THE TEMPERATURE DEPENDENCE OF THE HIGH CYCLE FATIGUE PROPERTIES OF AN Al-Li-Cu-Mg ALLOY

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THE TEMPERATURE DEPENDENCE OF THE HIGH CYCLE FATIGUE PROPERTIES OF AN

Al-Li-Cu-Mg ALLOY

P. Bischler, J. Martin

To cite this version:

P. Bischler, J. Martin. THE TEMPERATURE DEPENDENCE OF THE HIGH CYCLE FATIGUE

PROPERTIES OF AN Al-Li-Cu-Mg ALLOY. Journal de Physique Colloques, 1987, 48 (C3), pp.C3-

761-C3-767. �10.1051/jphyscol:1987389�. �jpa-00226621�

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T H E TEMPERATURE DEPENDENCE OF T H E HIGH CYCLE FATIGUE PROPERTIES OF AN Al-Li-Cu-Mg ALLOY

P.J.E. B I S C H L E R ( ~ ) and J.W. M A R T I N

Oxford University Department of Metallurgy and Science of Materials, Parks Road, GB-Oxford, OX1 3PH, Great-Britain

Abstract

It is known that a post-solution treatment stretch gives a finer, more uniformly distributed S-phase when ageing Al-Li-Cu-Mg alloys. Fatigue S - N curves have been determined, and tensile tests have been carried out on stretched (P) and unstretched (N) samples of 8090 alloy in the peak-aged condition at temperatures of 288K, 358K and 425K. A fractographic comparison has been made of the fatigued samples in the two conditions.

It is found that magnitude of the Endurance Ratio (of Endurance Limit to UTS) is relatively high, due to the homogenisation of slip by the S-phase, and is higher in the P than in the N material. A decline in Endurance Ratio with increase in temperature is observed in P material, but not in the N material. The results are discussed in terms of the dislocation interactions in the two microstructures, and is supported by previously observed differences in dynamic recovery characteristics between N and P material.

Introduction

The alloy studied in the present investigation is an Al-Li-Cu-Mg-Zr alloy designated 8090, and it may be regarded as a replacement for 2014 T651, but with a stiffness increase and density reduction of at least 10 percent. The main strengthening phase in this system is 6' (A13 Li), which forms a homogeneously distributed precipitate possessing the ordered L12 structure. This phase encourages planar, inhomogeneous slip, which may lead to premature failure.

Cu and Mg have been added to homogenise the slip by the formation of S-phase precipitates (A12 CuMg), which is observed to nucleate preferentially on stray dislocations and upon subgrain boundaries. A more uniform distribution of S-phase is produced by stretching the solution-treated material prior to artificial ageing.

This process produces uniform dislocation networks throughout the structure which act as S-phase nucleation sites.

The present authors have investigated the variation in elevated temperature tensile properties of this alloy (I), by comparing the behaviour of the stretched (P) and unstretched (N) material in the peak-aged condition at temperatures between 288K and 550K. In particular it was noted that the work-hardening rate begins to fall at a lower temperature in P than in N material.

In view of this greater resistance to dynamic recovery exhibited by material in the N condition compared with that in the P condition, it might be expected that other elevated temperature properties would also show differences. The object of the present investigation has been to compare the high-cycle fatigue behaviour of the peak-aged alloy in both N and P conditions, by determining the form of the S - N curves at three temperatures. The temperatures selected were room temperature, 358K and 425K which is a range in which there is known to be marked differences in dynamic recovery characteristics as established from work-hardening rate measurements (1).

("present Address:

C.E.G.B.

, B .N.L. , Berkeley, Glos.

Great-Britain

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1987389

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C3-762 JOURNAL DE PHYSIQUE

Experimental Procedure Material

The composition of the alloy is given in Table 1. The material was in the form of rolled plate of 25mm thickness, supplied in the stretched (2.5%) condition.

Table 1: Composition of 8090 alloy Element: Li Cu Mg Zr Fe Si Na A1 Wt. % : 2.41 1.16 0.61 0.11 0.14 0.10 0.0015 bal.

Testing procedure

Cylindrical tensile specimens of gauge length 18 mm and diameter 5 mm, and plane bending fatigue testpieces of thickness 3mm were machined from the centre-line of the plate to avoid any through-thickness inhomogeneity (2). All specimens were made with the longitudinal direction parallel with the longitudinal axis of the testpiece. Half the specimens were given a further 45 min at 803K to remove the effect of stretching, followed by a cold water quench. Both sets were then aged at 463K for 20 hours to produce peak strength.

Tensile tests were conducted in air at 288K, 358K and 425K at a strain rate of 1 x s-I to BS 6388 pt. 1. Specimens were maintained at test temperature for 0.5 hour prior to commencment of testing. S - N curves were also determined at 288K, 358K and 425K in air using an Avery Plane Bending Machine operating at a frequency of 1420 c.p.m., with R

=

-1. The elevated temperature tests were effected by placing a resistance heater round the waist of the specimen in each case.

Results (a) Microstructures

The microstructures of the peak-aged N and P material have been studied by TEM, and reported in detail elsewhere (1). In both conditions there is a homogeneous distribution present of the

6 '

strengthening phase, of average particle size 30nm.

A distinct difference in S-phase distribution is evident as a result of the stretching process. N material contains S-phase nucleated at the subgrain boundaries, and also decorating the residual dislocations (fig. 1). The latter are present as corrugated chains of interconnected laths whose individual length is - approximately 350nm. P material, on the other hand, contains a uniformly distributed

Fig.1. Bright field TEM image showing S- Fig.2. S-phase distribution in stretched

phase distribution on subgrain boundaries ( P ) material aged 20 hr at 463K. Dark

and within grains, for unstretched (N) field image using a I1121 S-phase reflec-

material aged 20 hr at 463K. tion; beam direction <001> .

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CYCLES TO FAILURE CYCLES T O FAILURE

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C3-764 JOURNAL DE PHYSIQUE

(a) The Endurance Ratio (E.R.) at Room Temperature

It is well established that age-hardening due to the presence of metastable, coherent precipitates, which has a large effect upon the tensile strength, may have relatively little effect upon the fatigue strength of aluminium alloys. This results in a low value of the endurance ratio, i.e. the ratio of the fatigue stress giving failure in lo6 or 10' cycles to the UTS.

For example (71, a peak-aged ternary Al-Mg-Si alloy exhibits an endurance limit of 75 MPa and a UTS of 323 MPa ( 8 ) , thus displaying an E.R. of 0.23. This low ratio is usually accounted for in terms of the progressive cutting or 'scrambling' of the precipitate particles by dislocations during fatigue. The coherent phases tend to promote inhomogeneity of slip, i.e. its concentration into intense bands, thus exascerbating the local softening process.

When, for example through the addition of dispersoid phases, slip is made more homogeneous, the E.R. is seen to increase. Thus in the A?-Mg-Si system referred to (7,8), an addition of Mn-bearing dispersoids raises the E.R. in this system to 0.29 and 0.33 for dispersoid volume percentages of 0.22 and 0.61 respectively.

In the present alloy, slip homogenisation is effected by the S-phase, which is obviously present in a considerably higher volume fraction than the disperoids in the Al-Mg-Si system referred to above. A higher E.R. might thus be expected in the 8090 material. Again, when the fracture paths illustrated in figs. 4a and 4b are compared, the same effect of slip homogenization can be discerned. In the N material (Fig. 4a) the more planar slip has resulted in a more serrated, crystallographic crack path than in the P material (Fig. 4b).

The values of the E.R. have been calculated from the present experimental data, and are shown in Table 2 as a function of the temperature of test. In the case of the room temperature data, the values of 0.44 in the N material and 0.47 in P material reflect the more effective slip homogenization by the S-phase in the present alloy than the Mn-disperoids in Al-Mg-Si, and also indicate that the more uniformly distributed particles in P material are more effective slip homogenizers than those in the N material.

T E M P E R A T U R E

(K)

Fjg.5. Variation in the modulus-normalised initial work harden- ing rate with temperature for N and P material.

(From Reference 1)

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heat-treated at temperatures up to 473K for periods of up to 1000 hours. The changes in microstructure after 24 hr, lOOhr and lOOOhr over this temperature range were identified. The tensile properties of the heat-treated materials were also measured both at room temperature and at the temperature of soak.

The maximum duration of the present fatigue tests was of the order 100 hours, and it is concluded that no significant changes in microstructure are likely to occur due to the specimens being held at elevated temperature during the course of the fatigue tests.

(b) Tensile Properties

The tensile data are shown in Table 2.

Table 2: Mechanical Test Results

Material Temperature 0.2% PS (MPa) UTS (MPa) % Elong. Endurance Ratio

Unstretched 288K 394 463 4.9 0.44

Stretched 288K 457 508 5.7 0.47

These data are consistent with those already reported in the literature (1). In particular it should be noted that tKe stretched (P) material not only exhibits an improvement in both Proof Stress and UTS at each temperature in comparison with the unstretched (N) material, but also shows an increase in ductility over this temperature range.

These changes are consistent with the changes in distribution of the S-phase brought about by the stretch, as discussed above, and the resultant slip homogenisation.

(c) Fatigue Properties

The S - N curves for stretched and unstretched material are shown in figs 3(a-c) for test temperatures of 288K, 358K and 425K respectively.

The micrographs of fig.4 illustrate the observed differences in crack propagation path at 288K between the stretched and unstretched material. They consist of montages of optical micrographs of the surfaces of specimens electropolished prior to testing. It is obvious that in the unstretched (N) material

(Fig. 4a) the surface slip bands are longer, straighter and more widely separated than in the stretched (PI material (Fig. 4b). This difference in slip distribution is accompanied by a difference in the crack profile, which appears more serrated and crystallographic in character in N material than in P.

Discussion

The S - N data obtained at room temperature (fig. 3a) are consistent with

those reported in the literature on other A1-Li alloys (e.g. refs 4-6), although no

other systematic work on the effect of stretching has been noted.

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C3-766 JOURNAL DE PHYSIQUE

(b) The Change in E.R. with Temperature

It is well-recognized that the E.R. decreases in a given alloy system with decreasing resistance to cyclic softening. Thus the highest E.R. is observed in annealed single-phase material, while work-hardened materials show a lower ratio due to instability of the dislocation structure under fatigue stressing.

Precipitation-hardened materials show even lower E.R.s due to the cyclic instability already referred to.

One would therefore expect that the E.R. would decline least with increase in temperature in those materials which exhibit the greatest resistance to recovery.

From Table 3 , it may be seen that only the stretched (P) material shows a continuous decline in E.R. This behaviour is consistent with the data of fig.5 (from Ref.l), which warnpares the change in the initial work hardening rate (normalised with respect to Young's modulus) with change in temperature of N and P material. The stretched alloy shows an initial change in whr at approximately 330K, where it decreases slowly to reach the lower level at about 490K. Unstretched material maintains its high whr to a higher temperature (375K), where it decreases more rapidly'to reach the lower whr value at approximately 500K.

The large difference in the temperature-dependence of the whr between N and P material, with dynamic recovery occurring at a lower temperature in the latter specimens, could be due either to the presence of the dislocation substructure in the stretched material, or to the difference in S-phase distribution in the two microstructures. Xia Xiaoxin and Martin (9) have investigated systematically the effect of the degree of prior stretch and of S-phase distribution upon dynamic recovery in this alloy. Their data suggest that it is essentially the presence of the dislocation substructure introduced by the stretch which promotes dynamic recovery in this temperature range.

It may be concluded, therefore, that the differences in the temperature variation of the Endurance Ratios of N and P material may be accounted for on the same basis as the difference in dynamic recovery behaviour illustrated in fig.5.

Under the conditions of the fatigue tests in the present work, the N material clearly shows a greater inherent stability and resistance to recovery than P material, and it may be concluded that the effect may be accounted for on the same basis as the whr behaviour.

Summary and Conclusions

1. A high value of the Endurance Ratio is observed at room temperature in peak-aged 8090 material. The E.R. is higher in stretched (P) than in unstretched ( N ) material, and this is interpreted as arising from the differences in S-phase distribution in the two materials: the more uniformly distributed particles of S-phase in the P material producing more effective slip homogenisation at this temperature.

2. A decline in E.R. with increasing temperature is observed in P material but not in N material. This can be accounted for by the increased resistance to dynamic recovery exhibited by N material in comparison with P material.

Acknowledgements

The authors are grateful to Professor Sir P.B. Hirsch FRS for the laboratory

facilitites made available, and to the Procurement Executive, Ministry of Defence for

support.

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