ASM Metals HandBook Volume 12 - Fractography
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ASM Metals HandBook Volume 12 - Fractography

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strongly influenced by the alloy, the microstructure, and whether the hydrogen is present in the lattice before testing or is 
introduced during the test. For example, a Ti-8Al-1Mo-1V alloy that was annealed at 1050 °C (1920 °F), cooled to 850 
°C (1560 °F), and water quenched to produce a coarse Widmanstätten structure exhibited cracking along the \u3b1-\u3b2 
interfaces when tested in 1-atm hydrogen gas at room temperature (Ref 137). The fracture surface, which exhibited crack-
arrest markings, is shown in Fig. 46(a). The arrest markings are believed to be due to the discontinuous crack propagation 
as a result of the repeated rupture of titanium hydride phase at the crack tip (Ref 137). Also, Fig. 46(b) shows a hydrogen 
embrittlement fracture in a Ti-5Al-2.5Sn alloy containing 90 ppm H that was \u3b2 processed at 1065 °C (1950 °F) and aged 
for 8 h at 950 °C (1740 °F). The fracture occurred by cleavage. 
Fig. 46 Examples of hydrogen-embrittled titanium alloys. (a) Hydrogen embrittlement fracture in a Ti-8Al-1Mo-
1V alloy in gaseous hydrogen. Note crack-arrest marks. Source: Ref 137. (b) Cleavage fracture in hydrogen-
embrittled Ti-5Al-2.5Sn alloy containing 90 ppm H. Source: Ref 141 
Cleavage was also the mode of fracture for a Ti-6Al-4V alloy having a microstructure consisting of a continuous, 
equiaxed \u3b1 phase with a fine, dispersed \u3b2 phase at the \u3b1 grain boundaries embrittled by exposure to hydrogen gas at a 
pressure of 1 atm (Fig. 47a). However, when the same Ti-6Al-4V alloy having a microstructure consisting of a medium, 
equiaxed \u3b1 phase with a continuous \u3b2 network was embrittled by 1-atm hydrogen gas, the fracture occurred by 
intergranular decohesion along the \u3b1-\u3b2 boundaries (Fig. 47b and c). 
Fig. 47 Influence of heat treatment and resulting microstructure on the fracture appearance of a hydrogen-
embrittled Ti-6A-4V alloy. Specimens tested in gaseous hydrogen at a pressure of 1 atm. (a) Transgranular 
fracture in a specimen heat treated at 705 °C (1300 °F) for 2 h, then air cooled. (b) Intergranular decohesion 
along \u3b1 -\u3b2 boundaries in a specimen heat treated at 955 °C (1750 °F) for 40 min, then stabilized. (c) Coarse 
acicular structure resulting from heating specimen at 1040 °C (1900 °F) for 40 min, followed by stabilizing. The 
relatively flat areas of the terraced structure are the prior-\u3b2 grain boundaries. See text for a discussion of the 
microstructures of these specimens. Source: Ref 142. 
Hydrogen Embrittlement of Aluminum. There is conclusive evidence (Ref 99, 100, 101, 102, 103, 104, 105, 106, 
107) that some aluminum alloys, such as 2124, 7050, 7075, and even 5083 (Ref 143), are embrittled by hydrogen and that 
the embrittlement is apparently due to some of the mechanisms already discussed, namely enhanced slip and trapping of 
hydrogen at precipitates within grain boundaries. The embrittlement in aluminum alloys depends on such variables as the 
microstructure, strain rate, and temperature. In general, underaged microstructures are more susceptible to hydrogen 
embrittlement than the peak or overaged structures. For the 7050 aluminum alloy, a low (0.01%) copper content renders 
all microstructures more susceptible to embrittlement than those of normal (2.1%) copper content (Ref 106). Also, 
hydrogen embrittlement in aluminum alloys is more likely to occur at lower strain rates and at lower temperatures. 
The effect of hydrogen on the fracture appearance in aluminum alloys can vary from no significant change in an 
embrittled 2124 alloy (Ref 99) to a dramatic change from the normal dimple rupture to a combination of cleavagelike 
transgranular fracture and intergranular decohesion in the high-strength 7050 (Ref 106) and 7075 (Ref 105) aluminum 
alloys. Figure 48 shows an example of a fracture in a hydrogen-embrittled (as measured by a 21% decrease in the 
reduction of area at fracture) 2124-UT (underaged temper: aged 4 h at 190 °C, or 375 °F) aluminum alloy. It can be seen 
that there is little difference in fracture appearance between the nonembrittled and embrittled specimens. However, when 
a low-copper (0.01%) 7050 in the peak-aged condition (aged 24 h at 120 °C, or 245 °F) is hydrogen embrittled, a 
cleavagelike transgranular fracture results (Fig. 49a). This same alloy in the underaged condition (aged 10 h at 100 °C, or 
212 °F) fails by a combination of intergranular decohesion and cleavagelike fracture (Fig. 49b). 
Fig. 48 Hydrogen-embrittled 2124-UT aluminum alloy that shows no significant change in the fracture 
appearance. (a) Not embrittled. (b) Hydrogen embrittled. Source: Ref 99 
Fig. 49 Effect of heat treatment on the fracture appearance of a hydrogen-embrittled low-copper 7050 
aluminum alloy. (a) Transgranular cleavagelike fracture in a peak-aged specimen. (b) Combined intergranular 
decohesion and transgranular cleavagelike fracture in an underaged specimen. Source: Ref 106 
The Effect of a Corrosive Environment. When a metal is exposed to a corrosive environment while under stress, 
SCC, which is a form of delayed failure, can occur. Corrosive environments include moist air; distilled and tap water; 
seawater; gaseous, ammonia and ammonia in solutions; solutions containing chlorides or nitrides; basic, acidic, and 
organic solutions; and molten salts. The susceptibility of a material to SCC depends on such variables as strength, 
microstructure, magnitude of the applied stress, grain orientation (longitudinal or short transverse) with respect to the 
principal applied stress, and the nature of the corrosive environment. Similar to the Kth in hydrogen embrittlement, there is 
also a threshold crack tip stress intensity factor, KISCC, below which a normally susceptible material at a certain strength, 
microstructure, and testing environment does not initiate or propagate stress-corrosion cracks. Stress-corrosion cracks 
normally initiate and propagate by tensile stress; however, compression-stress SCC has been observed in a 7075-T6 
aluminum alloy and a type 304 austenitic stainless steel (Ref 144). 
Stress-corrosion cracking is a complex phenomenon, and the basic fracture mechanisms are still not completely 
understood. Although such processes as dealloying (Ref 145, 146, 147, 148) in brass and anodic dissolution (Ref 149, 
150, 151) in other alloy systems are important SCC mechanisms, it is apparent that the principal SCC mechanism in 
steels, titanium, and aluminum alloys is hydrogen embrittlement (Ref 38, 100, 107, 137, 143, 152, 153, 154, 155, 156, 
157, 158, 159, 160, 161, 162, 163, 164, 165, 166). In these alloys, SCC occurs when the hydrogen generated as a result of 
corrosion diffuses into and embrittles the material. In these cases, SCC is used to describe the test or failure environment, 
rather than a unique fracture mechanism. 
Mechanisms of SCC. The basic processes that lead to SCC, especially in environments containing water, involve a 
series of events that begin with the rupture of a passive surface film usually an oxide), followed by metal dissolution, 
which results in the formation of a pit or crevice where a crack eventually initiates and propagates. When the passive film 
formed during exposure to the environment is ruptured by chemical attack or mechanical action (creep-strain), a clean, 
unoxidized metal surface is exposed. As a result of an electrochemical potential difference between the new exposed 
metal surface and the passive film, a small electrical current is generated between the anodic metal and the cathodic film. 
The relatively small area of the new metal surface compared to the large surface area of the surrounding passive film 
results in an unfavorable anode-to-cathode ratio. This causes a high local current density and induces high metal 
dissolution (anodic dissolution) at the anode as the new metal protects the adjacent film from corrosion; that is, the metal 
surface acts as a sacrificial