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Serrated yielding and the localized shear failure mode in aluminium alloys
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Vol Brllam
29. pp
Oool-6160.81:091553-14102.00.0 CopyrIght 0 1981 Pergamon Press Ltd
1553 to 1566, 1981
All rights reserred
SERRATED YIELDING FAILURE MODE J. E. KING, Department
of Metallurgy
AND THE LOCALIZED SHEAR IN ALUMINIUM ALLOYS C. P. YOU
and J. F. KNOTT
and Materials Science, University of Cambridge, Cambridge CB2 3Q2, U.K.
(Received 24 October
1980; in recisedform
12 March
Pembroke
Street.
1981)
Abstract-A study has been made of serrated yielding in two commercial Al-Zn-Mg alloys in the as-quenched condition. The different serration types produced in the two alloys and the shear failure mechanism observed in both notched-bend and tensile testing are related to the mechanisms of dynamic strain ageing occurring during the test. An estimate of 19.7 kJjmole for the activation energy for exchange of a solute atom and a vacancy in Al-6.2 wt% Zn, 2.5 wt”/; Mg has been made.
R&rm&Nous avons ttudie les hachures de la courbe de deformation dans deux alliages Al-ZnMg commerciaux bruts de trempe. Nous relions les differents types de hachures produites dans les deux alliages et le mecanisme de la rupture au cisaillement observe au tours de la flexion d’eprouvettes entaillees ou de la traction, aux mecanismes de vieillissement dynamique se produisant au tours des essais. Nous avons estime a 19.7 kJ/mole I’energie d’activation pour I’echange d’un atome de solute et d’une lacune dans AI-6,2’,; Zn-2.5”/, Mg [en poids). Zusammenfassung-Das ruckweise FlieDen wurde an zwei kommerziellen Al&Zn-Mg-Legierungen im abgeschrecktem Zustand untersucht. Die verschiedenen Arten dieses FlieSens und der Mechanismus des Scherbruches bei Biegung gekerbter Proben und beim Zugversuch werden auf die Mechanismen der wahrend der Versuche auftretenden dynamischen Reckalterung zurtickgeftihrt, Die Aktivierungsenergie fur das Auswechseln eines Losungsatomes mit einer Leerstelle wurde fir Al-6.2 Gew.-“” Zn -2.5 Gew.-“,, Mg zu 19,7 kJ/mole abgeschltzt.
INTRODUCTION Serrated yielding, or the Portevin-Le Chatelier [1] effect has been observed in a large number of alloy systems [Z]. including steels [3], aluminium alloys [&6], nickel-base materials [7] and copper [S-lo], magnesium [ 11. 121 and molybdenum [ 131 alloys. The effect can therefore operate in f.c.c., b.c.c. and hexagonal crystal structures in both single crystals and polycrystalline aggregates. It is generally attributed to a dynamic strain ageing process which occurs when solute atoms are diffusing sufficiently rapidly to slow down dislocations, moving under the action of an applied stress by forming ‘atmospheres’ around them. If the solutes are interstitial, for example C and N in b.c.c. iron, serrated yielding may be observed in tests carried out at temperatures close to room temperature. For substitutional solutes the effects are normally seen only at elevated temperatures, unless diffusion has been artificially accelerated, for example: by quenching from a high temperature to retain excess vacancies; by generating vacancies during plastic deformation; or by radiation damage. In a recent paper [5] the occurrence of shear bands and catastrophic, low-strain fracture in 7075 AIlZn-Mg alloy was associated with negative strain-rate sensitivity of the flow stress. Although no physical reason for the negative strain rate sensitivity was given the alloy was tested at 30°C in the as-
quenched condition, and the slow strain-rate tests exhibited serrated stress-strain curves. It seemed possible therefore that dynamic strain-ageing was responsible for the effects observed and the present research was planned to explore this point in other Al-Zn-Mg alloys in the as-quenched condition. Two alloys were studied: 7010, a high-purity, highstrength Al-Zn-Mg-Cu alloy with good toughness and stress-corrosion resistance used for aircraft structures, and a lower strength, weldable AIlZn-Mg alloy, designated CA, which has extensive military applications, including armour plating and lightweight bridges [1416]. The heat-affected-zone (HAZ) around a weld is subjected to a thermal cycle which involves a re-solution treatment followed by rapid cooling, due to metallic conduction. If the weldment is not artificially aged (as is usually the case for the welding of large structures and when welding is carried out on-site) the HAZ in service is then essentially in the as-quenched condition. The results obtained in the present research may therefore be of relevance to service application of Al-ZnMg weldments particularly at high ambient temperatures. The work has, then, attempted to investigate the relationship between dynamic strain ageing and the serrated yielding which results, and the micromechanisms of the final, rather unusual, failure mode, in both tensile and notched bend tests.
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n
132mm 27mm 334mm
Jai 43 69 Inn mm
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Metallographic specimens of both materials were etched in Keller’s Etch and photographed using a Zeiss Universal microscope. The fracture surfaces were examined in a Cambridge Instruments Stereoscan 2A scanning electron microscope operating at an accelerating voltage of 30 kV and a beam current of 150 PA.
. 391mm
RESULTS
Fig. 1. Test-piece designs. (a) Tensile test-piece, (b) SEN bend test-piece.
EXPERIMENTAL The compositions of the two alloys used are given below (wt.%): 7010 6.2Zn, 2.5Mg, 1.7Cu, O.l4Zr, 0.11 Fe, 0.07Si. CA 4.2Zn, 2.56 Mg, O.O4Cu, 0.17 Fe, 0.09 Si, 0.32 Mn, 0.01 Ti, 0.09 Cr. The designs for the tensile test-pieces and the single-edge notch bend test-pieces are shown in Fig. 1. For both materials the specimens were cut from plate with the rolling directions as indicated in Fig. 1. In 7010, the bend test-pieces were in the LT orientation [17], and the tensile test-pieces were in the L orientation: in alloy CA, the bend test-pieces were in the TL orientation and the tensile test-pieces were in the T orientation. All specimens were solution treated in a vacuum furnace for a period of 1 h, quenched into iced brine and then held under liquid nitrogen until they were required for testing, to minimize the extent of ageing during storage. The solution treatment temperatures used were 460°C for 7010 and 51O’C for CA. A small number of tensile specimens of 7010 were also aged at temperatures up to 150°C in a fluidized bed before testing at room temperature. The tensile tests were performed on a Mand screw driven testing machine at crosshead speeds ranging from 0.1 to 5 mm/min. corresponding to strain rates between -6.3 x 1O-5 - 3 x 10m3 per s. Test temperatures of 44°C. 65°C and 100°C were obtained by wrapping the specimens in heating tape, testing at - 196°C was carried out in a liquid nitrogen bath. The bend specimens were tested in three-point loading at room temperature under both load- and tests displacement-control. The load-contolled employed a Mand servo-controlled electro-hydraulic testing machine, and an Instron screw driven machine was used for the displacement-controlled tests, at ramp rates between 0.05 and 20 mm/min.
Optical micrographs of the two alloys in various orientations are presented in Fig. 2. Figure 3 shows the nominal stress-strain curves produced on testtng as-quenched 7010 at -196°C. 2O’C and 44°C and after ageing for 1: h at 100°C and testing at 20°C. At - 196°C [Fig. 3(a)] the normal strain rate dependence of cry and UTS is observed: values of these two parameters increasing with increasing strain rate; and the nominal strain to fracture at both strain rates is -407:. At 20°C [Fig. 3(b)] the strain rate dependence is reversed and the tests at slower strain rates show marked ‘serrations’, which become larger as the strain rate falls. The nominal fracture strain is lower than that at - 196”C, being about 30% for all strain-rates. At 44°C [Fig. 3(c)] behaviour similar to that at room temperature is observed, although, for the slower strain-rates, large serrations appear quite suddenly as a critical value of strain (E,) is reached. The average strain to failure was slightly less than at 20°C. Tests were also carried out at 65 and 1OOC. The results at 65’C resembled those obtained at 44’C, but E, was found to increase with decreasing strain rate and increasing temperature. An average strain to failure of -250,: was obtained. At 100°C the inverse strain-rate dependence of ur and UTS was still present, but serrations were observed only on the stress-strain curve at the fastest strain rate (2.5 x 1O-3 per s): for slower rates the curves were almost completely smooth. At i = 2.5 x 10m3 per s. the strain to failure was 23x, whereas for the curves which did not show serrations (; = 1.6 x 10e4 and 6.3 x 10m5 per s) it had increased to - 337;. After ageing for 1: h at 100°C [Fig. 3(d)] the reverse strain rate dependence was still observed although no distinct serrations were seen in room temperature testing. The strain to failure was less than 254,. Figure 4 shows the variation of UTS with strain rate for as-quenched 7010. An approximate work-hardening exponent, n, was calculated from the expression Q, = kE: and this was found to be independent of strain-rate at all temperatures investigated, having values of -0.22 at 20 and 44°C and 0.29 at - 196°C. Tensile tests were also carried out on as-quenched CA, in a different orientation to 7010 [Fig. l(a)]. Figure 5(a) shows the nominal stress-strain curves obtained at -196°C and Fig. 5(b) shows those at 20°C. At - 196°C a normal strain rate dependence is
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(a)
100
0
005
010 015 NCtllNAL STRAIN
QM
025
030
LM
m-
(cl nto rrsTs 0
0
005
010
OK
020
A, ‘I%
025
oa
35
Id)
Fig. 3. 7010 tensile tests, nominal
of
curves. (a) - 196°C; (b) 20°C; (c) 44°C; (d) After ageing for 1: h. at 100°C.
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Fig. 4. 7010 tensile test results, UTS vsi.
observed, but by 20°C this has inverted. Testing was also performed at 44 and 65”C, and these results showed similar trends to those at 20°C. As for 7010, the nominal strain to failure decreased with increasing temperature from -- 35% at - 196°C to -22% at 65”C, the maximum serration size increased with increase in temperature and decrease in strain rate, and the work-hardening exponent was independent of strain rate, having values of -0.57 at - 196°C and -0.43 between 20 and WC. The appearance of the serrations is, however, somewhat different from those in 7010. Instead of the continuous serrations falling betow the level of the stress-strain curve seen in 7010,
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‘periodic’ serrations, initially rising above the level of the curve and increasing in size as deformation proceeded, were produced in CA. It was not possible to measure any E, values for these serrations in the ranges of temperature and strain rate investigated. Figure 6 shows values of UTS as a function of strain rate in CA. The strain rate dependence of UTS is less marked than in 7010 (Fig. 4). The tensile failures in as-quenched 7010 and CA tested at and above room temperature occurred by a shear process, along a plane at between 45 -+ 60” to the tensile axis. From 20 -+ 65°C this produced macroscopically fiat fracture surfaces; by 100°C they were still of the same form but noticeably rougher. Just before the final failure parallel shear bands became visible along the specimen gauge length, most distinctly in the region of the neck, and failure then occurred in one of these bands. Fig. 7(a) and (b) shows a broken specimen of 7010 on which these bands can be seen, and Fig. 7(c) shows the intense deformation associated with the shear failure in this specimen. At - 196°C cup and cone tensile failures resulted. Figures 8(a) and (b) and 9(a) and (b) show the cup and cone failures produced at - 196°C in 7010 and CA. They consist of a central region of large, inclusion nucleated voids (8(b) and 9(b)) which have coalesced by internal necking, surrounded by a shear wall. The void forming inclusions in the CA specimen can be seen to be aligned along the rolling direction (i.e. across the test-piece diameter-see Figs 1 and 2). After ageing at 100°C and testing at 20°C the fracture surfaces were again of the shear type, but exhibited
Fig. 5. CA tensile tests, nominal u-e curves. (a) - 196°C; (b) 20°C.
KING et al.: SERRATED YIELDING AND SHEAR FAILURE
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350 -
5
-
_Y
9300 CA . . I .
65Y WC 2ooc -i-%7
250 -
I
10s
I
1
lo*
STRAIN RATE. s-1
10'1
10"
Fig. 6. CA tensile test results; UTS vs 2.
much rougher surfaces than those in the as-quenched material. On testing at temperatures between 20 and 100°C the as-quenched 7010 produced very smooth shear fracture surfaces [Fig. S(c)-(h)] with regions of heavily sheared voids near the edges [Fig. 8(d)], more equiaxed voids around large, central inclusions and almost featureless areas [Fig. 8( f)]:Above 44°C the failures become slightly rougher, with very fine voids appearing in the previously smooth regions [Fig. 8(g) and (h)]; the occurrence of these voids became more frequent with increasing temperature and decreasing strain rate. The aged specimens [Fig. 8(i) and (i)] failed in a manner similar to that of the as-quenched specimens but rougher fracture surfaces resulted, mainly covered by inclusion initiated voids with only small ‘smooth regions which, at high magnification, can be seen to be covered in very fine voids [Fig. S(i)]. The failures produced in as-quenched CA in the range 2065°C [Fig 9(c) and (d)] are essentially similar to those in 7010 although the arrangement of parallel diametral bands of shear and voids across the surface [Fig. 9(c)] reflects the change in specimen orientation. All the fracture surfaces exhibited regions of voiding around inclusions, heavily sheared voids and ‘featureless’ regions in which very fine voids started to appear with increasing temperature [Fig.
‘WI. On many of the specimens parallel ‘tram-line’ markings could be seen coming from some of the inclusion initiated voids [Fig. 8(e)] where the hard silicate inclusions were dragged across the fracture surface during the final stage of failure as the two halves of the specimen slid over each other. Figure lo(a) shows a typical trace of load vs crosshead displacement from a displacement-controlled test on a bend specimen of CA at room temperature. Similar traces were obtained in 7010. Very fine serrations are seen up to a maximum load and then the load drops off very sharply as if sudden increments of crack propagation were occurring. In Fig. 10(b) similar behaviour is shown in a load-controlled test
where, as load increases steadily, sudden increases in crosshead displacement are seen, again suggesting small bursts of crack growth. In the displacement controlled tests the crosshead displacement at which the load drops appeared was observed to increase with increasing crosshead speed, from 3.55 mm at 0.05 mm/min to 4.9 mm at 20 mm/min. Figure 11 shows the fracture surfaces associated with the behaviour shown in Fig. 10. Numerous ‘shear’ walls similar in appearance to the tensile fracture surfaces form around the bottom of the notch in both materials [Fig. 1l(a) and (bFCA, Fig. 1l(c) and (dt_7010] some of which become part of the final fracture surface. It is suggested that these regions form quite suddenly, causing the load and displacement jumps seen in Fig. 10. DISCUSSION The tensile stress-strain curves of both as-quenched 7010 and as-quenched CA exhibited serrations for test temperatures between 20°C and lOO”C, and the stress levels increased with decrease in strain rate. The curves for a test temperature of - 196°C were smooth and the stress levels increased with strain rate. This behaviour is characteristic of dynamic strain-ageing associated with one or more of the substitutional elements (presumably Mg and Zn) in the alloys. Table 1 [18] gives diffusion coefficients, Do, and activation energies, Q, for a number of common alloying elements in aluminium. By analogy with serrated yielding in iron caused by the rapid diffusion of nitrogen to dislocations, Cottrell[19] estimates that serrations appear in aluminium alloys when D 1
lo-i4im2/s
(1)
For the slowest strain rate used during these tests (6.3 x lo-‘per s) this gives D = 6 x lo-i9 m’/s. At room temperature D for Zn and Mg, the two substitutional elements present in large quantities and involved in precipitate formation in both alloys is
Fig. 7. As-quenched 7010 (a), (b) Shear bands parallel and perpendicular to the final failure; (c) Intense deformation associated with the final failure. x 410.
-2 x 1W3’ m2/s, (taking De 1 lOa m2/s and Q 5 150 kJ/mole from Table 1). No serrations would therefore be expected on the stress-strain curve until plastic deformation had increased the vacancy concentration. The diffusion coefficient is enhanced as a result of these extra vacancies according to the equation [20] : D 2: 0.12 C, exp( -E/RT) (2)
where E is the activation energy for the exchange of vacancies and solute atoms and C, is the vacancy concentration, which increases with strain as [21] : C,=&”
(3)
where B and m are constants. This suggests that, at each strain rate and temperature there will be a critical strain, E,, at which serrations
begin.
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(d)
(4
x1460
I
Fig. B(aHf).
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Fig. Q+(j). Fig. 8. 7010 fracture surfaces. (a) - 196°C x 33; (b) sheared voids at the edge of a specimen. x 1424; Featureless region. x 700; (g) 65°C. Very fine voiding shear and voiding. x 710; (i) Aged 7010 specimen.
- 196°C x 660; (c) 20°C x 34; (d) 20°C. Heavily (e) 2O’C. ‘Tram-line’ markings. x 1460; (fj 65’C. on ‘smooth’ region. x 4500; (h) 100°C. Regions of x 28; (i) Aged 7010. Very fine voids on ‘smooth’ region. x 1400.
However the specimens tested during this work were stored in liquid nitrogen after quenching from _ 500°C and the value of D at 500°C is 27 x lo- I5 for Mg and Zn. For the fastest strain rate used (6.3 x 10m3 per s) equation (1) predicts that D needs for serrations, so if the to be -6 x 10-17mZ/s vacancy concentration characteristic of 500°C is retained on quenching, enough vacancies would be present initially for serrations to appear from the start of the test, without the need for the generation of extra vacancies through plastic deformation. The presence of quenched-in vacancies can, therefore, explain why serrations started immediately in many
of the tests, without the need for a critical value of strain, lcr to be reached [Figs 3(b) and (c) and Fig.
W)l. Combining
equations
(l), (2) and (3) gives
and Fig. 12 shows a plot of log E, vs l/T at constant & (6.3 x lo-’ per s) for 7010. Equation (4) predicts that this graph should have a positive gradient of E/Rm and a positive gradient has been observed in some systems [2,6], but our results show a negative gradient, cE increasing with increasing temperature.
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Fig. 9. CA fracture surfaces. (a) - 196°C. x 32; (b) - 196°C. x 90; (c) 20°C. Note effect of orientation. x 34; (d) 65°C. Very fine voids in ‘featureless’ regions. x 1560.
Charnock’s [22] explanation for this behaviour, in aluminium and other substitutional alloys, is that Cottrell’s expressions only hold at ‘low’ temperatures, where ‘low’ means about 0.3 T,, depending on the strain rate and substitutional solute content. He suggests that there is a critical temperature, T,, at which the equilibrium vacancy concentration is sufficiently high to allow serrated yielding at E, = 0. At ‘low’ temperatures, below T,, Cottrell’s analysis applies, and E, is associated with the pinning of free dislocations by an increasingly mobile solute. At temperatures above T, the dislocations are fully aged from the start of the test and serrations begin, at a different critical strain value, when a critical number of dislocations, N, (which is a f(e)) achieve a critical
velocity, u, (u, =f(o)), at which they can break away from their solute atmospheres and multiply. He suggests that behaviour above T, is described by: rzBS = K exp( - Q’IRT)
(5)
where a, 8, y and K are constants, and Q’ = E, + E,, the activation energy for formation and movement of vacancies. From plots of E, vs l/T at constant > and taking c$ = 1 (because reported values of u and /3 lie between -0.6+ 1.2) Charnock derives values of Q between 58 and 116 kJ/mole, which are substantially larger than the value of 19.7 kJ/mole obtained from the graph in Fig. 12. The reason for this may be that the quenched-in vacancies present at the start of the tests cause the E, term to be unimportant, so
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(a)
(b)
Fig. 10. SEN bend test results. (a) Displacement-controlled test. Load vs crosshead displacement; (b) Load-controlled test. Load and crosshead displacement vs time.
Q’ 5 E,. Harris [23] has measured E, for vacancies in an Al-7 wt% Mg alloy as 19.3 kJ/mole, in excellent agreement with our experimental value of Q’. Such a mechanism not only explains the negative gradient of the E, vs l/T plot, but may be used also to rationalise the different appearance of the serrations in the two different alloys, assuming that both Zn and Mg atoms are responsible for the pinning. In 7010 the high alloy content (6.2 wt% Zn, 2.5 wt% Mg) means that in the temperature range 20+ 100°C the material is above T, (T, 4 0°C in Al, 5 wt% Mg [22]). The critical strain for the start of serrations increases with increasing temperature, but the serrations themselves do not increase in size as deformation proceeds because they are caused by fully aged dislocations breaking away from their atmospheres. The serrations are of a form such that the load falls below the general level of the stress-strain curve [Fig. 3(c)]. These observations are consistent with dislocations which are pinned from the start of deformation, so that the load drops below the general level of the curve as dislocations temporarily free themselves from their atmospheres. In CA the lower Zn + Mg content (4.2 wt%Zn, 2.56 wt?; Mg) causes the material to be just below,
but very close to, T, (r = 60°C in Al, 1 wt% Mg [22]) in the range of test temperatures from 20-65”C. No E, is measurable for these ‘Cottrell-type’ serrations, but this may be due to the quenched-in vacancies, as shown earlier, or to the proximity to T,. The serrations increase in size as the test proceeds and initially rise above the level of the curve [Fig. 5(b)], suggesting that, at first, the dislocations are not fully pinned, and as the extent of pinning increases and their atmospheres become larger an increased stress is needed to free them. Towards the end of the test the serrations start to fall below the curve and look more like those observed in 7010, as the precipitation processes proceed and the normal condition of the dislocations becomes fully aged. Serrations similar in form to these in copperindium have been designated types A and C [2,24], A being the ‘low’ temperature serrations which rise above the level of the stress-strain curve, and C the ‘high’ temperature form. They have also been called type I and type II serrations in Al-3 wt% Mg [25] but no explanation of their different appearances and behaviour was given. An alternative approach to the analysis of dynamic strain ageing has been used recently [26], based on a
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Fig. 11. CA bend specimen fracture surfaces. (a) ‘Shear’ wall running from notch. x 70; (b) Surface of ‘shear’ wall. x 300. 7010 bend specimen fracture surfaces; (c) ‘Shear’ wail. x 83; (d) Surface of ‘shear’ wall x 670.
stress concept rather than the critical strain parameter used here. Mulford and Kocks [26] suggest, as a micromechanical interpretation for this, that solute atmospheres form on forest dislocations and then drain
by pipe diffusion to the mobile dislocations whilst these are held up at the forest dislocations. The good
Table 1 Element
Cr CU
Do m’,Js
3 x lo-”
Mg
8.4 x W6 - 2.9 x lo-’ 4.1 x lo-l3 1.2 x lo-s-+1.05 x lo-4
Si Zn
9 x 10-S 1 x lo-6+l.1 x lo--+
Fe Mn
2.2 x lo-5
Q
kJ/mole 63
126-147 59-160 118-160 122
126-160 84-193
29 30
31
32 33 34 1000 /r,K
35
Fig. 12. log,, E, vs l/T.
36
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agreement between our value of Q’ and the value of E, measured by Harris [23], however, suggests that the dislocations are indeed pinned by general solute diffusion through the lattice rather than by pipe diffusion in forest walls. In addition the model of Mulford and Kocks provides no obvious explanation for the different types of serrations observed in the different strain rate and temperature regimes. Comparison of the nominal strains to failure (Ed) at - 196’C, where a normal strain-rate dependence of cry and UTS is seen, with those at room temperature where dynamic strain ageing and serrated yielding occur (7010: c 196 C = 40”,,, c, 65 C = 25”;; CA: E/ - 196’ C = ;5- ” 0) c/ 65 C = 22”,,) shows that dynamic strain ageing in these alloys may be regarded as an embrittling effect, similar to ‘Blue Brittleness’ in steels [27]. This embrittlement is also seen as a reduction in specimen deflection, at the start of failure by shear, with decreasing strain rate in room temperature notched bend tests, from 4.9 mm at a crosshead speed of 20 mm/min to 3.55 mm at 0.05 mm/min. The shear failure mode observed in tensile and bend tests on the as-quenched and quenched and aged specimens between 2O’C and 100 C appears to be a direct result of the strain ageing process. It is very similar in appearance to the much smaller, structureless regions which have been observed on the fracture surfaces of steels after tests involving extensive strain ageing [27]. It is absent in these aluminium alloys at - 196°C where cup and cone failures occur because diffusion rates are too low, even after considerable plastic deformation. to allow strain ageing to take place during the test. At - 196-C deformation continues beyond the UTS [Fig. 3(a)] whereas when dynamic strain ageing is occurring final failure occurs at or just beyond the maximum load [Fig. 3(b) and (c)l. It is suggested that the constant pinning and unpinning of dislocations, although not increasing the work-hardening exponent in the material. nevertheless reduces its work-hardening capacity. The explanation for the appearance of the failures is that, at an early stage in the tensile test (or ahead of the notch in bend) voids initiate around the weakly bonded silicate
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inclusions and these start to grow. At very low temperatures where strain ageing cannot occur the voids continue to grow until they finally link up by internal necking, producing classical ductile failure by microvoid coalescence. At higher temperatures, especially if these are close to or above T,, where the dislocations are aged from the start of the test, void growth and coalescence does not occur because the deformation localizes very rapidly. Normally dislocations start to move along the planes of maximum shear stress, and as work-hardening increases the stress required to keep them moving on these planes, dislocation motion is activated along less favourably oriented planes and flow does not localize. Where this type of dynamic strain ageing is occurring, however, once dislocation motion has started along a plane of maximum shear stress it continues because the band in which the dislocations are moving effectively becomes softer than the surrounding material where ageing and precipitation processes are continuing on the stationary dislocations, and so localized deformation can occur at stresses lower than those required to activate slip on other planes. A ‘catastrophic’ type of ductile failure finally occurs in the localized shear bands in the direction of maximum shear stress 1281, i.e. at -45 to the tensile axis and along logarithmic spirals emanating from the blunt notch in bend, as can be seen in Fig. 13. (This can happen in both the as-quenched and the aged material, but occurs to a lesser extent after ageing because the concentration of solute remaining in solution and therefore available to pin dislocations is much lower). The fine voids which start to appear in the shear regions of the fracture surfaces at higher temperatures (and in the aged specimens) may be formed around MgZn* and Mg,Zn,Al* [18] precipitates which are gradually forming as the tests proceed.
SUMMARY The two different serration types observed in asquenched 7010 and CA have been explained in terms of the different mechanisms of serrated yielding which
shear along logarithmic spirals
x80 Fig. 13. Shear failure along logarithmic spirals emanating from a blunt notch. x 80.
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occur above and below a composition-dependent and strain-rate-dependent critical temperature, T,. Below T, the dislocations are free of atmospheres at the start of deformation and serrations occur as solute atoms catch up with dislocations. The serrations then initially rise above the stress-strain curve. Above T, the dislocations are aged from the start of the test and serrations start when a whole group of dislocations reaches a critical velocity, breaks free of its atmosphere and crosses the specimen as a narrow deformation bend. These serrations drop below the level of the stress-strain curve. Although below T, a plot of E, for the start of serrations vs l/T has a positive gradient [2,6], above T,, it has a negative gradient as a consequence of the change in mechanism. The intensely sheared failure mode often observed in tensile tests on aluminium alloys at room temperature can also occur in notched bend tests and is a direct result of the dynamic
strain
ageing processes,
which cause serrated yielding, producing flow localization at an early stage during the test. Acknowledgements-Thanks are due to Professor R. W. K. Honeycombe for the provision of research facilities, Mr R. J. H. Smyth for experimental assistance, Rolls Royce Ltd., Aero Division for financial support of one of the authors (J.E.K.) and Professor R. S. You for financial support for another (C.P.Y.).
REFERENCES A. Portevin and F. Le Chatelier, C.r. hebd. Skanc Acad. Sci. Paris 176, 507 (1923). B. J. Brindley and P. J. Worthington, Metal/. Rev. 145, 101 (1970). J. D. Baird, Metall. Rev. 149, 1 (1971).
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4. V. A. Phillips, A. J. Swain and R. Eborall, J. Inst. Metals 81 103 (1961). 5. N. Chung, J. C. Embury, J. D. Evensen, R. G. Hoagland and C. M. Sargent, Acta metal/. 25, 377 (1977). 6. S. Miura, Proc. Int. Conf: on Mechanical Behaviour of Materials, Kyoto, Vol. 1, p. 128 (1971). 7. R. A. Mulford, Metal/. T&s. A 10, 1527 (1979). 8. L. M. T. Hookin. J. Inst. Metals 84. 102 (1955-56). 9. M. M. Hut&in&n and F. P. Bullen,‘Phil. Mag. 7, i535 (1962).
10. D. J. H. Corderoy, Ph.D. Thesis, University of Sheffield (1962). 11. S. L. Couling, Acta metall. 7, 133 (1959). 12. K. G. Kent and A. Kelly, J. Inst. Metals. 93, 536 (1964-5). 13. A. Lawley and R. Maddin, Trans. metal/. Sot. A.I.M.E. 224, 573 (1962).
14. Y. Bresson, Rev. Aluminium 341, 481 (1966). 15. B. Track, Metals Mater. 6, 263 (1972). 16. J. G. Whitman, Welding Institute Select Conference on Weldable AI-Zn-Mg Alloys, W.I. Abingdon. Paper 8. (1970). 17. British Standards Institution, BS5447: 1977, Methods jar Plane Strain Fracture
18 Metallic Materials.
Toughness (K,,)
Testing of
L. F. Mondolfo, Metal/. Rev. 153, 95 (1971). 19: A. H. Cottrell, Phi/. Mag. 44, 31 (1953). 20. A. H. Cottrell, Phil. Mng. 44, 897 (1953). z:: F. Seitz, Ado. Phys. 1, 43 (1952). 23 W. Charnock, Phil. Mug. 20, 427 (1969). . S. G. Harris, Vacancies and Orher Point Defects in Metals and Alloys, p. 220. Institute of Metals, London (1958). 24. P. J. Worthington, Ph.D. Thesis, University of Sheffield, 1962. 25. B. J. Brindley and P. J. Worthington, Acta metall. 17, 1357 (1969). 26. R. A. Mulford and U. F. Kocks, Acra metal/. 27, 1125 (1979). 27. I. D. G. Groom, Ph.D. Thesis, University of Cambridge (1971). 28. J. F. Knott, Fracture 1977 Vol 1. p. 61. ICM, Waterloo, Canada, June (1977).
,M~‘~‘illurqI‘~r
Prmted
m Great
Vol Brllam
29. pp
Oool-6160.81:091553-14102.00.0 CopyrIght 0 1981 Pergamon Press Ltd
1553 to 1566, 1981
All rights reserred
SERRATED YIELDING FAILURE MODE J. E. KING, Department
of Metallurgy
AND THE LOCALIZED SHEAR IN ALUMINIUM ALLOYS C. P. YOU
and J. F. KNOTT
and Materials Science, University of Cambridge, Cambridge CB2 3Q2, U.K.
(Received 24 October
1980; in recisedform
12 March
Pembroke
Street.
1981)
Abstract-A study has been made of serrated yielding in two commercial Al-Zn-Mg alloys in the as-quenched condition. The different serration types produced in the two alloys and the shear failure mechanism observed in both notched-bend and tensile testing are related to the mechanisms of dynamic strain ageing occurring during the test. An estimate of 19.7 kJjmole for the activation energy for exchange of a solute atom and a vacancy in Al-6.2 wt% Zn, 2.5 wt”/; Mg has been made.
R&rm&Nous avons ttudie les hachures de la courbe de deformation dans deux alliages Al-ZnMg commerciaux bruts de trempe. Nous relions les differents types de hachures produites dans les deux alliages et le mecanisme de la rupture au cisaillement observe au tours de la flexion d’eprouvettes entaillees ou de la traction, aux mecanismes de vieillissement dynamique se produisant au tours des essais. Nous avons estime a 19.7 kJ/mole I’energie d’activation pour I’echange d’un atome de solute et d’une lacune dans AI-6,2’,; Zn-2.5”/, Mg [en poids). Zusammenfassung-Das ruckweise FlieDen wurde an zwei kommerziellen Al&Zn-Mg-Legierungen im abgeschrecktem Zustand untersucht. Die verschiedenen Arten dieses FlieSens und der Mechanismus des Scherbruches bei Biegung gekerbter Proben und beim Zugversuch werden auf die Mechanismen der wahrend der Versuche auftretenden dynamischen Reckalterung zurtickgeftihrt, Die Aktivierungsenergie fur das Auswechseln eines Losungsatomes mit einer Leerstelle wurde fir Al-6.2 Gew.-“” Zn -2.5 Gew.-“,, Mg zu 19,7 kJ/mole abgeschltzt.
INTRODUCTION Serrated yielding, or the Portevin-Le Chatelier [1] effect has been observed in a large number of alloy systems [Z]. including steels [3], aluminium alloys [&6], nickel-base materials [7] and copper [S-lo], magnesium [ 11. 121 and molybdenum [ 131 alloys. The effect can therefore operate in f.c.c., b.c.c. and hexagonal crystal structures in both single crystals and polycrystalline aggregates. It is generally attributed to a dynamic strain ageing process which occurs when solute atoms are diffusing sufficiently rapidly to slow down dislocations, moving under the action of an applied stress by forming ‘atmospheres’ around them. If the solutes are interstitial, for example C and N in b.c.c. iron, serrated yielding may be observed in tests carried out at temperatures close to room temperature. For substitutional solutes the effects are normally seen only at elevated temperatures, unless diffusion has been artificially accelerated, for example: by quenching from a high temperature to retain excess vacancies; by generating vacancies during plastic deformation; or by radiation damage. In a recent paper [5] the occurrence of shear bands and catastrophic, low-strain fracture in 7075 AIlZn-Mg alloy was associated with negative strain-rate sensitivity of the flow stress. Although no physical reason for the negative strain rate sensitivity was given the alloy was tested at 30°C in the as-
quenched condition, and the slow strain-rate tests exhibited serrated stress-strain curves. It seemed possible therefore that dynamic strain-ageing was responsible for the effects observed and the present research was planned to explore this point in other Al-Zn-Mg alloys in the as-quenched condition. Two alloys were studied: 7010, a high-purity, highstrength Al-Zn-Mg-Cu alloy with good toughness and stress-corrosion resistance used for aircraft structures, and a lower strength, weldable AIlZn-Mg alloy, designated CA, which has extensive military applications, including armour plating and lightweight bridges [1416]. The heat-affected-zone (HAZ) around a weld is subjected to a thermal cycle which involves a re-solution treatment followed by rapid cooling, due to metallic conduction. If the weldment is not artificially aged (as is usually the case for the welding of large structures and when welding is carried out on-site) the HAZ in service is then essentially in the as-quenched condition. The results obtained in the present research may therefore be of relevance to service application of Al-ZnMg weldments particularly at high ambient temperatures. The work has, then, attempted to investigate the relationship between dynamic strain ageing and the serrated yielding which results, and the micromechanisms of the final, rather unusual, failure mode, in both tensile and notched bend tests.
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Jai 43 69 Inn mm
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Metallographic specimens of both materials were etched in Keller’s Etch and photographed using a Zeiss Universal microscope. The fracture surfaces were examined in a Cambridge Instruments Stereoscan 2A scanning electron microscope operating at an accelerating voltage of 30 kV and a beam current of 150 PA.
. 391mm
RESULTS
Fig. 1. Test-piece designs. (a) Tensile test-piece, (b) SEN bend test-piece.
EXPERIMENTAL The compositions of the two alloys used are given below (wt.%): 7010 6.2Zn, 2.5Mg, 1.7Cu, O.l4Zr, 0.11 Fe, 0.07Si. CA 4.2Zn, 2.56 Mg, O.O4Cu, 0.17 Fe, 0.09 Si, 0.32 Mn, 0.01 Ti, 0.09 Cr. The designs for the tensile test-pieces and the single-edge notch bend test-pieces are shown in Fig. 1. For both materials the specimens were cut from plate with the rolling directions as indicated in Fig. 1. In 7010, the bend test-pieces were in the LT orientation [17], and the tensile test-pieces were in the L orientation: in alloy CA, the bend test-pieces were in the TL orientation and the tensile test-pieces were in the T orientation. All specimens were solution treated in a vacuum furnace for a period of 1 h, quenched into iced brine and then held under liquid nitrogen until they were required for testing, to minimize the extent of ageing during storage. The solution treatment temperatures used were 460°C for 7010 and 51O’C for CA. A small number of tensile specimens of 7010 were also aged at temperatures up to 150°C in a fluidized bed before testing at room temperature. The tensile tests were performed on a Mand screw driven testing machine at crosshead speeds ranging from 0.1 to 5 mm/min. corresponding to strain rates between -6.3 x 1O-5 - 3 x 10m3 per s. Test temperatures of 44°C. 65°C and 100°C were obtained by wrapping the specimens in heating tape, testing at - 196°C was carried out in a liquid nitrogen bath. The bend specimens were tested in three-point loading at room temperature under both load- and tests displacement-control. The load-contolled employed a Mand servo-controlled electro-hydraulic testing machine, and an Instron screw driven machine was used for the displacement-controlled tests, at ramp rates between 0.05 and 20 mm/min.
Optical micrographs of the two alloys in various orientations are presented in Fig. 2. Figure 3 shows the nominal stress-strain curves produced on testtng as-quenched 7010 at -196°C. 2O’C and 44°C and after ageing for 1: h at 100°C and testing at 20°C. At - 196°C [Fig. 3(a)] the normal strain rate dependence of cry and UTS is observed: values of these two parameters increasing with increasing strain rate; and the nominal strain to fracture at both strain rates is -407:. At 20°C [Fig. 3(b)] the strain rate dependence is reversed and the tests at slower strain rates show marked ‘serrations’, which become larger as the strain rate falls. The nominal fracture strain is lower than that at - 196”C, being about 30% for all strain-rates. At 44°C [Fig. 3(c)] behaviour similar to that at room temperature is observed, although, for the slower strain-rates, large serrations appear quite suddenly as a critical value of strain (E,) is reached. The average strain to failure was slightly less than at 20°C. Tests were also carried out at 65 and 1OOC. The results at 65’C resembled those obtained at 44’C, but E, was found to increase with decreasing strain rate and increasing temperature. An average strain to failure of -250,: was obtained. At 100°C the inverse strain-rate dependence of ur and UTS was still present, but serrations were observed only on the stress-strain curve at the fastest strain rate (2.5 x 1O-3 per s): for slower rates the curves were almost completely smooth. At i = 2.5 x 10m3 per s. the strain to failure was 23x, whereas for the curves which did not show serrations (; = 1.6 x 10e4 and 6.3 x 10m5 per s) it had increased to - 337;. After ageing for 1: h at 100°C [Fig. 3(d)] the reverse strain rate dependence was still observed although no distinct serrations were seen in room temperature testing. The strain to failure was less than 254,. Figure 4 shows the variation of UTS with strain rate for as-quenched 7010. An approximate work-hardening exponent, n, was calculated from the expression Q, = kE: and this was found to be independent of strain-rate at all temperatures investigated, having values of -0.22 at 20 and 44°C and 0.29 at - 196°C. Tensile tests were also carried out on as-quenched CA, in a different orientation to 7010 [Fig. l(a)]. Figure 5(a) shows the nominal stress-strain curves obtained at -196°C and Fig. 5(b) shows those at 20°C. At - 196°C a normal strain rate dependence is
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100
0
005
010 015 NCtllNAL STRAIN
QM
025
030
LM
m-
(cl nto rrsTs 0
0
005
010
OK
020
A, ‘I%
025
oa
35
Id)
Fig. 3. 7010 tensile tests, nominal
of
curves. (a) - 196°C; (b) 20°C; (c) 44°C; (d) After ageing for 1: h. at 100°C.
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Fig. 4. 7010 tensile test results, UTS vsi.
observed, but by 20°C this has inverted. Testing was also performed at 44 and 65”C, and these results showed similar trends to those at 20°C. As for 7010, the nominal strain to failure decreased with increasing temperature from -- 35% at - 196°C to -22% at 65”C, the maximum serration size increased with increase in temperature and decrease in strain rate, and the work-hardening exponent was independent of strain rate, having values of -0.57 at - 196°C and -0.43 between 20 and WC. The appearance of the serrations is, however, somewhat different from those in 7010. Instead of the continuous serrations falling betow the level of the stress-strain curve seen in 7010,
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‘periodic’ serrations, initially rising above the level of the curve and increasing in size as deformation proceeded, were produced in CA. It was not possible to measure any E, values for these serrations in the ranges of temperature and strain rate investigated. Figure 6 shows values of UTS as a function of strain rate in CA. The strain rate dependence of UTS is less marked than in 7010 (Fig. 4). The tensile failures in as-quenched 7010 and CA tested at and above room temperature occurred by a shear process, along a plane at between 45 -+ 60” to the tensile axis. From 20 -+ 65°C this produced macroscopically fiat fracture surfaces; by 100°C they were still of the same form but noticeably rougher. Just before the final failure parallel shear bands became visible along the specimen gauge length, most distinctly in the region of the neck, and failure then occurred in one of these bands. Fig. 7(a) and (b) shows a broken specimen of 7010 on which these bands can be seen, and Fig. 7(c) shows the intense deformation associated with the shear failure in this specimen. At - 196°C cup and cone tensile failures resulted. Figures 8(a) and (b) and 9(a) and (b) show the cup and cone failures produced at - 196°C in 7010 and CA. They consist of a central region of large, inclusion nucleated voids (8(b) and 9(b)) which have coalesced by internal necking, surrounded by a shear wall. The void forming inclusions in the CA specimen can be seen to be aligned along the rolling direction (i.e. across the test-piece diameter-see Figs 1 and 2). After ageing at 100°C and testing at 20°C the fracture surfaces were again of the shear type, but exhibited
Fig. 5. CA tensile tests, nominal u-e curves. (a) - 196°C; (b) 20°C.
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350 -
5
-
_Y
9300 CA . . I .
65Y WC 2ooc -i-%7
250 -
I
10s
I
1
lo*
STRAIN RATE. s-1
10'1
10"
Fig. 6. CA tensile test results; UTS vs 2.
much rougher surfaces than those in the as-quenched material. On testing at temperatures between 20 and 100°C the as-quenched 7010 produced very smooth shear fracture surfaces [Fig. S(c)-(h)] with regions of heavily sheared voids near the edges [Fig. 8(d)], more equiaxed voids around large, central inclusions and almost featureless areas [Fig. 8( f)]:Above 44°C the failures become slightly rougher, with very fine voids appearing in the previously smooth regions [Fig. 8(g) and (h)]; the occurrence of these voids became more frequent with increasing temperature and decreasing strain rate. The aged specimens [Fig. 8(i) and (i)] failed in a manner similar to that of the as-quenched specimens but rougher fracture surfaces resulted, mainly covered by inclusion initiated voids with only small ‘smooth regions which, at high magnification, can be seen to be covered in very fine voids [Fig. S(i)]. The failures produced in as-quenched CA in the range 2065°C [Fig 9(c) and (d)] are essentially similar to those in 7010 although the arrangement of parallel diametral bands of shear and voids across the surface [Fig. 9(c)] reflects the change in specimen orientation. All the fracture surfaces exhibited regions of voiding around inclusions, heavily sheared voids and ‘featureless’ regions in which very fine voids started to appear with increasing temperature [Fig.
‘WI. On many of the specimens parallel ‘tram-line’ markings could be seen coming from some of the inclusion initiated voids [Fig. 8(e)] where the hard silicate inclusions were dragged across the fracture surface during the final stage of failure as the two halves of the specimen slid over each other. Figure lo(a) shows a typical trace of load vs crosshead displacement from a displacement-controlled test on a bend specimen of CA at room temperature. Similar traces were obtained in 7010. Very fine serrations are seen up to a maximum load and then the load drops off very sharply as if sudden increments of crack propagation were occurring. In Fig. 10(b) similar behaviour is shown in a load-controlled test
where, as load increases steadily, sudden increases in crosshead displacement are seen, again suggesting small bursts of crack growth. In the displacement controlled tests the crosshead displacement at which the load drops appeared was observed to increase with increasing crosshead speed, from 3.55 mm at 0.05 mm/min to 4.9 mm at 20 mm/min. Figure 11 shows the fracture surfaces associated with the behaviour shown in Fig. 10. Numerous ‘shear’ walls similar in appearance to the tensile fracture surfaces form around the bottom of the notch in both materials [Fig. 1l(a) and (bFCA, Fig. 1l(c) and (dt_7010] some of which become part of the final fracture surface. It is suggested that these regions form quite suddenly, causing the load and displacement jumps seen in Fig. 10. DISCUSSION The tensile stress-strain curves of both as-quenched 7010 and as-quenched CA exhibited serrations for test temperatures between 20°C and lOO”C, and the stress levels increased with decrease in strain rate. The curves for a test temperature of - 196°C were smooth and the stress levels increased with strain rate. This behaviour is characteristic of dynamic strain-ageing associated with one or more of the substitutional elements (presumably Mg and Zn) in the alloys. Table 1 [18] gives diffusion coefficients, Do, and activation energies, Q, for a number of common alloying elements in aluminium. By analogy with serrated yielding in iron caused by the rapid diffusion of nitrogen to dislocations, Cottrell[19] estimates that serrations appear in aluminium alloys when D 1
lo-i4im2/s
(1)
For the slowest strain rate used during these tests (6.3 x lo-‘per s) this gives D = 6 x lo-i9 m’/s. At room temperature D for Zn and Mg, the two substitutional elements present in large quantities and involved in precipitate formation in both alloys is
Fig. 7. As-quenched 7010 (a), (b) Shear bands parallel and perpendicular to the final failure; (c) Intense deformation associated with the final failure. x 410.
-2 x 1W3’ m2/s, (taking De 1 lOa m2/s and Q 5 150 kJ/mole from Table 1). No serrations would therefore be expected on the stress-strain curve until plastic deformation had increased the vacancy concentration. The diffusion coefficient is enhanced as a result of these extra vacancies according to the equation [20] : D 2: 0.12 C, exp( -E/RT) (2)
where E is the activation energy for the exchange of vacancies and solute atoms and C, is the vacancy concentration, which increases with strain as [21] : C,=&”
(3)
where B and m are constants. This suggests that, at each strain rate and temperature there will be a critical strain, E,, at which serrations
begin.
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Fig. B(aHf).
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Fig. Q+(j). Fig. 8. 7010 fracture surfaces. (a) - 196°C x 33; (b) sheared voids at the edge of a specimen. x 1424; Featureless region. x 700; (g) 65°C. Very fine voiding shear and voiding. x 710; (i) Aged 7010 specimen.
- 196°C x 660; (c) 20°C x 34; (d) 20°C. Heavily (e) 2O’C. ‘Tram-line’ markings. x 1460; (fj 65’C. on ‘smooth’ region. x 4500; (h) 100°C. Regions of x 28; (i) Aged 7010. Very fine voids on ‘smooth’ region. x 1400.
However the specimens tested during this work were stored in liquid nitrogen after quenching from _ 500°C and the value of D at 500°C is 27 x lo- I5 for Mg and Zn. For the fastest strain rate used (6.3 x 10m3 per s) equation (1) predicts that D needs for serrations, so if the to be -6 x 10-17mZ/s vacancy concentration characteristic of 500°C is retained on quenching, enough vacancies would be present initially for serrations to appear from the start of the test, without the need for the generation of extra vacancies through plastic deformation. The presence of quenched-in vacancies can, therefore, explain why serrations started immediately in many
of the tests, without the need for a critical value of strain, lcr to be reached [Figs 3(b) and (c) and Fig.
W)l. Combining
equations
(l), (2) and (3) gives
and Fig. 12 shows a plot of log E, vs l/T at constant & (6.3 x lo-’ per s) for 7010. Equation (4) predicts that this graph should have a positive gradient of E/Rm and a positive gradient has been observed in some systems [2,6], but our results show a negative gradient, cE increasing with increasing temperature.
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Fig. 9. CA fracture surfaces. (a) - 196°C. x 32; (b) - 196°C. x 90; (c) 20°C. Note effect of orientation. x 34; (d) 65°C. Very fine voids in ‘featureless’ regions. x 1560.
Charnock’s [22] explanation for this behaviour, in aluminium and other substitutional alloys, is that Cottrell’s expressions only hold at ‘low’ temperatures, where ‘low’ means about 0.3 T,, depending on the strain rate and substitutional solute content. He suggests that there is a critical temperature, T,, at which the equilibrium vacancy concentration is sufficiently high to allow serrated yielding at E, = 0. At ‘low’ temperatures, below T,, Cottrell’s analysis applies, and E, is associated with the pinning of free dislocations by an increasingly mobile solute. At temperatures above T, the dislocations are fully aged from the start of the test and serrations begin, at a different critical strain value, when a critical number of dislocations, N, (which is a f(e)) achieve a critical
velocity, u, (u, =f(o)), at which they can break away from their solute atmospheres and multiply. He suggests that behaviour above T, is described by: rzBS = K exp( - Q’IRT)
(5)
where a, 8, y and K are constants, and Q’ = E, + E,, the activation energy for formation and movement of vacancies. From plots of E, vs l/T at constant > and taking c$ = 1 (because reported values of u and /3 lie between -0.6+ 1.2) Charnock derives values of Q between 58 and 116 kJ/mole, which are substantially larger than the value of 19.7 kJ/mole obtained from the graph in Fig. 12. The reason for this may be that the quenched-in vacancies present at the start of the tests cause the E, term to be unimportant, so
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(a)
(b)
Fig. 10. SEN bend test results. (a) Displacement-controlled test. Load vs crosshead displacement; (b) Load-controlled test. Load and crosshead displacement vs time.
Q’ 5 E,. Harris [23] has measured E, for vacancies in an Al-7 wt% Mg alloy as 19.3 kJ/mole, in excellent agreement with our experimental value of Q’. Such a mechanism not only explains the negative gradient of the E, vs l/T plot, but may be used also to rationalise the different appearance of the serrations in the two different alloys, assuming that both Zn and Mg atoms are responsible for the pinning. In 7010 the high alloy content (6.2 wt% Zn, 2.5 wt% Mg) means that in the temperature range 20+ 100°C the material is above T, (T, 4 0°C in Al, 5 wt% Mg [22]). The critical strain for the start of serrations increases with increasing temperature, but the serrations themselves do not increase in size as deformation proceeds because they are caused by fully aged dislocations breaking away from their atmospheres. The serrations are of a form such that the load falls below the general level of the stress-strain curve [Fig. 3(c)]. These observations are consistent with dislocations which are pinned from the start of deformation, so that the load drops below the general level of the curve as dislocations temporarily free themselves from their atmospheres. In CA the lower Zn + Mg content (4.2 wt%Zn, 2.56 wt?; Mg) causes the material to be just below,
but very close to, T, (r = 60°C in Al, 1 wt% Mg [22]) in the range of test temperatures from 20-65”C. No E, is measurable for these ‘Cottrell-type’ serrations, but this may be due to the quenched-in vacancies, as shown earlier, or to the proximity to T,. The serrations increase in size as the test proceeds and initially rise above the level of the curve [Fig. 5(b)], suggesting that, at first, the dislocations are not fully pinned, and as the extent of pinning increases and their atmospheres become larger an increased stress is needed to free them. Towards the end of the test the serrations start to fall below the curve and look more like those observed in 7010, as the precipitation processes proceed and the normal condition of the dislocations becomes fully aged. Serrations similar in form to these in copperindium have been designated types A and C [2,24], A being the ‘low’ temperature serrations which rise above the level of the stress-strain curve, and C the ‘high’ temperature form. They have also been called type I and type II serrations in Al-3 wt% Mg [25] but no explanation of their different appearances and behaviour was given. An alternative approach to the analysis of dynamic strain ageing has been used recently [26], based on a
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Fig. 11. CA bend specimen fracture surfaces. (a) ‘Shear’ wall running from notch. x 70; (b) Surface of ‘shear’ wall. x 300. 7010 bend specimen fracture surfaces; (c) ‘Shear’ wail. x 83; (d) Surface of ‘shear’ wall x 670.
stress concept rather than the critical strain parameter used here. Mulford and Kocks [26] suggest, as a micromechanical interpretation for this, that solute atmospheres form on forest dislocations and then drain
by pipe diffusion to the mobile dislocations whilst these are held up at the forest dislocations. The good
Table 1 Element
Cr CU
Do m’,Js
3 x lo-”
Mg
8.4 x W6 - 2.9 x lo-’ 4.1 x lo-l3 1.2 x lo-s-+1.05 x lo-4
Si Zn
9 x 10-S 1 x lo-6+l.1 x lo--+
Fe Mn
2.2 x lo-5
Q
kJ/mole 63
126-147 59-160 118-160 122
126-160 84-193
29 30
31
32 33 34 1000 /r,K
35
Fig. 12. log,, E, vs l/T.
36
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agreement between our value of Q’ and the value of E, measured by Harris [23], however, suggests that the dislocations are indeed pinned by general solute diffusion through the lattice rather than by pipe diffusion in forest walls. In addition the model of Mulford and Kocks provides no obvious explanation for the different types of serrations observed in the different strain rate and temperature regimes. Comparison of the nominal strains to failure (Ed) at - 196’C, where a normal strain-rate dependence of cry and UTS is seen, with those at room temperature where dynamic strain ageing and serrated yielding occur (7010: c 196 C = 40”,,, c, 65 C = 25”;; CA: E/ - 196’ C = ;5- ” 0) c/ 65 C = 22”,,) shows that dynamic strain ageing in these alloys may be regarded as an embrittling effect, similar to ‘Blue Brittleness’ in steels [27]. This embrittlement is also seen as a reduction in specimen deflection, at the start of failure by shear, with decreasing strain rate in room temperature notched bend tests, from 4.9 mm at a crosshead speed of 20 mm/min to 3.55 mm at 0.05 mm/min. The shear failure mode observed in tensile and bend tests on the as-quenched and quenched and aged specimens between 2O’C and 100 C appears to be a direct result of the strain ageing process. It is very similar in appearance to the much smaller, structureless regions which have been observed on the fracture surfaces of steels after tests involving extensive strain ageing [27]. It is absent in these aluminium alloys at - 196°C where cup and cone failures occur because diffusion rates are too low, even after considerable plastic deformation. to allow strain ageing to take place during the test. At - 196-C deformation continues beyond the UTS [Fig. 3(a)] whereas when dynamic strain ageing is occurring final failure occurs at or just beyond the maximum load [Fig. 3(b) and (c)l. It is suggested that the constant pinning and unpinning of dislocations, although not increasing the work-hardening exponent in the material. nevertheless reduces its work-hardening capacity. The explanation for the appearance of the failures is that, at an early stage in the tensile test (or ahead of the notch in bend) voids initiate around the weakly bonded silicate
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inclusions and these start to grow. At very low temperatures where strain ageing cannot occur the voids continue to grow until they finally link up by internal necking, producing classical ductile failure by microvoid coalescence. At higher temperatures, especially if these are close to or above T,, where the dislocations are aged from the start of the test, void growth and coalescence does not occur because the deformation localizes very rapidly. Normally dislocations start to move along the planes of maximum shear stress, and as work-hardening increases the stress required to keep them moving on these planes, dislocation motion is activated along less favourably oriented planes and flow does not localize. Where this type of dynamic strain ageing is occurring, however, once dislocation motion has started along a plane of maximum shear stress it continues because the band in which the dislocations are moving effectively becomes softer than the surrounding material where ageing and precipitation processes are continuing on the stationary dislocations, and so localized deformation can occur at stresses lower than those required to activate slip on other planes. A ‘catastrophic’ type of ductile failure finally occurs in the localized shear bands in the direction of maximum shear stress 1281, i.e. at -45 to the tensile axis and along logarithmic spirals emanating from the blunt notch in bend, as can be seen in Fig. 13. (This can happen in both the as-quenched and the aged material, but occurs to a lesser extent after ageing because the concentration of solute remaining in solution and therefore available to pin dislocations is much lower). The fine voids which start to appear in the shear regions of the fracture surfaces at higher temperatures (and in the aged specimens) may be formed around MgZn* and Mg,Zn,Al* [18] precipitates which are gradually forming as the tests proceed.
SUMMARY The two different serration types observed in asquenched 7010 and CA have been explained in terms of the different mechanisms of serrated yielding which
shear along logarithmic spirals
x80 Fig. 13. Shear failure along logarithmic spirals emanating from a blunt notch. x 80.
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KING et al.:
SERRATED YIELDING
occur above and below a composition-dependent and strain-rate-dependent critical temperature, T,. Below T, the dislocations are free of atmospheres at the start of deformation and serrations occur as solute atoms catch up with dislocations. The serrations then initially rise above the stress-strain curve. Above T, the dislocations are aged from the start of the test and serrations start when a whole group of dislocations reaches a critical velocity, breaks free of its atmosphere and crosses the specimen as a narrow deformation bend. These serrations drop below the level of the stress-strain curve. Although below T, a plot of E, for the start of serrations vs l/T has a positive gradient [2,6], above T,, it has a negative gradient as a consequence of the change in mechanism. The intensely sheared failure mode often observed in tensile tests on aluminium alloys at room temperature can also occur in notched bend tests and is a direct result of the dynamic
strain
ageing processes,
which cause serrated yielding, producing flow localization at an early stage during the test. Acknowledgements-Thanks are due to Professor R. W. K. Honeycombe for the provision of research facilities, Mr R. J. H. Smyth for experimental assistance, Rolls Royce Ltd., Aero Division for financial support of one of the authors (J.E.K.) and Professor R. S. You for financial support for another (C.P.Y.).
REFERENCES A. Portevin and F. Le Chatelier, C.r. hebd. Skanc Acad. Sci. Paris 176, 507 (1923). B. J. Brindley and P. J. Worthington, Metal/. Rev. 145, 101 (1970). J. D. Baird, Metall. Rev. 149, 1 (1971).
AND SHEAR FAILURE
4. V. A. Phillips, A. J. Swain and R. Eborall, J. Inst. Metals 81 103 (1961). 5. N. Chung, J. C. Embury, J. D. Evensen, R. G. Hoagland and C. M. Sargent, Acta metal/. 25, 377 (1977). 6. S. Miura, Proc. Int. Conf: on Mechanical Behaviour of Materials, Kyoto, Vol. 1, p. 128 (1971). 7. R. A. Mulford, Metal/. T&s. A 10, 1527 (1979). 8. L. M. T. Hookin. J. Inst. Metals 84. 102 (1955-56). 9. M. M. Hut&in&n and F. P. Bullen,‘Phil. Mag. 7, i535 (1962).
10. D. J. H. Corderoy, Ph.D. Thesis, University of Sheffield (1962). 11. S. L. Couling, Acta metall. 7, 133 (1959). 12. K. G. Kent and A. Kelly, J. Inst. Metals. 93, 536 (1964-5). 13. A. Lawley and R. Maddin, Trans. metal/. Sot. A.I.M.E. 224, 573 (1962).
14. Y. Bresson, Rev. Aluminium 341, 481 (1966). 15. B. Track, Metals Mater. 6, 263 (1972). 16. J. G. Whitman, Welding Institute Select Conference on Weldable AI-Zn-Mg Alloys, W.I. Abingdon. Paper 8. (1970). 17. British Standards Institution, BS5447: 1977, Methods jar Plane Strain Fracture
18 Metallic Materials.
Toughness (K,,)
Testing of
L. F. Mondolfo, Metal/. Rev. 153, 95 (1971). 19: A. H. Cottrell, Phi/. Mag. 44, 31 (1953). 20. A. H. Cottrell, Phil. Mng. 44, 897 (1953). z:: F. Seitz, Ado. Phys. 1, 43 (1952). 23 W. Charnock, Phil. Mug. 20, 427 (1969). . S. G. Harris, Vacancies and Orher Point Defects in Metals and Alloys, p. 220. Institute of Metals, London (1958). 24. P. J. Worthington, Ph.D. Thesis, University of Sheffield, 1962. 25. B. J. Brindley and P. J. Worthington, Acta metall. 17, 1357 (1969). 26. R. A. Mulford and U. F. Kocks, Acra metal/. 27, 1125 (1979). 27. I. D. G. Groom, Ph.D. Thesis, University of Cambridge (1971). 28. J. F. Knott, Fracture 1977 Vol 1. p. 61. ICM, Waterloo, Canada, June (1977).