Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
ADVANCES IN SELF HEALING OF METALS
R.N. Lumley1* & I.J. Polmear2
1. CSIRO Manufacturing and Materials Technology, Private Bag 33, Clayton South MDC,
Victoria 3169, Australia.
2. Professor Emeritus, School of Physics & Materials Engineering, Monash University,
Melbourne, Vic. 3168, Australia
Tel: +613 95452894
Fax: +613 95441128
e-mail: Roger.Lumley@csiro.au
Press-and-sinter powder metallurgy is an example of an industrial processing technique that relies entirely on
natural “healing” behavior to bond particles and provide mechanical strength. During sintering of the porous
body at close to the melting point of the alloy, mass transfer occurs producing liquid phase or solid state bonding
across material interfaces. With some aluminium alloys, it has now been shown that significant densification
can also occur at lower temperatures by inducing precipitation in pores through controlled cooling from the
sintering temperatures.
Recently it has been shown that a conventionally produced, wrought Al-Cu-Mg-Ag alloy, underaged to ~85% of
its peak strength, outperforms the fully hardened alloy in accelerated creep tests despite its lower tensile
properties. This behaviour has been attributed to a form of self healing whereby retained solute facilitates
dynamic precipitation throughout the matrix which inhibits dislocation motion during creep.
A comparative study has also been made in which alloy specimens in either the underaged or fully hardened
conditions were subjected to fatigue loading. Dynamic precipitation again occurred in the underaged alloy and
fatigue properties were substantially improved. By analogy with the mechanisms of precipitation induced
densification of sintered aluminium alloys, it is proposed that dynamic precipitation in the underaged condition
may facilitate the healing of incipient fatigue cracks as they develop.
Keywords: Aluminium, Creep, Fatigue, Precipitation
1
Introduction
Biological structural materials such as in-vivo bone repair by a process that, in essence,
involves material being dissolved from where it is not required and re-precipitated to where it
is needed, in response to mechanical stimuli and damage. This process was first described in
the late 19th century, and is now known as the Wolff-Roux law [1,2].
Metals and alloys deteriorate in service. Surface damage due to corrosion and wear, and the
development of external and internal cracks leads to a gradual loss of properties that may
eventually cause failure of a component or structure. Recent studies have suggested that these
inanimate materials may also benefit from healing processes to repair defects or
microstructural damage which develops during manufacture, processing, or in service.
Press and sinter powder metallurgy is an example of an industrial processing technique in
which pores and other defects are effectively healed during their manufacture.
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© Springer 2007
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
When a powder compact is sintered at a temperature of at least 0.5Tm, or more commonly at
higher temperatures such as 0.7- 0.9Tm, mass transfer of metal occurs producing either solid
state or liquid phase bonding across powder interfaces. Particles bond together by one or
more mechanisms, and the mechanical strength of the compact increases. Sintering of
powdered materials aims to achieve an increase in density through the removal of porosity.
This is not always the case however, and transient liquid phase sintered materials in
particular, often display particle bonding and simultaneous density reduction via the
Kirkendall effect where unfavorable diffusivity conditions exist.
Recent work on some aluminium powder alloys has shown that porosity may be substantially
reduced if a compact is slowly cooled rather than quenched after being exposed to a high
temperature [3]. This situation has been shown to apply to alloys in which there is a high
propensity for precipitation to occur, and is illustrated in Table 1 [4]. For example, slow
cooling a powder compact of the alloy Al-8Zn-2.5Mg-1Cu from 620°C can result in a
reduction of 8% in porosity. This densification is attributed to removal of solute from the
aluminium solid solution combined with heterogeneous precipitation of the η phase (MgZn2),
particularly on pore surfaces. Examples of this phenomenon are shown in Figure 1.
Table 1: Volume Changes In Different Aluminium Alloy Systems By The Process Of Precipitation Induced
Densification
Alloy
Density after Density after Densification
Solute
Predicted
sintering
slow cooling
(+/-)
content
Precipitate
(%Theoretical) (%Theoretical)
(atomic %)
species
89
88
3.5
γ
Al8Zn
90.5
91
+
Al2.5Mg
+/-♦
88
90
Al2.5Mg
-1Cu
90
93
+
Al8Zn2.5Mg
90
98
+
Al8Zn2.5Mg
-1Cu
89
93
+
Al4.5Cu1.6Mg
♦Note: Test samples showed a mix of expansion and shrinkage
Propensity for
precipitation
Very Low
β
Al3Mg2
S
Al2CuMg
Very Low to
Nil
moderate
6.2
η
MgZn2
High
6.6
η
MgZn2
Very high
3.9
S
Al2CuMg
Very High
2.9
3.4
Evidence for another type of healing has been obtained in studies of conventionally produced,
wrought aluminium alloys that also undergo precipitation. It has been found that such alloys
may show superior creep and fatigue properties if they are heat treated to the underaged
condition rather than being fully age hardened (T6 temper) [5-10]. As described below, this
behavior has recently been attributed to retained solute not yet committed to form precipitates,
which is available to undergo dynamic precipitation throughout the matrix during creep or
fatigue loading.
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© Springer 2007
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
Figure 1: Backscattered SEM micrographs of the sintered Al-8Zn-2.5Mg-1Cu alloy revealing the process of
precipitation induced densification. (a) shows precipitation on the surface of a pore, and (b) shows a higher
magnification image of precipitation into the tip of an elongated pore. The η phase precipitates are shown as the
white phase present in the micrographs
2
Creep studies
Earlier studies of the creep performance of a high purity, experimental alloy Al-4Cu-0.3Mg0.4Ag revealed that the secondary creep rate at 125°C was reduced by approximately 65% if
the alloy was first tested in the underaged (UA) condition (1h at 170°C) rather than the fully
hardened, T6 temper (12h at 170°C). This result was unexpected because the underaged alloy
had a lower 0.2% proof stress (313 MPa) when compared with the T6 value (363 MPa).
Effects of underageing were investigated further using extrusions prepared from the
experimental composition Al-5.6Cu-0.45Mg-0.45Ag-0.3Mn-0.18Zr and comparative creep
curves for tests conducted at 150°C and a stress of 300MPa are shown in Figure 2(a) [7]. The
secondary creep rate for the underaged (UA) condition (2h at 185°C) was found to be 3.5 ×
10-10/s which was close to one third of the creep rate for the fully hardened (T6) alloy (1.12 ×
10-9/s). Subsequently, prolonged tests revealed that the UA alloy showed zero secondary
creep after 20000h at 130°C and a stress of 200 MPa [11]. Improved creep performance in
accelerated tests was also recorded for the commercial alloy 2024 (Al-4.4Cu-1.5Mg-0.6Mn)
in the UA condition (Fig.2b). In this case, tertiary creep was delayed and the time to failure
was increased from ~260h in the T6 condition to ~480h by underageing prior to testing [7].
Strain to failure was also increased. Results from these investigations suggested that the
beneficial effects of underageing on creep properties may be applied more generally to aged
aluminium alloys.
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© Springer 2007
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
Figure 2: Creep curves for the T6 and UA condition in (a) Al-Cu-Mg-Ag alloy, and (b) 2024 alloy
Reasons for the enhanced creep behavior of the underaged alloys were sought by making a
detailed comparison of the microstructures of the Al-Cu-Mg-Ag alloy in the UA and T6
conditions,
(a) after the initial underageing treatment,
(b) at the beginning of creep at 150°C after loading to a stress of 300 MPa, equilibrating
for 0.5h, then cooling to room temperature, and
(c) after accelerated creep testing for 500h at 150°C and a stress of 300 MPa.
Specimens for electron microscopy were prepared from the gauge length of the creep
specimens parallel to the longitudinal direction of the extruded bars. These were produced as
thin, 3mm discs by an electrical discharge machine to avoid any mechanical deformation.
The discs were then thinned by standard electropolishing techniques and examined in a Jeol
2000EX microscope operating at 200kV.
Comparative microstructures of the UA and T6 conditions after creep testing 500h at 150°C
are shown in Figure 3. Observations suggested that the beneficial effect of underageing is due
to one or more of the following factors:
1) The ability of residual solute in the matrix of the UA alloy to form solute atmospheres
around dislocations that impede or inhibit their motion;
2) Retention of plates of the θ′ phase (on the {001}α planes) within the matrix during
creep rather than its dissolution, as happens in the T6 alloy;
3) Additional dynamic precipitation of θ′ that occurs on mobile dislocations throughout
the matrix;
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© Springer 2007
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
4) Retention of a finer dispersion of plates of the Ω precipitate within the matrix;
5) Retention of a fine subgrain structure with particles of θ′ and Ω densely dispersed
along subgrain boundaries.
Other more recent suggestions of effects that may be occurring during creep from the
underaged state have also been proposed, and include:
6) Modeling which predicts that a minimum in the dislocation velocity occurs when the
alloy is in the underaged condition due to a combination of solid solution
strengthening and precipitation hardening [12],
7) A release of vacancies into the aluminium matrix from the surface of Ω particles in the
underaged material that will alter precipitation occurring during creep [13].
Figure 3: Microstructures of the Al-Cu-Mg-Ag alloy in [101]α and [001]α orientations, following 500h creep at
150°C and 300 MPa. In the T6 material, the θ′ phase is removed from the matrix aluminium whereas for the UA
material, the θ′ phase is retained through being continually replenished by a process of dynamic precipitation
3
Fatigue studies
The nature and dispersion of precipitates can sometimes effect fatigue properties in a way that
is different from their effect on simple tensile properties. While there is a general trend of
increased fatigue resistance with increased tensile strength, that relationship is not always
followed [14-18]. Fatigue crack initiation is normally preceded by strain localization so that
deformation becomes progressively more concentrated in preferred (persistent) slip bands
[e.g. 15]. Dislocations that move back and forth within these bands shear precipitates and
repeated cutting may then cause them to re-dissolve [e.g. 16,17].
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© Springer 2007
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
This process causes softening in these bands leading to a further concentration of strain during
the process of fatigue thereby accelerating crack initiation. On the other hand, the presence of
larger non shearable precipitates, or fine intermetallic compounds, causes dislocations to be
dispersed more uniformly which may delay or inhibit crack initiation [14,18].
Early studies of the fatigue properties of aged Al-Cu alloys revealed that times to failure were
longest if the alloy was in the UA condition (Figure 4(a))[9]. Observations using scanning
electron microscopy and backscattered electron diffraction from grains on the fracture surface
showed that through a process of restricted slip, the fatigue crack was constrained to
propagate along {001}α planes (Figure 4b). It is also interesting to note that shear banding
occurred on the {111}α planes close to the crack only in the UA condition, although no
observations were made of precipitation in this region [9].
Figure 4: (a) shows the comparative crack length during fatigue testing of two alloys under a range of different
conditions from underaged to overaged. U,P and O correspond to underaged, peak aged or overaged tempers
respectively. Alloy 3 is Al-3Cu and alloy 4 is Al-4.5Cu. (b) shows propagation of a crack in the UA alloy
following closely to an {001}α matrix plane. Shear banding on {111}α planes is observed in association with the
propagating crack. T6 treated alloy did not show either of the features observed for (b)[9]
More recently, comparisons have been made between the fatigue performance of the extruded
Al-Cu-Mg-Ag alloy, in either the UA and fully hardened conditions, where the UA time
corresponded to the optimum underageing period for maximum creep resistance mentioned
earlier. A particular aim was to test possible correlations between creep and fatigue
performance in the UA condition.
The extruded bars were prepared in the same manner as for the creep study. UA specimens
were solution treated for 6h at 525°C, quenched into cold water, and aged 2h at 185°C.
Fatigue tests were carried out for the following conditions:
(a) Within 24h at room temperature following quenching from the ageing temperature (in
order to minimize secondary precipitation)
(b) Slow cooling from the ageing temperature and holding at 25°C for >1200h (designated
SUA)
(c) The so-called T6I4 treatment [19] in which specimens were quenched after
underageing 2h at 185°C, and then held >1200h before testing
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© Springer 2007
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
(d) The T6I6 treatment [19] which involved quenching after underageing 2h at 185°C,
holding for ≥ 340h at 25°C, then continuing ageing at 150°C until peak hardness was
recorded.
These conditions were compared to the standard T6 temper for this alloy, which involved
ageing for 10h at 185°C following solution treatment and quenching.
Fatigue tests were conducted in an axial loading machine with a stress value of R=-1 (i.e. zero
mean stress). Initial evaluations were made at a stress amplitude of ± 235 MPa in order to
examine a possible correlation between creep and fatigue performance. A minimum of four
tests for each condition were evaluated and fatigue samples were prepared to a highly
polished finish to minimize surface defects. Results are shown in Figure 5 and it is interesting
to note that the reciprocals of the secondary creep rates for tests at 150°C and a stress of 300
MPa do reveal similar trends to those observed for the fatigue results. In both fatigue and
creep testing, the UA and SUA tempers outperformed the T6, T6I4, and T6I6 tempers. The
major difference between the two underaged tempers and the other tempers resided in the fact
that the underaged tempers contained considerable residual solute, whereas the other three
tempers did not.
Figure 5: (a) the reciprocal of secondary creep rate for five tempers examined, and (b), mean fatigue lives for
fatigue testing at ±235 MPa
More detailed stress versus number of cycles (S/N) plots were then determined for the UA
and T6 conditions, in which multiple tests were conducted at differing stress levels leading to
failure within the approximate range of 105 to 107 cycles. As shown in Figure 6, the UA alloy
generally sustains longer lives over a range of stress levels. This is despite the fact that the
0.2% proof stress for the UA alloy (417 MPa) was less than that of the T6 alloy (470 MPa).
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© Springer 2007
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
Figure 6: Fatigue S-N curves for the T6 and UA material. 107 values are runouts
Transmission electron microcopy was also used to study the microstructural damage
accumulated during fatigue in the UA and T6 specimens. Similarities were observed between
the microstructures that developed during fatigue and creep testing. In the T6 condition, the
matrix became progressively depleted of θ′ precipitates as dislocations generated by the
fatigue process moved through the matrix. On the other hand, the presence of residual solute
in the matrix of the UA alloy facilitated dynamic precipitation of θ′ during fatigue which
appeared to reduce dislocation mobility and maintain a more uniform dispersion of
precipitates (Figure 7). Furthermore, dynamic precipitation during fatigue indicates that
mobile dislocations have become saturated with solute.
Figure 7: Microstructures of T6 and UA material having undergone 100,000 fatigue cycles at ±160 MPa, shown
in an [001]α orientation. The T6 material has had selected variants of the θ′ phase dissolved, whereas the UA
material is continually replenished by the action of dynamic precipitation (examples arrowed)
Pipe diffusion of solute atoms associated with dislocations is typically some 105 to 107 times
faster than vacancy diffusion in the bulk material [20], and estimates have suggested a similar
increase (of up to 106 times) is likely to apply to the diffusion of Cu in aluminium at ambient
temperature [21]. This means pipe diffusion rates are equivalent to vacancy diffusion rates
occurring at much higher temperatures.
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© Springer 2007
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
Such high solute mobility in association with mobile shear bands or dislocation structures
such as those shown in Figures 3 and 7 suggest that relatively significant quantities of solute
have the ability to be readily delivered to a crack site and contribute to closure of the crack by
a mechanism analogous to precipitation induced densification as shown in Figure 1.
In addition, it should be noted that the overall microstructure of the alloy is susceptible to
deterioration during creep and fatigue testing. In particular, dissolution of the θ′ phase in the
T6 condition leaves the alloy with inferior properties when compared with the UA condition
in which underlying damage is being continually repaired as this phase is replenished.
Finally, it also seems possible that self healing may be facilitated as a result of volume
changes associated with precipitation, as detailed by Hunsicker [22,23]. In alloys based on
the Al-Cu system, precipitation of the θ′ phase is known to cause a volume expansion. This
volume change is most pronounced when Mg content is relatively low, since more θ′ forms.
It therefore follows that concentrated dynamic precipitation of θ′ precipitates in bands will
also cause a localized heterogeneous expansion of the aluminium matrix. When this is
considered in light of the configuration of a crack with associated shear bands as shown in
Figure 4b, it also seems clear that the region surrounding the crack should in fact experience
compression, which may slow fatigue crack propagation.
Progressive dilatometric changes associated with ageing the Al-Cu-Mg-Ag alloy at 185°C are
shown in Figure 8 together with the thermal cycle occurring during heating to, and at, this
temperature. It is clear that dimensional changes occur in three stages. The first is a rapid
expansion that shows a peak when the alloy reaches 140°C during heating, which may
correspond to the initial formation of solute clusters prior to precipitation. A rapid contraction
then occurs that may be attributed to nucleation of the Ω phase on the {111}α planes. This
contraction then occurs more slowly up to an ageing time of 2-2.5h, as competitive
precipitation of Ω and θ′ occurs. At about 2-2.5h, Ω precipitation ends [7] and from that time
on, the alloy expands at a rate that gradually decreases as precipitation of θ′ continues.
Figure 8: Dimensional changes associated with ageing Al-Cu-Mg-Ag alloy at 185°C. Sample length 25mm. See
text for details. (Differential dilatometry courtesy W.Wright, CSIRO)
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© Springer 2007
Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
R.N. Lumley et al.
What is also apparent from Figure 8 is that the UA alloy has a much greater capacity to
undergo localized expansion, should dynamic precipitation take place. Conversely, it may
also be predicted that dissolution of the θ′ phase in the T6 alloy during fatigue stressing will
lead to a dimensional contraction in the vicinity of an advancing crack, which may lead to
faster growth rates and earlier failure.
4
Conclusions
Self healing in UA aluminium alloys may occur due to three proposed mechanisms:
1. Continual replenishment of strengthening phases within the microstructure by a
process of dynamic precipitation;
2. heterogenous precipitation into short crack sites in a manner analogous to precipitation
induced densification;
3. crack closure due to volume changes causing localized compression, associated with
heavy precipitation of θ′ in the plastic zone of a crack or damage related defect.
ACKNOWLEDGEMENTS
The authors would like to thank Drs John Griffiths, Allan Morton, Rob O’Donnell, Cameron Davidson, Wayne
Wright and Prof. Graham Schaffer for their contributions, and CSIRO for financial support of this work. The
authors would also like to thank Dr Russell Wanhill for helpful discussions.
REFERENCES
1.
2.
3.
4.
5.
6.
7.
8.
9.
10.
11.
12.
13.
14.
15.
16.
17.
18.
19.
20.
J. Wolff, Das Gesetz der Transformation der Knochen, Hirschwald, Berlin, (1892).
W. Roux, Arch. Anat. Physiol., Anat. Abt., 120 (1885).
Lumley, R.N., Schaffer, G.B. Scripta Mat., 55, p.207, (2006).
Lumley, R.N., unpublished research, (1996).
Arumulla, S.R & Polmear, I.J. Proc. 7th Int Conf. on Strength of Metals and Alloys, Montreal, Canada,
Pergamon Press, Oxford, 1, p.453, (1985).
Lumley, R.N., Morton, A.J and Polmear, I.J, Mat. Sci. Forum Vols. 331-337 p.1495-1500, (2000).
Lumley, R.N, Morton, A.J, Polmear, I.J., Acta Mat., 50, p.3597, (2002).
J.M. Finney, The Relationship Between Aged Structure And The Fatigue Behavior Of Aluminium
Alloys, Part 2, Experimental Results Of Aluminium-Copper-Magnesium Alloys, Aeronautical Research
Laboratories, Structures and Material Report 325, Melbourne Australia, July (1969).
Garrett, G.G & Knott, J.F., Acta Met., 23, p.841, (1975.)
Lumley, R.N, O’Donnell, R.G., Polmear, I.J., Griffiths, J.R., Mat. Sci. Forum, Vol.29 p.256 (2005).
R.N. Lumley and I.J. Polmear, Scripta Mat., 50 #9, p.1227, (2004).
C. R. Hutchinson, P. Cornall and M. Gouné, Mat. Sci. Forum, Vol. 519-521, p.1029, (2006).
Somoza, A.., Macchi , C.E.., Lumley, R.N.., Polmear, I.J.., Dupasquier A., & Ferragut, R.. 14th ICPA International Conference on Positron Annihilation, Hamilton, Ontario, Canada., (2006), To be published
in Physica Status Solidi, (2007).
K. Boyapati and I.J. Polmear, Fatigue of Eng. Mater. And Struct., 2, p.23. (1979)
P.J.E. Forsyth and C.A. Stubbington, J. Inst. Metals, 83, p.395, (1954-55).
I.J. Polmear and I.F. Bainbridge, Phil. Mag., 4, p.1253, (1959).
C.A. Stubbington, Acta Met., 12, p.931, (1964).
G. Lubgering, H. Doker and D. Munz, Proc. 3rd Int. Conf. of Strength of Metals and Alloys, Cambridge,
England, Vol.1, p.427, (1973).
Lumley, R.N., Polmear, I.J., Morton, A.J. Mat. Sci Tech., 19, 11 p.1483, (2003).
Landolt-Börnstein Numerical Data and Functional Relationships in Science and Technology, 3, 26,
Diffusion in Metals and Alloys, Ed. H. Mehrer, Springer-Verlag, p.195-196, (1991).
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Proceedings of the First International Conference on Self Healing Materials
18-20 April 2007, Noordwijk aan Zee, The Netherlands
21.
22.
23.
R.N. Lumley et al.
Jannot, E.J., Private communication, Institut für Metallkunde und Metallphysik, RWTH Aachen, 52074
Aachen, Germany, (2006).
Hunsicker, H.Y. in “Aluminum, Properties, Physical Metallurgy and Phase Diagrams”, Vol.1, K.R. Van
Horn ed., p.158-159, (1967).
Hunsicker, H.Y., Met. Trans. A, 11A, 759, 1980.
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