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'Fatigue crack initiation in aerospace aluminum alloys, components and structures

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The concept of self-healing fatigue cracks in aluminium alloys has aroused much interest due to a recent paper by Lumley et al. [1]. The results of fatigue tests, fractography, and microstructural analysis for an underaged (UA) and peak aged (PA) experimental Al-Cu-Mg-Ag alloy were interpreted as evidence for possible self-healing of small internal fatigue cracks in the UA material. The present paper reviews some fatigue studies on UA, PA and overaged (OA) aluminium alloys, and surveys fatigue crack initiation in aerospace aluminium alloys, components and structures. It is concluded that whether or not self-healing of fatigue cracks actually occurs in some aluminium alloys, it should be considered inapplicable in practice.
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Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands R.J.H. Wanhill
1 © Springer 2007
FATIGUE CRACK INITIATION IN AEROSPACE
ALUMINIUM ALLOYS, COMPONENTS AND
STRUCTURES
R.J.H. Wanhill
National Aerospace Laboratory NLR
P.O. Box 153, 8300 AD Emmeloord
The Netherlands
Tel: (031) 522-8294
e-mail: wanhill@nlr.nl
The concept of self-healing fatigue cracks in aluminium alloys has aroused much interest due to a recent paper
by Lumley et al. [1]. The results of fatigue tests, fractography, and microstructural analysis for an underaged
(UA) and peak aged (PA) experimental Al-Cu-Mg-Ag alloy were interpreted as evidence for possible self-
healing of small internal fatigue cracks in the UA material. The present paper reviews some fatigue studies on
UA, PA and overaged (OA) aluminium alloys, and surveys fatigue crack initiation in aerospace aluminium
alloys, components and structures. It is concluded that whether or not self-healing of fatigue cracks actually
occurs in some aluminium alloys, it should be considered inapplicable in practice.
Keywords: aluminium alloys, aerospace, fatigue, self-healing
1 Introduction
There is currently much interest in the concept of self-healing fatigue cracks in aluminium
alloys, owing to a recent paper by Lumley et al. [1]. The results of fatigue tests, fractography
and microstructural analysis for an experimental Al-Cu-Mg-Ag alloy in an underaged (UA)
and peak aged (PA) condition were interpreted as evidence for possible self-healing of small
internal fatigue cracks in the UA material.
The present paper first considers fatigue studies of UA, PA and also overaged (OA)
aluminium alloys, particularly the work of Lumley et al. [1] and Finney [2]. Then there are
surveys of fatigue crack initiation in aerospace aluminium alloys, components and structures.
The subsequent discussion shows that self-healing of small internal fatigue cracks, whether or
not it occurs in some aluminium alloys, should be considered inapplicable in practice to
aerospace aluminium alloys, components and structures.
2 The fatigue study by Lumley et al. [1]
Lumley et al. [1] conducted high-cycle reversed-stress unnotched fatigue tests on extruded
bars of an experimental silver-containing aluminium alloy with the composition Al-5.6Cu-
0.45Mg-0.45Ag-0.3Mn-0.18Zr by weight.
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands R.J.H. Wanhill
The alloy was in two heat-treatment conditions: underaged (UA) and T6 peak aged (PA). The
fatigue results are shown in figure 1 and demonstrate generally longer fatigue lives for the UA
material.
Figure: 1 Fatigue life data for the experimental Al–Cu–Mg–Ag alloy [1]
Scanning Electron Microscope fractography showed that fatigue crack initiation was internal
[1, 3], which is unusual for aluminium alloys. Microstructural analysis by Transmission
Electron Microscopy led to the conclusion that during fatigue the dislocations in the UA
material became saturated with free solute copper atoms. This would be expected to reduce
the general mobility of dislocations, but some might transport solute to small initiating fatigue
cracks and close them. This self-healing concept is illustrated schematically in figure 2.
Figure 2: Proposed self-healing mechanism for the UA experimental Al–Cu–Mg–Ag alloy [1]
Lumley et al. [1] proposed that one or both effects, i.e. a general reduction in dislocation
mobility and the closing of small internal fatigue cracks could delay crack initiation and result
in longer fatigue lives for the UA material. They also suggested a third possibility, that crack
initiation might be delayed by localised matrix hardening due to dynamic precipitation.
However, it is the self-healing concept that has gained the most attention.
3 The fatigue study by Finney [2]
Information on the effect of ageing on aluminium alloy fatigue properties was limited at the
time of Finney’s study [2], and is summarised in table 1. The only consistency is that OA
materials never had the best fatigue properties.
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Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands R.J.H. Wanhill
Table 1: Summary of ageing effects on aluminium alloy fatigue known before 1969
Alloy type Fatigue life ranking* Reference
D.T.D. 683: commercial AlZnMgCu alloy PA > OA >ST Broom et al. [4]
Al-7Mg: experimental alloy ST = UA = PA Stubbington [5]
2014: commercial AlCuMg alloy ST = UA > PA & OA Form [6]
2014: commercial AlCuMg alloy UA > PA Fricke [7]
Al-7.5Zn-2.5Mg: experimental alloy ST = PA > OA Stubbington [8]
* ST: solution treated; UA: underaged; PA: peak aged; OA: overaged
Owing to the limited and inconsistent prior data, Finney set up a study of UA, PA and OA
experimental and commercial variations of two AlCuMg alloys.
The experimental alloys were Al-2.35Cu-1.3Mg and Al-4.18Cu-0.69Mg by weight. The
commercial alloys were D.T.D. 5014 (equivalent to AA 2618) and B.S. 2L.65 (similar to AA
2014). Their compositions were respectively Al-2.4Cu-1.51Mg-0.87Fe-0.86Ni-0.24Si-
0.03Mn-0.02Zn and Al-4.0Cu-0.72Mg-0.2Fe-0.67Si-0.77Mn-0.06Zn by weight. The main
differences between the experimental and commercial alloys were the presence of iron, nickel
and silicon in D.T.D. 5014, and iron, silicon and manganese in 2L.65.
Finney conducted high-cycle rotating cantilever unnotched fatigue tests on the four alloys in
UA, PA and OA conditions. Figure 3 shows the fatigue results from 105 cycles onward. There
are several points to note:
Figure 3: Finney’s [2] fatigue life data from 105 cycles onward
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Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands R.J.H. Wanhill
4 © Springer 2007
(1) The fatigue rankings for the experimental alloys were UA > PA > OA.
(2) The fatigue rankings for the commercial alloys were different:
The Al-2.4Cu-1.51Mg (D.T.D. 5014) alloy fatigue rankings were PA > UA > OA out
to 107 cycles. From 107 to 108 cycles the differences gradually vanished.
The Al-4.0Cu-0.72Mg (2L.65) alloy fatigue rankings were UA = PA > OA.
(3) The fatigue strengths at 108 cycles are given in table 2.
Table 2: Summary of fatigue strengths from Finney’s test data [2]
Fatigue strength at 108 cycles (MPa)
UA PA OA
Al-2.35Cu-1.3Mg : experimental 166 133 93
Al-2.4Cu-1.51Mg : commercial 135 137 135
Al-4.18Cu-0.69Mg : experimental 169 ~ 140 88
Al-4.0Cu-0.72Mg : commercial 220 220 171
From table 2 it is seen that:
The experimental alloys had similar fatigue strengths in all three conditions.
The commercial Al-4.0Cu-0.72Mg (2L.65) alloy had significantly higher fatigue
strengths than the commercial Al-2.4Cu-1.51Mg (D.T.D. 5014) alloy in all three
conditions.
In the UA condition the Al-2.35Cu-1.3Mg experimental alloy had a higher fatigue
strength than its commercial equivalent, but the reverse held for the experimental Al-
4.18Cu-0.69Mg and its commercial equivalent.
As for the previous investigations, table 1, the only overall consistency is that OA
materials never had the best fatigue properties.
Finney used optical fractography to qualitatively determine the fatigue fracture modes. He
found a good correlation, not entirely consistent, with the fatigue rankings at 107 cycles, see
table 3. From this correlation he proposed that slip band (Stage I) fatigue is a slower cracking
process than tensile mode (Stage II) fatigue, since the plastic deformation is dispersed and not
concentrated at a single crack tip. Consequently, conditions resulting in slip band fatigue
should lead to better fatigue resistance [2].
Table 3: Fatigue fracture modes and rankings at 107 cycles: Stage I = slip band cracking
Al-2.35Cu-1.3Mg
Al-4.18Cu-0.69Mg : experimental UA
mainly Stage I > PA
small Stage I >OA
little/no Stage I
Al-2.4Cu-1.51Mg : commercial PA
small Stage II >UA
mainly Stage II >OA
little/no Stage I
Al-4.0Cu-0.72Mg : commercial UA
mainly Stage I = PA
mainly Stage I >OA
little/no Stage I
However, Finney’s data extend beyond 107 cycles, where one would expect the fatigue lives
and strengths to be dominated by crack initiation and possibly a small amount of microcrack
growth [9, 10]. Other work, to be discussed in section 5 of this paper, showed that inclusions
and dispersoids in commercial alloys can affect crack initiation, and it was noted above that
Finney’s commercial alloys contained iron, nickel, silicon and manganese, which are
inclusion- and dispersoid-forming elements.
Finally, Forsyth [11] commented briefly on the UA > PA > OA ranking of one of Finney’s
experimental alloys.
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands R.J.H. Wanhill
He suggested – somewhat indirectly – that the superior fatigue strength of the UA material
could be due to S' precipitation on dislocations, resulting in dislocation pinning.
4 Comparison of the data of Lumley et al. [1] and
Finney [2]
Figure 4 compares the UA and PA experimental alloy data points from Lumley et al. [1] with
the appropriate fatigue life curves from Finney [2]:
Figure 4: Comparison of the fatigue properties of the UA and PA experimental alloys of Lumley et al. [1] and
Finney [2]
(1) Beyond 106 cycles Lumley’s UA data points agree well with Finney’s UA curve for the
Al-2.35Cu-1.3Mg alloy.
(2) Over the whole range of fatigue lives, Lumley’s PA data agree with the PA curve for both
of Finney’s alloys.
The agreements shown in figure 4 are in the first instance remarkable, since the alloys had
relatively large differences in copper contents, and the alloy of Lumley et al. contained silver,
which is a potent age-hardening element when added to AlCuMg alloys [12]. On the other
hand, the solid solution content of copper in the UA alloys may have been similar [13]. In any
event, it seems that the high-cycle fatigue properties were not strongly determined by
differences between the alloys in the amount of precipitation hardening up to peak strength.
5 Fatigue crack initiation in aerospace aluminium alloy
specimens
5.1 External or internal initiation
Fatigue cracks in high-strength aerospace aluminium alloy specimens nearly always initiate at
external surfaces. This is the case even in gigacycle fatigue, where internal crack initiation is
otherwise the rule for high strength steels and titanium and nickel alloys [14]. One practical
exception is when the aluminium alloy surfaces have been shot-peened [15], see subsection
5.4.
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5.2 Large inclusions
Over the last 50 years there have been numerous investigations of fatigue crack initiation in
high-strength aluminium alloys, e.g. Refs. [2, 9, 11, 16-24]. In some of the earlier work there
was an understandable tendency to focus on crack initiation and development along slip bands
[2, 9], and it was not recognised that fatigue cracks could initiate at large intermetallic
particles (inclusions) unless they were already cracked [9].
However, further studies showed that fatigue cracks nucleate at both cracked and uncracked
inclusions, and that these are the predominant sites of fatigue crack initiation in commercial
alloys [16-18, 20-24].
Figure 5 gives examples of fatigue crack initiation at cracked and uncracked inclusions in two
commercial alloy specimens, tested under widely different conditions (unnotched versus
notched, constant amplitude loading versus flight simulation loading). Note that figure 5a
refers to the same alloy specification, D.T.D. 5014, of one of the commercial alloys studied
by Finney [2]. The inclusions in figure 5a are FeNiAl9, and are characteristic of this alloy,
though they need not always be cracked [11].
Figure 5: Fatigue crack initiation at (a) cracked inclusions in a smooth specimen of D.T.D. 5014 tested
presumably under constant amplitude loading, and (b) uncracked inclusions in a notched (Kt = 3.2) specimen of
2024-T3 tested under gust spectrum loading [11, 25]
The sizes of the inclusions in figure 5 are typical. FeNiAl9 is specific to iron- and nickel-
containing alloys, but there are several kinds of large inclusions that commonly occur in
commercial alloys, including Al7Cu2 (Fe,Cr), (Fe,Mn)Al6, Mg2Si and Al2CuMg [23, 24, 26,
27]. Recent quantification of the sizes of inclusions associated with fatigue crack initiation
shows that they often reach depths of 10 – 20 μm into the specimen surface [23], but their
overall dimensions can be much larger [23], see figure 6.
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Figure 6: Comparison between fatigue-initiating inclusions and those obtained from metallographic sectioning of
the short-transverse planes of (a) thin and (b) thick 2024-T3 [23]. The fatigue-initiating inclusions are indicated
by specimen numbers. The inclusions are arranged from left to right and from top to bottom in descending order
of size (area)
Although inclusions are evidently associated with fatigue crack initiation, and there is general
agreement that the cracks initially follow the inclusion/matrix interfaces [11, 16-18, 20-23,
25], opinions differ as to whether interfacial debonding is involved [18, 20-23]. Also, it
appears that some cracks may start from processing voids between closely-spaced inclusions
[21].
Figure 7 suggests that interfacial debonding could well be involved in fatigue crack initiation.
Several inclusion/matrix interfaces in an inclusion cluster have debonded during (R = 0.7)
near-threshold fatigue crack growth [28], where there is a minimum of possible crack closure
and cyclic plasticity to disturb the appearance of debonding. Note the magnification: optical
and low-magnification SEM microscopy, as used in Refs. [16, 18, 20-23], will not resolve
such details, if present, and TEM of replicas [16-18] appears generally unsuitable. The
interfacial debonding is most likely due to dislocation pile-ups at the inclusions.
Figure 7: Inclusion/matrix debonding during near-threshold fatigue crack growth in OA 7075-T7351 plate [28].
The crack front marking are due to low-frequency growth measurements every 10,000 cycles
Owing to the importance of inclusions in nucleating fatigue cracks, Grosskreutz and Shaw
investigated the effect of reduced inclusion content on high-cycle reversed-stress notched
fatigue of the AlCuMg alloy 2024 [17]. An experimental alloy, X2024-T4, was obtained with
more than ten times fewer inclusions larger than 10 μm.
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Figure 8 shows that a relative absence of large inclusions has a beneficial effect on fatigue
compared with commercial 2024 alloys, but the improvement is not large.
Figure 8: Effect of reduced inclusion content on notched fatigue of experimental (X) and commercial UA 2024
alloys [17]
5.3 Dispersoids and small inclusions
There are some data on the effects of dispersoids, about 1μm in size, and small inclusions in
the 0.1 – 0.2 μm size range, on high-cycle fatigue of high-strength aluminium alloys [11, 19,
29]. Unlike large inclusions, the presence of these smaller particles can have a beneficial
influence on fatigue. This is shown in figure 9, which may be compared with figure 8.
Figure 9: Effects of (a) adding Mn- and Ni-containing dispersoids to a ternary PA A1MgZn alloy [11, 29] and
(b) reducing the small inclusion content of UA 2024 [19] on the unnotched fatigue properties
The effects of dispersoids in figure 9a may be due partly to preventing recrystallization and
retention of a fine grain size [11, 29], which would have prevented long dislocation pile-ups
along fatigue-induced slip bands, thereby reducing the stress concentrations leading to crack
nucleation. However, the dispersoids would also have helped to prevent cyclic slip
concentration in narrow bands, as did the small inclusions in the commercial 2024 alloy of
figure 9b [19].
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5.4 Initiation in air and in vacuo
Lumley et al. [1, 3] observed that the fatigue cracks in their specimens initiated internally.
Since internal cracks initiate in vacuo [15, 30], it is appropriate to compare the effects of
vacuum and air environments on the fatigue of aluminium alloys. Figure 10 shows subsurface
fatigue crack initiation in a shot-peened specimen [15]. The faceted appearance of the
initiation site is typical for vacuum fatigue [31], and is due to extensive slip band (Stage I)
cracking [15].
Figure 10: Internal high-cycle fatigue crack initiation in a shot-peened unnotched specimen from a 7070-T652
(PA) forging, tested under axial loading [15]
Several investigations have shown that the high-cycle unnotched fatigue lives of aluminium
alloys can be longer in vacuo than in air. Some results show life differences for up to 107
108 cycles [31, 32, 35, 36], e.g. figures 11a and 11b. Other data indicate little or no effect at
long lives [32-34, 37], see figures 11c and 11d.
Figure 11: Effects of vacuum fatigue, as compared to fatigue in air, for several aluminium alloys: D.T.D. 5070A
is A1-2.5Cu-1.5Fe-1.2Ni-0.25Si sheet in the PA condition and clad with A1-1.0Zn on the sheet faces; 1100 A1
is effectively unalloyed aluminium
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Fatigue cracks initiate at the same time [31, 36, 38] or later [33, 39] in vacuo than in air.
However, even in the latter case it appears that the major environmental effect is on crack
growth, which is much slower in vacuo than in air, e.g. Refs. [31, 40-42].
6 Fatigue crack initiation in aerospace aluminium alloy
components and structures
Probably the most widespread, though not the most frequent, source of fatigue cracking in
aerospace aluminium alloy components and structures is corrosion, which has led to failures
in all types of aircraft and a variety of components [24, 43-49]. Corrosion pitting causing
fatigue failures of propellers and rotor blades is particularly dangerous [45, 49]. Only very
small defects can be tolerated before rapid high-cycle fatigue initiates, e.g. figure 12, and the
aircraft usually becomes uncontrollable.
Figure 12: Fatigue crack initiation from a corrosion pit on a 2014-T651 (PA) forged propeller blade [44]
Another common initiator of fatigue is fretting [47], notably at faying surfaces and in fastener
and pin-loaded holes [50]. Fretting fatigue cracks initiate very early in the fatigue life [51-53].
For example, figure 13 shows a fretting-initiated fatigue crack in a fastener hole from a
window frame of a transport aircraft full-scale test article [54]. This crack had initiated and
grown during 45,402 simulated flights. Fractographic analysis enabled tracing crack growth
back to 0.15 mm from the origin, and back-extrapolation to the origin indicated that fatigue
began effectively immediately, as if there were an initial defect about 0.06 mm in size.
Figure 13: Fatigue crack initiation owing to fretting in the bore of a fastener hole in a 7175-T73 (OA) forged
window frame [54]. The arrows point to (a) the fatigue origin and (b) a detail showing the fretting scar
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Other surface-related damages that act as fatigue initiation sites include manufacturing defects
in fastener holes [55, 56], etch pits from anodising and other surface treatments [24, 46, 55,
57, 58], and fatigue cracking in cladding layers [23, 58-61]. Crack initiation from large
inclusions, clusters of inclusions, porosity and other material discontinuities has also been
observed [46, 55, 62].
Besides all these sources of fatigue, crack initiation can also occur owing to the stress
concentrations at fastener-filled and open holes; double stress concentrations, e.g. knife-edges
at fastener holes; fillet and blend radii; more or less abrupt changes in section thickness;
cutouts; and higher than anticipated service stresses owing, for example, to secondary bending
and excessive load transfer [63-70].
Some “milestone case history” examples of aircraft accidents involving fatigue in aluminium
alloys are shown in figures 14-16 [71]. These, and others, notably the General Dynamics F-
111 crash in 1969 [69, 71], have led to adoption of the Damage Tolerance design philosophy
for military and transport fixed-wing aircraft, and Aviation Regulations requiring its use. This
is also likely to be the case for helicopters [72].
Figure 14: Details of the probable failure origin in the pressure cabin of DeHavilland comet G-ALYP [64,71],
which was lost in January 1954. Figures 14a and 14b show the frame cutouts to accommodate the stringers, and
the resulting stress concentration at the first fastener joining the frame and stringer flanges to the skin (and an
outside doubler). Figure 14c show that the first fastener was a countersunk bolt, creating knife-edges in both the
skin and outside doubler. There was thus a double stress-concentration. This combined with local unanticipated
secondary bending to cause early fatigue failure [64,71]
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Figure 15: Failure origin in the right-hand horizontal stabilizer of a Boeing 707-321C freighter [68, 71]. which
was lost in May 1977. Fatigue initiated at a fastener hole in the upper chord of the rear spar attachment to the
skin, owing to higher loads than those anticipated in the desingn
Figure 16: Structural aspects of the Boeing 737-200 accident in April 1988 (the famous Aloha Airlines case) [65,
71]. Defective cold bonding allowed moisture to enter the skin lap area. This led to corrosion-induced
disbonding, both in the cold-bonded skin splice and the associated hot-bonded tear straps. The loss of skin splice
integrity caused the pressure cabin loads to be transferred through the rivets. These had countersunk heads
causing knife-edges in the upper skin, and the knife-edges caused stress-induced Multiple Site fatigue Damage
(MDS) in the upper skin along the upper rivet row of the skin lap area [65, 71]
There are some essential points to be made about the Damage Tolerance philosophy in the
context of the present paper:
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(1) The concept of an initial crack-free fatigue life is abandoned. Instead it is assumed that
defects (cracks and flaws) are already present at critical locations in new aerospace
components and structures, and that these assumed defects must be treated as cracks that
are immediately capable of growing by fatigue, albeit slowly.
(2) The assumed defects are all surface-connected, for example table 4. When designing for
safety the assumed defects are all larger than 0.5 mm. When designing for durability the
assumed defects are generally larger than 0.1 mm [73, 74] and unlikely to be smaller than
0.03 mm [25].
(3) The assumed defects will be mostly at fastener holes. A large majority of fatigue cracks
in aluminium alloy aircraft structures initiate at fastener holes in joints [70].
Table 4: Example of Damage Tolerance safety requirements for assumed initial damage [73]
From the above information and remarks we may conclude that both actual and assumed
initial damage and fatigue cracks in aerospace aluminium alloy components and structures are
surface-related. The assumed initial defects are also much larger than actual ones, reflecting
the necessary conservatism when designing safety-critical structural elements.
7 Discussion
7.1 The self-healing concept
The possible self-healing of fatigue cracks in UA aluminium alloys is subject to some
constraints. Its occurrence should be limited to internal cracks, since cracks initiating at an
external surface in an air environment will be contaminated by adsorption of water vapour
and oxygen onto the crack surfaces [41]. This would likely prevent self-healing, and certainly
promotes crack growth compared to fatigue in vacuo [31, 40-42]. There is also the question of
crack size. If internal cracks were to nucleate at large inclusions, e.g. from dislocation pile-
ups at the inclusion/matrix interfaces, then they would be expected to rapidly assume
dimensions larger than 10 μm.
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This makes the idea of self-healing less feasible. On the other hand, internal cracks can
initiate from slip band cracking [15], with no lower-bound restriction on the initial crack size.
However, there appears to be another problem. Slip band crack initiation and self-healing
would both have to occur via the movement of dislocations, even if the self-healing
dislocations do not actually cross the crack, as sketched in figure 2, but only deliver solute by
pipe diffusion to the crack vicinity [75]. Lumley et al. [1] obtained evidence to suggest that
the dislocations in their UA alloy became saturated with free solute copper atoms relatively
early in the fatigue process. Since high-cycle fatigue crack initiation usually occurs late in
life, there does not seem to be a reason why there should be two types of dislocation
(unsaturated and saturated) that compete in initiating and self-healing events.
Be that as it may, there will be an upper-bound to the size of slip band crack that can be self-
healed. This upper-bound could be when the crack becomes capable of generating
dislocations from its tips. An analysis to this effect has been given by Tanaka and Akiniwa
[76]. They considered a Stage I slip band crack in aluminium, and calculated that it can emit
dislocations when the total length reaches 1.4 μm. This seems entirely reasonable for the
upper-bound size of a self-healing crack.
7.2 High-cycle fatigue of UA and PA experimental aluminium alloy
specimens
At the beginning of section 3 of this paper it was stated that the only consistency from pre-
1969 studies of ageing effects on aluminium alloy fatigue is that overaged (OA) materials
never had the best fatigue properties, see table 1. Finney’s 1969 study [2] reinforces this
conclusion, as may be seen from figure 3.
However, Finney’s study and that of Lumley et al. [1] show a similar result, namely that their
experimental alloys had better underaged (UA) high-cycle fatigue properties than the
equivalent peak aged (PA) alloys, see figures 1, 3 and 4. The significance of this result in the
present context is as follows:
(1) Because these alloys were experimental, they were relatively large-grained owing to
recrystallization during solution treatment [2] and grain growth [77]. They would also
have been relatively free from large inclusions, and Finney’s alloys would have had few
dispersoids in the 0.1 – 0.2 μm size range. In the UA condition such relatively clean
microstructures should favour the development of long slip bands during fatigue, and
extensive and long Stage I fatigue cracks, as found for Finney’s alloys [2], see table 3
also.
(2) While Lumley et al. reported internal fatigue crack initiation for their reversed-stress UA
and PA specimens [1, 3], Finney reported surface initiation for rotating cantilever
specimens [2]. (This would be expected anyway, since the maximum fatigue stresses for
rotating cantilever specimens are at the surface.) From (1) and (2) one may conclude that
the better fatigue properties of the UA experimental alloys cannot depend only on the
delay of crack initiation by self-healing of small internal fatigue cracks. There are other
possible mechanisms, already mentioned in sections 2 and 3.
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These are (a) reductions in dislocation mobility owing to their becoming saturated by free
solute [1]; (b) dislocation pinning by precipitates [11]; and (c) localised matrix hardening due
to dynamic precipitation [1]. Note that these mechanisms could apply to all the experimental
alloy results, i.e. self-healing is not required.
7.3 High-cycle fatigue of UA and PA commercial aluminium alloy specimens
Depending on the alloy, and to some extent on the fatigue lifetimes between 106 and 108
cycles, see figure 3, the underaged (UA) commercial alloys had better [6, 7], similar [2] or
worse [2, 4] high-cycle fatigue properties than their peak aged (PA) equivalents.
This lack of consistency suggests that the high-cycle fatigue properties of commercial alloys
can be determined more by differences in grain size and inclusion and dispersoid contents
than by the ageing condition up to peak strength. Also, the surveys in subsections 5.1 and 5.2
indicate that the predominant sites of fatigue crack initiation are large surface-connected
inclusions. All in all, there seems to be little scope for the self-healing of small internal cracks
to influence the basic fatigue properties of commercial alloys.
7.4 Aerospace aluminium alloy components and structures
The survey in section 6 of this paper shows that actual initial damage and fatigue cracks in
aerospace aluminium alloy components and structures are surface-related, and a large
majority will be at fastener holes in joints. Inevitably, this is also the case for the assumed
defects when designing according to Damage Tolerance requirements. The minimum size of
the assumed initial defects is at least 0.03 mm, and generally more than 0.1 mm.
The actual initial damages and fatigue cracks and Damage Tolerance requirements show that
the concept of self-healing of small internal fatigue cracks is inappropriate to aerospace
aluminium alloy components and structures.
8 Conclusions
The possible self-healing of fatigue cracks in underaged (UA) aluminium alloys is
(1) Limited to internal slip band cracks probably shorter than about 1.4 μm.
2) Not capable of explaining the similar high cycle fatigue behaviour and/or rankings of UA
and peak aged (PA) experimental alloys tested by Lumley et al. [1] and Finney [2]. Nor is
it required for explaining why any UA experimental alloys can have better high-cycle
fatigue properties than their PA equivalents.
(3) Unlikely to influence the basic fatigue properties of commercial alloys.
(4) Inapplicable in practice to aerospace aluminium alloy components and structures.
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands R.J.H. Wanhill
16 © Springer 2007
ACKNOWLEDGEMENTS
My thanks to Roger Lumley, Ali Merati, Co Dalemans, Tim Hattenberg and the NLR Library Staff for their
assistance. Thanks also, in retrospect, to Jack Finney.
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... Aircraft grade aluminum alloys are principal structural material for aircraft parts because they combine high strength and low weight [1,2]. Their mechanical characteristics such as high strength, toughness, and excellent anti-corrosive properties stem from their microstructures which result in a combination of composition, deformation, natural and artificial aging treatments [2,3]. The 7xxx and 2xxx series aluminum alloys are hardenable by precipitation which assigning higher mechanical properties. ...
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The corrosion fatigue behavior of three metallurgical conditions of α-brass (as- received, cold- rolled and annealed) in 1 M NaNO2 solution under free corrosion potential and under a constant applied anodic po- tential of 300 mVNHE (Normal Hydrogen Electrode) was investigated by applying the reverse bending technique at 70 cycle /min. The fatigue life and fatigue strength of α-brass in 1 M NaNO2 solution under free corro- sion is considerably shorter than that obtained in air, while under applied anodic potential value 300 mVNHE a further reduction in fatigue life was obtained. The longest fatigue life is that corresponding to the cold- rolled followed by the fatigue life of the as- received and finally the fatigue life to the annealed specimens. The same order was also obtained in sodi- um nitrite solution at the free corrosion condition as well as under the applied anodic potential. The fracture mode of the fatigued specimens changed from transgranular in air to mixed inter-and transgranular in nitrite solution for the as- received and cold- rolled. Annealed specimens show a mixed mode in both air and nitrite solution. The conjoint action of film rupture and adsorption models is the operating mechanism of corrosion fatigue of brass in nitrite solution.
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