhibiting large dimples or flutes (elongated dimples) [Fig. 5(e)].

The weld at the other end of the curved part of the duct was examined for signs of

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264 S.P. LYNCH et al.

Fig. 5. SEM of fracture surface of failed duct showing wide variation in detailed appearance: (a) small, shallow dimples [boxed area in 4(b)], (b) fine-scale terraces and relatively featureless steps, (c) shallow dimples and fatigue striations, (d) terraces and secondary cracking, and (e) large flutes.

cracking by f luorescent dye penetrants . No cracks were found initially but sectioning the duct to gain ease of access and fur ther meticulous examinat ion revealed a series of fine cracks with a max imum depth - 0 . 3 3 mm. Metal lographic sections also revealed cracks (some as small as - 0 . 0 5 m m deep). Howeve r , only small, isolated hydride needles were detected a round cracks (Fig. 6). Fractur ing pieces cut f rom the weld also showed that blue-t inted (pre-existing) cracks - 0 . 2 5 m m deep and 0.5 m m long had been present but had not been detected by the dye-pene t ran t inspection.

Fig. 6. Optical micrograph of section through ~'intact" weld showing small crack initiated from the HAZ-weld-metal junction on the inside surface. Inset shows small hydride needles present in the weld metal.

Failures of welded titanium aircraft ducts

3. TESTS ON SPECIMENS CUT FROM DUCTS

265

Specimens (40 x 10 x t mm) containing an intact (or largely intact) weld were cut from the duct and notched so that testing in cantilever or three-point bending produced crack growth in a circumferential direction in the HAZs. Specimens were notched so that cracking occurred either in the H A Z of the thinner (0.5 mm) section of the duct (where the in-service failure occurred) or in the HAZ of the thicker (0.8 mm) section.

Fatigue fractures (produced at cyclic frequencies of 0.3 and 24 Hz in air at 20 °C) through the thin-section HAZ exhibited a wide variety of features, with well-defined fatigue striations present only in a few areas. Other areas exhibited terraces and small dimples or cleavage-like areas with river lines (Fig. 7). Polished and etched sections through cracks showed extensive crack-branching and secondary cracking, with cracks

Fig. 7(a)-(c). Caption overleaf.


Fig. 7. SEM of fatigue fracture surface for the thin-section HAZ showing: (a) cleavage-like appearance at low magnification, and (b)-(d) patches of fatigue striations, small dimples, and relatively featureless areas at high magnification.

often following acicular 0l interfaces (Fig. 8). For the thick-section H A Z , on the o ther hand, there was a fine needle-like (possibly martensitic) 0~ microstructure, fatigue fractures exhibited striations in most areas, and there was little crack branching (Figs 9 and 10).

Over load fractures for the thin-section H A Z were macroscopical ly brittle and exhibited cleavage-like regions and faceted dimpled areas. There were no obvious

Fig. 8. Optical micrograph of section through a thin-section HAZ specimen after fatigue, showing extensive crack branching with cracks following acicular 0l interfaces.


Fig. 9. SEM of fatigue fracture surface for the thick-section HAZ, showing fatigue striations in most areas.

differences in appearance for overload fractures produced at 20 °C and those produced at 175 °C (where any hydrides present should have dissolved) (Fig. 11). For the thick-section HAZ, overload fractures resulted in significant lateral contraction of the side surfaces of specimens and fracture surfaces exhibited large dimples/flutes (Fig. 12).

Fatigue pre-cracked specimens, loaded until "clicks" were heard (indicating the stress intensity factor was close to that for unstable fracture), and then left for about 6 weeks in air at 20 °C, did not show any signs of sustained-load cracking.

4. DISCUSSION

4.1. Pre-existing cracks

The observation that there were small blue-tinted areas of the fracture surfaces adjacent to the inside surface of the duct suggests that small cracks were present prior to service. Blue tinting on the fracture surfaces (and on the surfaces of ducts) is

Fig. 10. Optical micrograph of section through a thick-section HAZ specimen after fatigue, showing crack traversing the fine, needle-like 0r-phase.


Fig. 11. SEM of overload fracture (at 175 °C) in the thin-section HAZ, showing cleavage-like areas.

probably due to heating the ducts to 450 °C during the gold-coating process.* Ducts which are not gold-coated are not tinted, and heating portions of those ducts at 450 °C for 10 rain resulted in blue tints. Pre-existing cracks with cleavage-like fracture surfaces could be produced by various processes, as outlined in the following. Which process is most likely to be responsible for cracking is unclear.

4.1.1. Overload fractures. A small extent of overload cracking could have occurred due to the brittleness of the H A Z material, either after welding due to a high near-surface residual tensile stress or during proof testing. The axial stress, or, resulting f rom proof testing at a pressure, p , of 2.65 MPa should be - 2 0 0 M P a (or = pr/2t , where r is the tube radius and t is the wall thickness): this stress is about 50% of the proof strength, but would be locally higher due to the stress-concentrating profile of the weld and the superposition of tensile residual stresses. The brittleness of the thin-section H A Z is discussed further in Section 4.3.

4.1.2. Stress-corrosion cracking (SCC). SCC could have occurred due to the combined presence of residual tensile stresses and a specific environment. Environ- ments known to induce SCC in titanium alloys, especially in notched or pre-cracked material , include aqueous chloride solutions, methanol and ethylene glycol at ambient temperatures , and halogenated hydrocarbons at elevated temperatures [4]. Rapid SCC of titanium T I G weldments in trichloroethylene vapour has been reported [5]. The likelihood of the ducts being exposed to SCC environments prior to service is not known.

*The use of gold paint for protecting parts of the duct from fuel is somewhat surprising given that gold is one of a number of metals known to produce metal-induced embrittlement of titanium. Solid-metal-induced embrittlement of titanium alloys by gold (and silver) has been observed at temperatures as low as 230 °C [3]. However, no problems seem to have arisen from the use of gold paint on the ducts.


Fig. 12. SEM of overload fracture (at 20 °C) in the thick-section HAZ, showing equiaxed dimples.

4.1.3. Hot-salt cracking. Sustained-load cracking at elevated temperatures (285-425 °C) in the presence of chloride, bromide and iodide salts can occur in all commercial titanium alloys except grades 1, 2, 7, 11 and 12 [4]. Salt contamination from finger prints has been known to result in cracking during heat-treatment. Cracking of the ducts could have possibly occurred during heating to 450 °C (as part of the gold-coating process) if salt contamination were present. Salt contamination from finger prints, however, is unlikely on the inside surfaces of ducts where cracking occurred.

4.1.4. Fatigue. The ducts could have experienced significant cyclic stresses resulting in fatigue cracking if they were not well supported and subject to vibration or transportation after welding (and prior to gold coating and heating to 450 °C). Whether such events occurred is unknown but seems unlikely.

4.2. Cracking in service

Crack growth in service most probably occurred by fatigue since:

(i) the duct is probably subject to significant cyclic stresses, (ii) the appearance of fracture surfaces for the in-service failure was similar to that

for fatigued specimens, with both exhibiting cleavage-like features plus patches of fatigue striations, and

(iii) other fracture modes can largely be discounted, e.g. sustained-load fracture did not occur in test specimens, and there was no convincing evidence that fracture involved repetitive formation and fracture of hydrides.

The cyclic stresses responsible for fatigue could arise from the changes in pressure/temperature or from vibration. Axial stresses due to pressurisation would be

270 S.P. LYNUH et al .

low (~23 MPa) and other welds would have been subject to this stress but were not substantially cracked. Thus, factors other than or in addition to pressure cycles are probably involved. The fact that the cracked weld was close to a bend in the duct could result in higher stresses due to bending that would occur, for example, if changes in length due to temperature changes were constrained.

The presence of fatigue striations, albeit in only a few areas, is clear evidence that cracking occurred by fatigue. The absence of striations in most areas, for both the in-service failure and the fatigued specimens, is clearly associated with the brittleness of the thin-section HAZ which was characterised by a coarse acicular cr microstructure. (Striations were observed in most areas for fatigue cracks grown through the thick-section HAZ specimens.) In general, fatigue striations are difficult or impossible to detect when fatigue occurs through a brittle phase or along a brittle interface. Striations can also be difficult to detect when fatigue cracking occurs across lamellar microstructures since the fracture topography is dominated by the different fracture behaviour of the different phases. Many of the fine details on the fracture surfaces of the duct are probably associated with the presence of thin strips of [3 between the o:-phase. Striations are also not resolved on fatigue fracture surfaces produced at low AK in titanium alloys with "ductile" microstructures: cleavage-like areas, shallow furrows, and blocky, undercut steps are observed [6-10].

Sustained-load cracking (SLC) at 20°C in air (and vacuum) occurs in many titanium alloys due to the presence of internal hydrogen [11-15]. SLC occurs most readily in o:-/3 titanium alloys such as Ti6AI4V in which hydrogen levels as low as 8 ppm can cause cracking [11]. Crack growth rates are higher for higher hydrogen contents and, for a given hydrogen content, are highest at around ambient temperatures (-23 to 27 °C) [14, 15]. SLC is usually attributed to diffusion of hydrogen to crack tips followed by formation and fracture of hydrides [14, 15]. For c~-/3 alloys with <~ 100 ppm hydrogen, cracking occurs by cleavage of the 0:-phase (on basal or near-basal planes), with increasing extents of microvoid coalescence at higher K values. For alloys with >~ 100 ppm hydrogen, brittle cracking occurs along a~-/3 interfaces as well as by cleavage of the ot-phase [11, 13]. The /3-phase has a higher hydrogen solubility and diffusion rate than the o~-phase so that hydrides tend to form at or-/3 interfaces. There is, however, little fractographic evidence that SLC in air involves repeated formation and fracture of hydrides and, hence, it has been proposed by some workers that solute hydrogen is somehow responsible [11-13]. In-situ TEM observations of cracking in thin foils indicate that either solute hydrogen or hydrides can facilitate cracking depending on the circumstances [16].

Recent failures of welded Ti6A14V pressure vessels used in the Space Shuttle have been attributed to SLC (plus fatigue) initiated from gas pores in the weld metal [17]. However, commercially pure titanium appears to be much more resistant than titanium alloys to SLC, and cracking has only occasionally been reported (in material with > 100 ppm hydrogen) [18]. For pre-cracked grade 2 and grade 12 titanium specimens tested at slow strain rates--a more severe test than sustained loading-- slow, subcritical crack growth was observed only for material with ~>400ppm hydrogen [19]. The absence of SLC at 20 °C in specimens cut from the ducts, where hydrogen contents in the HAZ are probably < 100 ppm, is therefore consistent with previous studies, and suggests that SLC does not contribute to cracking in the duct.

Failures of welded titanium-alloy pressure vessels have also been attributed to a high density of pre-existing hydrides at fusion-zone-HAZ interfaces [20]. Extensive hydride precipitation in these cases probably occurred because hydrogen diffusion to the interface was promoted by a difference in composition (and hydrogen activity) between the base metal (Ti6A14V) and the filler metal (CP Ti) [20]. Low-energy cleavage-like fractures are observed when there is an almost continuous fracture path along brittle hydrides [21-23].

For the duct examined in the present study, metallographic observations showed some small, pre-existing hydrides, and the secondary cracks observed on fracture surfaces are probably associated with these hydrides. There was, however, no


fractographic evidence that cracking primarily involved repetitive formation and fracture of hydrides, as proposed [1] for previous duct failures. Such a process would probably produce striations on fracture surfaces (due to crack arrest and blunting), as observed for SLC of Ti alloys in hydrogen gas [24, 25] and for SLC of zirconium alloys containing solute hydrogen [26]. The observation that cleavage-like overload fracture of the thin-section HAZ occurred at 175 °C as well as at 20 °C also suggests that hydrides were not responsible for embrittlement since hydrides should have dissolved at 175 °C.

4.3. Brittleness of thin-section HAZ

The brittleness of the thin-section (0.5 mm) HAZ compared with the thick-section (0.8 mm) HAZ is almost certainly associated with the differences in microstructure arising from differences in temperature/cooling rate of the two sections during welding. The thinner section would probably reach higher temperatures, and would cool more slowly because heat would be conducted away more slowly along the thinner section than along the thicker section. Convective cooling probably makes a relatively small contribution, resulting in faster cooling of the near-surface material. These factors would account for the formation of coarse acicular tr in the thin-section HAZ, and the formation of a fine needle-like (possibly martensitic) tr, with some near-surface acicular tr, in the thick-section HAZ.

Previous work [27] has shown that slow cooling rates from above the /3 transus (~940 °C) produce a coarse acicular c~ with serrated boundaries, whereas fast cooling rates produce a fine needle and lamellar-shaped or-phase. Furthermore, ductilities were lower for the slowly cooled material. Studies [9] have also shown that ductilities and fatigue properties of thin welded titanium tubes can be degraded by oxygen contamination. Contamination by oxygen (and nitrogen or carbon) can occur due to inadequate inert-gas shielding during welding or inadequate cleaning prior to welding. Fatigue fracture surfaces for an uncontaminated serrated c~ HAZ microstructure exhibited striations whereas fatigue fractures for an oxygen-contaminated acicular tr microstructure exhibited brittle, cleavage-like fractures without striations [9]. The formation of the acicular tr microstructure also appeared to be promoted by the oxygen contamination.

The above studies suggest that the absence of any macroscopic ductility and cleavage-like fractures of the thin-section HAZ could be due the combined effect of the acicular tr microstructure (with an intrinsically low ductility) and contamination by oxygen (and possibly nitrogen and carbon). More contamination could occur for the thin section because it reaches a higher temperature than the thick section and because a given amount of contamination is absorbed by a smaller volume of material. Interference colours due to oxide films around welds can give some indication of the degree of contamination that has occurred [28, 29]. Light blue through grey blue to grey indicates very heavy (unacceptable) contamination. The fact that the failed duct was generally tinted blue due to heating to 450 °C in air after welding (as part of the gold painting process) complicates the interpretation of the colours around the welds. [Heating to 450 °C after welding probably does not significantly increase contamination by oxygen and only produces thickening of the oxide film: rapid absorption of oxygen (and nitrogen) occurs only at temperatures > 500 °C]. Interpretation is also complicated because titanium dissolves its own oxide at high temperatures and because oxide-growth kinetics can depend on whether the initial oxide forms at high temperatures or at room temperature. However, the darker blue tints (thicker oxide) nearer to welds in some areas suggest that some contamination occurred during welding. The variability of colours around welds also suggests that the efficacy of the inert-gas shielding was variable. Large amounts of oxygen/ nitrogen contamination during welding would be expected to increase the hardness of the weld zones, but no significant differences in hardness between the HAZs and the parent sheet were observed. Thus, there is some uncertainty regarding the causes of the brittleness of the thin-section HAZ and further studies are warranted.


4.4. R e m e d i a l measures

Previously r e c o m m e n d e d measures [1] to prevent failures have included:

(i) stress relieving the ducts after welding, and (ii) replacing the ducts with ones manufac tu red f rom the lower-s t rength grade 2

t i tanium, and increasing wall thickness.

These measures were apparent ly based on the premise that hydrogen-assis ted cracking (of one form or another) was responsible for failures. Stress relieving would decrease tensile residual stresses so that there was less driving force for hydrogen diffusion to the H A Z . Grade 2 t i tanium has a lower specified max imum hydrogen content* than grade 3 t i tanium. These measures may well reduce the l ikelihood of failure but are apparent ly not sufficient since failures have occurred in stress-relieved grade 3 ducts with low ( 3 0 p p m ) hydrogen contents [30] and in grade 2 ducts [1], A n o t h e r r e c o m m e n d e d measure [1] was to move the position of the circumferential weld fur ther away f rom the bend, presumably on the basis that welds might experience lower bending stresses.

The present investigation indicates that fatigue is the pr imary fracture mode (at least for the failure studied) and, hence, moving the posit ion of the weld would probably reduce any cyclic bending stresses experienced by the weld. However , increasing the fatigue pe r fo rmance of the welds by eliminating pre-cracks, improving the weld profile, and increasing the ductility of the thin-section H A Z are also likely to be necessary to prevent failures. If oxygen contamina t ion and acicular c~ microstructures are the reasons for poor ductility, then changes in welding parameters and procedures might be necessary.

A n ability to detect smaller cracks than the dye-pene t ran t me thod current ly used would be desirable to prevent precracked ducts being used and to detect cracks at inspections during service. Exper iments showed that an X-ray technique was capable of detect ing cracks with depths greater than 0.13 mm. F o a m rubber was used to hold the film in posit ion, and about one-sixth of the interior c i rcumference was radio- g raphed at a time (using 40 kV, 7.5 mA-min , 800 m m focal film distance). Installing addit ional heat sensors a round the ducts to detect leaks before catastrophic failure would also be advisable.

R E F E R E N C E S

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*Note that reducing hydrogen contents to very low values requires vacuum annealing at -760 °C for -8 h and is not economically viable.


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failures of welded titanium aircraft ducts

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