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Origin of acoustic emission in AlZnMg alloys II: Copper-containing quaternary alloys
Materials Science and Engineering, 59 (1983) 169-183
169
Origin of Acoustic Emission in Al-Zn-Mg Alloys II: Copper-containing Quaternary Alloys C. B. SCRUBY, H. N. G. WADLEY* and K. L. RUSBRIDGE
Atomic Energy Research Establishment, Harwell, Oxon. (Gt. Britain) (Received October 2, 1982)
SUMMARY
Acoustic emission has been measured during the deformation and fracture o f t w o quaternary A I - 5 . 5 w t. % Z n - 2 . 5 w t. %Mg- 1.6w t. %Cu alloys with different grain sizes as a function o f aging. The results have been compared with a previous study o f the ternary alloy (Part I). It has been f o u n d that in the larger-grain-size alloy the m o s t energetic acoustic emission source is the formation o f intense slip bands in the peak-aged condition similar to the ternary alloy. However, reducing the grain size (by the addition o f zirconium to the quaternary alloy) suppressed both this emission source and weaker sources in overaged states. The effects are interpreted in terms o f a reduction in the area o f slip events. S phase inclusion fracture was observed to have occurred in the quaternary alloys and may have generated small-amplitude emissions in peak-aged material. 1. INTRODUCTION
When metals are deformed, they emit elastic waves. In pure metals these elastic waves are generated by moving dislocations. The properties (amplitude and frequency) of the elastic waves which determine their detectability axe controlled by the velocity and propagation distance of individual or coordinated groups of dislocations [1]. Impure metals contain inclusions and the brittle fracture of these has been suggested as an additional possible source of acoustic emission [2-5]. Previous work [6] (Part I) has shown that in inclusion-free A1-5.5wt.%Zn2.5wt.%Mg the generation of elastic waves *Present address: National Bureau of Standards, Washington, DC 20234, U.S.A. 0025-5416/83/0000-0000/$03.00
(acoustic emission) by moving dislocations was critically dependent on the precipitate distribution, which in turn was determined by the aging treatment. In particular, it was found that in the peak-aged condition where dislocations sheared 77' precipitates, resulting in the unstable formation of intense slip bands, very energetic acoustic emission signals were generated whilst other conditions generated, in the main, much weaker emission. In order to investigate further the influence of metallurgical variables and the presence of inclusions, t w o further alloys based on the A1-5.5wt.%Zn-2.5wt.%Mg system (F68) were prepared. In the first of these, referred to as F69, 1.6 wt.% Cu was added. This resulted in the formation of insoluble S phase inclusions, a slight modification to precipitate composition and distribution on aging, a coppersaturated matrix and a reduced grain size. The second alloy, designated F70, was further modified by the addition of 0.16 wt.% Zr, an alloy otherwise identical in composition with F69. The zirconium forms small insoluble particles which pin grain boundaries during solution treatments restricting grain growth, thus enabling the effect of grain size to be determined independently. In this paper we report the effects of these additional metallurgical variables on acoustic emission from alloys in the same aged states as in Part I. Using mechanical property measurements and microstructural characterization techniques, we attempt to relate the emission sensitivity to these variables to the changes in deformation mechanism. 2. MATERIALS
Two ingots, F69 and F70, were prepared using the technique described in Part I for © Elsevier Sequoia/Printed in The Netherlands
170 TABLE 1 Bulk chemical compositions of A1-Zn-Mg alloys Alloy
Zn (wt.%)
Mg (wt.%)
Cu (wt.%)
Zr (wt.%)
Fe (wt.%)
AI (wt.%)
F69 Top F69 Bottom
5.51 5.48
2.49 2.47
1.59 1.59
---
0.02 0.02
Balance Balance
F70 Top F70 Bottom
5.36 5.54
2.36 2.47
1.48 1.60
0.16 0.15
0.03 0.03
Balance Balance
All other elements were below the detection level of 0.01 wt.%.
alloy F68 [6]. The ingot compositions measured at the top and bottom of each ingot are given in Table 1. The ingots were rolled to 50 mm plate, using procedures identical with those for F68. An electron microprobe was used to seek evidence of microsegregation in both the as-rolled condition and following solution treatment for 2 h at 465 °Co Insignificant levels of segregation were observed in these two alloys, unlike F68. Inclusions were present in alloys F69 and F70 (Fig. 1). Electron microprobe chemical analysis of the inclusions in alloy F69 indicated the presence of two inclusion types: Cu0.4Alo.05Mg0.1vZn0.03, type I; Cu0.21A10.33Mg0.2sZn0.3, type II. The type II inclusions were of similar composition to the S phase, A12MgCu, but with zinc substitution of some aluminium. The inclusions were up to about 10 pm in size and were distributed
I
20pm
!
Fig. 1. Inclusions in undeformed AI-5.5wt.%Zn2.5wt.%Mg-l.6wt.%Cu alloy (F69).
in a dendritic structure with an interdendrite spacing of 50-100 ~zm. Tensile samples of dumb-beU geometry [6] were machined from the rolled plates and solution treated at 465 °C for 2 h and quenched into warm water. In alloy F69 this resulted in a fully recrystallized microstructure with equiaxed grains whose average size (by the line intercept method) was about 230 pm (compared with 420 #m in alloy F68 studied previously). In alloy FT0, recrystallization had been inhibited, as anticipated, by the addition of zirconium and the microstructure was that of a recovered rolled structure, i.e. elongated grains up to 200 pm in length and 50 pm across, containing a subgrain structure with an average subgrain size of about I0 ~m. 3. EXPERIMENTAL METHODS
The solution-treated samples were divided into three groups, and pairs of samples were given one of three types of aging treatment: (a) isothermal aging at 20 °C (natural aging) for up to 30 days; (b) isothermal aging at 120 °C (120 °C aging) for up to 41 days; (c) 24 h at 120 °C plus isothermal aging at 180 °C (double aging) for up to 12 h. The samples were electropolished to remove any oxide layer formed during heat treatment. One of each pair was used for acoustic emission measurement and the other for detailed mechanical property studies, using the techniques and procedures described in Part I. After testing, the fracture surfaces of the samples were examined to characterize the fracture mode. Additionally, transmission electron microscopy samples were cut from the gauge regions to observe the precipitate distribution and the dislocation structures of the deformed material.
171 4. ACOUSTIC EMISSION RESULTS
4.1. Alloy F68 The detailed acoustic emission results for the ternary alloy (F68) are given in Part I [6] but are summarized here for completeness. The behaviour of the quenched material was characterized by the generation of many small emission signals at yield and during the load drops accompanying serrated yielding (the Portevin-Le Chatelier effect) and by the generation of a few large-amplitude signals randomly throughout the test. In contrast, material that had been naturally aged or aged at 120 °C emitted isolated but individually energetic signals near yield and many very energetic signals during plastic deformation. These energetic signals were most intensely emitted from samples that had been aged for 24 h or longer at 120 °C. A second aging treatment at 180 °C caused the disappearance of energetic signals during plastic deformation and the reappearance of small emission signals at yield. 4.2. Alloy F69 The quenched condition generated very similar acoustic emission signals to the ternary alloy. A broad peak in activity composed of many small-amplitude signals was observed at yield (Fig. 2(a)). This was followed by further weak emission associated with serrated yielding during post-yield deformation. In this
region, isolated large-amplitude signals were also observed. The effect of natural aging is shown in Figs. 2(b)-2(d) and again was similar in many ways to that of alloy F68 in that the yield region activity quickly disappeared and highly energetic signals were emitted during plastic deformation. However, in contrast with F68, the naturally aged copper-containing alloy copiously emitted small-amplitude signals during serrated yielding (Fig. 3). The emission tended to be linked with the drops in load accompanying each serration. Aging at 120 °C again resulted in the generation of many extremely energetic signals (Fig. 4). This activity was most intense in the vicinity of peak hardness, similar to alloy F68. Double aging at 180 °C (Fig. 5) resulted in a gradual disappearance of the highly energetic burst emission, the reappearance of yield region emission and, after long aging times, continuous acoustic emission during postyield deformation. A cluster of events was observed during deformation as fracture was initiated and the nominal stress fell. To summarize, the effect of the copper addition seems to be rather small, the most noticeable effect being the more prominent so-called "continuous" emission during postyield deformation in underaged and overaged conditions.
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Fig. 3. A portion of the stress-strain, acoustic emission (AE)-strain data showing correlation between dynamic strain aging and acoustic emission for F69 (aged for 7 days at 20 °C).
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Fig. 4. Effect of aging at 120 °C on acoustic emission and mechanical properties for F69: (a) quenched; (b) aged for 6 rain; (c) aged for 24 min; (d) aged for 120 min; (e) aged for 1440 rain; (f) aged for 60 000 min.
4.3. Alloy F70 The addition of 0.16 wt.% Zr to F69 had a great effect on the acoustic emission activity. Whilst a yield peak was still emitted in the quenched condition (Fig. 6(a)), there was very little post-yield activity during serrated yielding. The yield activity decreased with natural aging (Figs. 6(b)-6(d)) but, in marked contrast with alloys F68 and F69, there was
only a small increase in activity during postyield deformation in naturally aged samples. Similarly, for samples aged at 120 °C, the highly energetic signals observed in F68 and F69 near peak hardness were almost totally absent apart from a small cluster of events associated with fracture (Figs. 7(e) and 7(f)). Overaging also had a different effect in F70 from in F68 and F69 (Fig. 8). No yield
173
60¢
¢0
region peak was observed, and the continuous acoustic emission during post-yield deformation of F69 was absent. The sample aged for 24 h at 120 °C plus 12 h at 180 °C was noteworthy because it generated no detectable elastic waves other than from the stress drop of final fracture.
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5. MECHANICAL PROPERTIES
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'1 ''~[-'L'J`2 "=
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The effect of natural aging for all three alloys is shown in Fig. 9. It can be seen that b o t h Oo.x and OUTSwere proportional to log(aging time) for all three alloys. In part the effect of the composition difference was to change the grain size and thus to displace the curves so that the strength of F70 was greater than that of F69 which in turn was greater
0
Extension(ram) Fig. 5. Effect of secondary aging at 180 °C on the acoustic emission and mechanical properties for F69 : (a) 24 h at 120 °C only; (b) secondary aging for 1 h; (c) secondary aging for 2 h; (d) secondary aging for 4 h; (e) secondary aging for 8 h; (f) secondary aging for 1 2 h .
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174 t.O 400 f200 ~ i hA~,i "AcOustic Emission Power
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Fig. 7. Effect of aging at 120 ~C on acoustic emission and mechanical properties for F70: (a) quenched; (b) aged for 6 rain; (c) aged for 24 min; (d) aged for 120 rain; (e) aged for 1440 rain; (f) aged for 60 000 rain. than that of F68. Similarly the copper addition did n o t greatly affect the kinetics of hardening at 120 °C (Fig. 10); all three alloys had a similar time dependence, but at each time the strength decreased from F70 to F69 to F68. Double aging caused a loss of strength (Fig. 11). The rate of decrease in alloy F69 was lower, whilst in F70 it was greater, than that of F68.
5.2. Ductility We used the plastic strain to maximum true stress as the ductility indicator. From Tables 2, 3 and Part I it is seen that natural aging in all three alloys tended to reduce ductility, although this effect was less evident in F70. Isothermal aging at 120 °C caused an almost m o n o t o n i c loss of ductility with increasing strength in all three alloys. Double aging at 180 °C had surprisingly little systematic effect on the ductility. 5.3. Work-hardening rates The stress-strain curves were numerically differentiated with respect to strain to determine the work-hardening rate as a function
of plastic strain for all the tests and these are s h o w n in Figs. 12-14. Several effects of aging were observed. The work-hardening rate decreased with plastic strain in all cases to a value of about 1.2 GPa at 5% strain, for all heat treatments and compositions. At small strains (about 1%), natural and 120 °C aging tended to increase the work-hardening rate in all three alloys. At small strains, double aging at 180 °C causes a decrease in work-hardening rate that was most marked in F70 (Fig. 14). At larger strains (about 5%) the work-hardening rate tended to decrease as the peak hardness was approached. The work-hardening exponent n in the relation O"T ---~ k C T n
(where o T is the true stress, eT the true plastic strain and k a constant) is also c o m m o n l y used to describe the work-hardening characteristics of a metal. Values of n at three strains are given in Tables 2 and 3. Whilst it is usually considered that n is independent of strain, the precise recording techniques used here show that the n values usually increase with strain.
175
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Fig. 8. Effect of secondary aging at 180 °C on acoustic emission and mechanical properties for F70: (a) 24 h at 120 °C only; (b) secondary aging for 1 h; (c) sec()ndary aging for 2 h; (d) secondary aging for 4 h; (e) secondary aging for 8 h; (f) secondary aging for 12 h.
lOO
1101
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103
1104
105
Ag ing Time (rain)
The values at a strain of about 5% are probably the best to use for comparison with other published data. In all three alloys, natural aging in general decreased n, aging at 120 °C resulted in larger decreases whilst double aging led t o a recovery in the n value. For a given aging condition the
(b)
Fig. 10. Effect of aging at 120 °C on (a) the 0.1% proof stresses and (b) the ultimate tensile stresses of alloy F68 (o), F69 (A) and F70 (u). [ ] indicates premature failure.
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176 TABLE 2 Mechanical properties: F69
Heat treatment
Stress (MPa) O0.1
Isochronai aging at 20 °C Quenched 119 1 day 155 30 days 249 Isochronal aging Quenched 6min 12min 24min 36 min 60min 120 min 1440mm 6000 min 60000 min
O0.2
Ultimate strain a
Work-hardening rate (GPa) Work-hardening exponent for the following strains for the following strains
OUTS
60.1
61
e5
eO.1
61
65
129 166 259
274 312 385
0.140 0.164 0.110
9.2 12 13.5
2.5 2.7 4.1
1.7 1.1 1.2
0.07 0.07 0.06
0.12 0.10 0.08
0.38 0.21 0.17
at 120 °C 119 129 209 220 211 220 233 246 272 285 293 303 298 308 398 415 428 445 423 445
274 356 344 362 383 383 414 487 485 476
0.140 0.152 0.147 0.152 0.074 0.047 0.119 0.058 0.010 0.026
9.2 16.7 14.2 17.5 17.5 20.0 19.0 25.9 31.0 28.9
2.5 3.2 3.1 3.3 4.3 4.5 4.4 5.0 5.1 3.2
1.7 1.5 1.2 1.3 1.3 1.2 1.1 0.9 --
0.07 0.09 0.07 0.08 0.07 0.09 0.08 0.06 0.05
0.38 0.22 0.19 0.18 0.16 0.13 0.12 0.09 --
--
--
0.12 0.08 0.08 0.08 0.08 0.07 0.07 0.06 0.05 0.04
487 436 422 417 402
0.058 0.031 0.041 0.056 0.045
25.9 18.6 14.7 11.0 13.0
5.0 3.3 3.0 3.0 3.2
0.9 --1.0 1.0
0.06 0.04 0.04 0.03 0.04
0.06 0.04 0.04 0.04 0.05
0.09 --0.10 0.10
Doubleagingat180°C 0h 398 2h 376 4h 358 8h 341 12 h 321
415 382 363 345 329
--
aplastic strain corresponding to maximum true stress. 500
i[:: Z
z ,ooL
•
300
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100
_~ 100
n E
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0
(b)
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Fig. 11. Effect of secondary aging at 180 °C on (a) the 0.1% proof stresses and (b) the ultimate tensile stresses of alloy F68 (e), F69 (A) and F70 (m) in the 120 °C, 24 h aged condition. n v a l u e a t 5% s t r a i n d e c r e a s e d f r o m F 6 8 t o F69 to F70. 6. MICROSTRUCTURAL CHARACTERIZATION
6.1. P r e c i p i t a t e s Transmission electron microscopy examinat i o n o f s a m p l e s o f F 6 9 a n d F 7 0 a g e d a t 1 2 0 °C
s h o w e d t h a t t h e p r e c i p i t a t e size a n d i n t e r particle spacing distributions in these alloys were similar to those in F68. A high density of small precipitates formed within the grains a f t e r 2 4 h a t 1 2 0 °C w i t h a n a r r o w p r e c i p i t a t e ~ free z o n e a d j a c e n t to grain b o u n d a r i e s . The boundaries themselves, including the subgrain b o u n d a r i e s in F70, were d e c o r a t e d with larger
177 TABLE 3 M e c h a n i c a l p r o p e r t i e s : F 70
Heat treatment
Stress (MPa) O0.1
Isochronal aging at 20 °C Quenched 160 1 day 211 7 days 246 30 d a y s 275 Isochronal aging Quenched 6 rain 12 rain 24 r a i n 36 m i n 6 0 rain 1 2 0 rain 1 4 4 0 rain 6000 min 6 0 0 0 0 rain
Ultimate strain a
00.2
Work-hardening rate ( G P a ) Work-hardening exponent for the following strains for the following strains
OUTS eO.1
61
65
60,1
61
65
170 216 260 288
330 358 401 438
0.165 0.145 0.114 0.163
12 7 17 15
2.6 3.1 4.8 4.7
1.8 1.4 1.4 1.3
0.08 0.03 0.08 0.06
0.09 0.09 0.10 0.09
0.32 0.22 0.18 0.16
at 120 °C 160 170 220 232 239 249 263 277 265 297 316 323 336 340 407 424 419 456 426 453
330 382 385 397 392 444 454 508 505 502
0.165 0.176 0.149 0.109 0.099 0.109 0.098 0.081 0.073 0.077
12 15 17 16 19 19 18 28 31 32
2.6 3.7 4.0 4.1 4.2 4.6 5.0 5.3 3.7 3.6
1.8 1.5 1.2 1.2 1.2 1.3 1,2 1,0 0.9 0.7
0.08 0.07 0.08 0.07 0.08 0.07 0.06 0.06 -0.06
0.09 0.09 0.08 0.08 0.08 0.08 0.07 0.06 0.04 0.04
0.32 0.21 0.17 0.16 0.14 0.14 0.13 0.08 0.08 0.06
508 485 447 434 373
0.081 0.061 0.081 0.073 0.057
28 27 15 11 13
5.3 3.8 3.0 3.2 2.7
1,0 0,9 1.0 1.2 1,3
0.06 0.05 0.03 0.03 0.05
0.06 0.05 0.04 0.04 0.05
0.08 0.09 0.10 0.13 0.16
Double aging at 180°C 0 h 407 2 h 406 4 h 358 8 h 345 12 h 280
424 421 366 353 288
aplastic s t r a i n c o r r e s p o n d i n g t o m a x i m u m t r u e stress. 40 35
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0.01 Plastic Strain
0.02
F i g . 12. E f f e c t o f n a t u r a l aging ( 2 0 ° C ) o n t h e yield region w o r k - h a r d e n i n g r a t e : (a) F 6 8 ; (b) F 6 9 ; ( c ) F 7 0 .
precipitates. Prolonged aging at 120 °C gave precipitate coarsening (Fig. 15). Peel et al. [7] have found that aging at 120 °C for 24 h gives almost entirely 7' precipitates in an alloy
of the F69 composition, rather than a mixture of 7' and 77 as in F68. In contrast, aging at 180 °C for 8 h did produce a different precipitate size distribution
178
~. 151\ \n |1 _ \ \ ~ d o ~
@h
2
d~ys
~oy5 Ib h
0.01 Ptast,¢ Strain
002
24rain
0
0.01 Plastic Strain
0.02
(a)
0
0.01 PLastic Strain
0.02
(b)
0
(c)
Fig. 13. Effect of aging at 120 °C on the yield region work-hardening rate: (a) F68; (b) F69; (c) F70.
30 o Q_ ®
"-" 25
25
0 n~
¢F
~, 20
~ zo
2O n 0
"u
20 15
er
0h
i
Os.
g 0
"
~ (a)
~10
10
5
0
Sh 1 '~
4h
I 0.01 Plastic Stra,n
--
I 002
~
0
I
I
0.01 Plash¢ Strain
(b)
0.02
001 Plastic Strain
I
002
(c)
Fig. 14. Effect of double aging at 180°C on the yield region work-hardening rate: (a) F68; (b) F69; (c) F70.
!OOnm i
Fig. 15. Transmission electron micrograph of fine precipitates in alloy F69 aged at 120 °C for 41 days.
i
Fig. 16. Transmission electron micrograph showing coarse matrix precipitates and precipitate-free zone at grain boundary in alloy F69 aged at 120 °C for 24 h and at 180°C for 8 h.
179
(a)
(b) Fig. 17. (a) Transmission electron micrograph of precipitates in alloy F69 aged at 120 °C for 24 h and at 180 °C for 8 h; (b) transmission electron micrograph of precipitates in alloy F70 aged at 120 °C for 24 h and at 180°C for 8 h. in F69 and F70 compared with that in F68. In F68, rod-shaped precipitates with a precipitate density of 1016 cm -3 were observed. In F69 and F70 there was a similar density of the rod-shaped precipitates but a distribution of smaller spherical precipitates was also present, giving a total precipitate density of 3 × 1016 cm -3 (Figs. 16 and 17).
6.2. Dislocation structures The dislocation structures in F69 and F70 samples were also examined using transmission electron microscopy. F69 was found to be similar to F68, showing uniform dislocation structures in quenched and underaged conditions with a transition to localized slip in bands around peak hardness. The dislocation density within the bands was less in F69 than in F68 (Fig. 18). There was a reversion to uniform deformation after a secondary aging treatment at 180 °C. In contrast, F70 exhibited little heterogeneous slip.
Fig. 18. Transmission electron micrograph of lightly deformed bands in alloy F69 aged at 120 °C for 4 days.
Fig. 19. Scanning electron micrograph of fracture surface of alloy F69 aged at 120 °C for 6 min, showing mainly ductile fracture.
6.3. Fractography Fractography indicated that alloy F69 exhibited a tendency to undergo a transition from ductile dimple fracture in the underaged state (Fig. 19) to a mixed fracture consisting of small areas of intergranular failure linked by regions of ductile dimples when peak aged. The addition of copper reduced the extent of intergranular cracking in peak-aged material but did not totally eliminate it. A combination of intergranular and ductile fracture was also found in F70 alloys aged at 120 °C to peak hardness, with the proportion of intergranular fracture decreasing away from peak hardness. In F70 aged for 24 h at 120 °C, the fracture path appeared to follow the boundaries of the elongated grains. Fracture along strings of inclusions at grain boundaries and fracture along subgrain boundaries were also seen (Fig. 20).
180 that considerable plastic relaxation had accompanied the cracking.
7. DISCUSSION
Fig. 20. Scanning electron micrograph of fracture surface of alloy F70 aged at 120 °C for 24 h, showing intergranular fracture, fracture at inclusions A and subgrain cracking B.
Fig. 21. Scanning electron rnicrograph of a polished gauge region of alloy F69, showing cracks across S phase inclusions. The tensile axis was horizontal.
6.4. Inclusions
Scanning electron microscopy of fracture surface indicated that the majority of inclusions present on the surface had fractured in both F69 and F70. The gauge regions of fractured samples were then sectioned and polished to reveal the inclusions below the fracture surface. It was found (Fig. 21) that almost every inclusion had fractured, the larger ones containing several cracks. The fractures tended to be normal to the tensile axis, indicating mode I failure. The large crackopening displacements (about 1 pm) suggested
It was shown in Part I that during deformation the product of the propagation distance a and velocity v of individual (or a group of n) dislocations must exceed a threshold value (determined by experimental sensitivity) for detectable elastic waves to be emitted. For aluminium alloys, nay >/3.6 × 1 0 - 2 m 2 s -1 for detectable emission to be generated, where n is the number of dislocations moving together. The smallest detectable motion of a single dislocation might thus correspond to the creation of a loop 180 ~zm in radius whose mean radial velocity was 200 m s-1 [1]. Since b o t h the distance of dislocation propagation and the velocity are distributed quantities, only a fraction of the deformation processes of a sample are detectable, and the acoustic emission measurements indicate this to be very dependent on metallurgical variables such as precipitate t y p e and distribution and grain size. In a detailed analysis of results from the ternary alloy in Part I, it was shown that three types of deformation mechanism were able to generate detectable acoustic emission and each depended differently on microstructure. (i) In quenched or very lightly aged conditions, dislocation unpinning from solute atoms either at yield or during plastic flow (dynamic strain aging) generated, in the main, weak signals. (ii) In peak-aged samples the formation of intense slip bands associated with the shearing of fine precipitates was an intense source of individually energetic elastic waves. (iii) In overaged samples, with a wider precipitate spacing, the cooperative motion of groups of dislocations from one set of precipitates to the next could generate many weak signals. The copper-containing alloys could potentially have a fourth emission source: the fracture of S phase inclusions. In the following the acoustic emission characteristics of the basic ternary, coppercontaining quaternary and zirconium-doped quaternary alloys are compared and discussed in terms of the effect of microstructural differences on the ability of these deforma-
181
tion and fracture mechanisms to generate detectable acoustic emission. We shall consider each mechanism in turn.
7.1. Dislocation unpinning Similar yield region acoustic emission activity in the quenched state was observed in all three alloys. In the ternary alloy it was shown that, with the aid of internal stress concentrations (possibly ahead of dislocation pile-ups), a sufficient stress could be reached for individual dislocations to unpin themselves from impurity atoms segregated at dislocation cores. After escape the dislocations could multiply and propagate at a high speed, resulting in a sudden local stress relaxation. It seems reasonable to expect the magnitude of this stress relaxation to be determined by the distance of dislocation propagation, i.e. by the grain size at small strains. In the ternary alloy the grain size was a b o u t 430 pm and signals from single dislocations were predicted to be just detectable. In F69 the grain size was less (about 230 pm) b u t the flow stress and possibly the dislocation velocity were higher, and so we might again expect similar levels of acoustic emission. In alloy F70, however, whilst the grain size was still a b o u t 200 pm, there was also a well-developed subgrain structure with a spacing of a b o u t 10 pm which would severely limit the distance of propagation. In this case a single dislocation, even if it propagated at the unlikely limiting (transverse) wave speed of 3200 m s-1, would have an av product of only 3.2 × 10 -2 m 2 s-1 and so be barely detectable. The generation of the observed levels of acoustic emission in F70 could have three possible explanations. (i) There could be coordinated escape of many dislocations, each propagating over a mean subgrain diameter. (ii) Single dislocations could move in the largest subgrains at the top of the grain size distribution so that the propagation distance is greater than the mean subgrain size. (iii) The subgrains were not strong enough to arrest high speed dislocations, and so they propagated to the grain boundaries. It is n o t possible to determine the contribution of these various mechanisms of slip. It is possible that to some extent all three occurred.
During post-yield deformation of b o t h quenched and aged samples, more serrated yielding and its associated acoustic emission were observed in alloys F69 and F70 than in F68. The generation of acoustic emission after large plastic strains when the mean free path for individual dislocations has been reduced to the dislocation cell size (a few micrometres) can only come from the coordinated motion of many thousands of dislocations. Evidence for such widespread sudden motion is given by the stress-strain curves. Here, serrated yielding is analogous to very low frequency (less than 10 Hz) acoustic emission. Each time that sufficient unpinning occurred to give a detectable stress drop on the stress-strain curve, groups of acoustic emissions were observed. The load cell detected the static change in stress whilst the emission signals were from the micromechanisms of slip that a c c o m m o d a t e d the change. The enhanced activity in F69 and F70 could be related to the higher flow stress of these alloys which would aid the escape of dislocations and cause them to propagate at a high velocity. Whilst this flow stress increase was achieved in part b y a decrease in grain size, this did not limit the acoustic emission activity directly since this was determined by the much smaller dislocation cell size.
7.2. Precipitate shear It was argued in Part I for F68 that, in peak-aged material with a fine distribution of small precipitates, dislocation propagation involved precipitate shearing which in turn led to unstable slip localization in bands and the emission of energetic signals. The results from F69 and F70 support this conclusion as the energetic signals were only detected in peak-aged F69 in which slip bands were found using transmission electron microscopy, whereas F70 always showed a uniform dislocation structure and deformed quietly. The peak-aged F69 alloy showed a somewhat reduced level of activity compared with F68. This change could be brought about if the precipitate strength was greater in the copper-containing quaternary alloy, reducing the tendency for precipitate shearing, or it could be related to the smaller grain size which would reduce the length of individual dislocation pile-ups and the local stress on
182 precipitates. The loss of acoustic emission in FT0 might also be the result of the small grain size. A few signals were detected towards the end of deformation in this alloy, suggesting that only then could sufficient stress be applied to the leading dislocation in a pile-up to shear precipitates. However, these signals could also be related to fracture.
7.3. Orowan looping The results from double-aged F68 reported in Part I were surprising in that detectable acoustic emission was recorded during yield although the slip distance might be expected to be limited to the interparticle spacing kp. It was postulated that the acoustic emission was due to the sudden growth of dislocation loops that had been pinned b y an inner set of precipitates lying at a radius r to a set lying at a radius r + kp. The detection of loop growth may then be possible provided that the loop radius (which is limited by the grain size) and interparticle spacing are large. Both the grain size and the interparticle spacing were less in alloy F69 than those of F68 b u t this appears to have had only a small effect on elastic wave generation at yield. This could perhaps result from the larger yield Stress of F69 which could drive free dislocations at a higher velocity. Alloy F70, whilst having a similar interparticle spacing to F69, had a much smaller effective grain size (10 #m compared with 230 pm). Thus, the area swept out when a loop expanded from one set of particles to the next would have been reduced b y a factor of a b o u t 20, and this seems sufficient to ensure no emission of detectable elastic waves.
7.4. Inclusions There is good evidence that inclusion fracture (and possibly decohesion) is an important source of acoustic emission in commercial grade aluminium alloys [2, 4, 5]. Alloys F69 and F70 b o t h contained large numbers of S phase inclusions, but the acoustic emission results provide little evidence that they are an important source of detectable acoustic emission. Indeed, the overaged samples of alloy F70 generated little or no acoustic emission at all, suggesting that inclusion cracking is a slower process in the S phase inclusions. This is supported by recent results of Cousland and Scala [ 5 ] who have shown that the acoustic
emission from inclusion cracking in aluminium alloys is predominantly from brittle iron-containing inclusions. Let us try to estimate the emission amplitudes that might be generated by S phase inclusion fracture. It has been shown [8] that, for a detectable signal to be emitted by a mode I microcrack,
t 3C1Eh Uminl ~2
acVcl i 2 ~ ~
J
where ac is the crack radius (assuming a circular microcrack), vc the radial crack speed, C1 the longitudinal wave speed, E Young's modulus, h the crack-to-detector separation, UmL~ the smallest detectable displacement, v Poisson's ratio and o the applied stress of the instant of fracture. Using, for aluminium, C1 = 6400 m s-1, E = 80 GPa and v = 0.3 and assuming Umm = 10 -14 m, h = 40 m m and a = 400 MPa (for peak-aged material) gives the condition that acv¢~2 > 1.45 × 10 -~ m 3/2 s-1/2. Using a value for ac of a b o u t 3 #m (representing one of the larger microcracks), then the crack growth speed would only need to be more than 25 m s-1 for a detectable signal to be released. Let us assume a range of inclusion sizes with 0.5 pm < ac < 5 pm and a range of speeds with 10 m s-1 < v~ < 1000 m s-1. Then the model predicts a range of acoustic emission amplitudes given by 10-16 m ~ Umm < 10-12 m About half this range exceeds the detection threshold of 10-14 m. Thus it is concluded that S phase inclusion fracture is a possible weak emission source, but only when the
larger inclusions crack at moderate to high velocities. The absence of emission during the period of deformation when cracking of the majority of the inclusions was expected to be occurring [9] (at large plastic strains) suggests that crack growth speeds are generally well below 1000 m s -1. The stronger emission from iron-rich inclusions in the materials studied by Scala and Cousland [4] and by Cousland and Scala [5] is consistent with a larger inclusion size (1-14 pm) and possibly a higher fracture speed. Their work also supports the notion of the critical a~v~~2 value for detectable signals of a b o u t 3 × 10 -4 m 3/2 s -1/2. When material
183
with a smaller inclusion size (1-6 pm) was tested, very little emission was observed.
iron-rich inclusions in commercial alloys could well be a more important emission source.
8. CONCLUSIONS
ACKNOWLEDGMENTS
The acoustic emission activity of quaternary A1-5.5wt.%Zn-2.5wt.%Mg-l.6wt.%Cu alloy has been shown to be critically dependent on composition and heat treatment. The origin of these effects has been attributed to changes in deformation mechanism. The following facts have emerged from this study. (1) The addition of copper to the ternary alloy increases the weak emission associated with dynamic strain aging. The effect is greatest for small-grain-size (high flow stress) samples. (2) The most energetic emission source in the base ternary alloy, that associated with strain localization in the peak-aged state, is only slightly reduced by the addition of copper alone. (3) The addition of zirconium which greatly reduces the grain size causes almost complete disappearance of strain localization and the associated emission. It also leads to the disappearance of the emission associated with Orowan looping. Both effects are consistent with the restriction of the movement of dislocations by grain boundaries. (4) Fracture of the S phase inclusions appears to be a small-amplitude emission source, although other work has shown that
We wish to express our gratitude to D. Stockham-Jones for experimental assistance and Dr. J. Hudson for discussion of this work. The study was funded by the Royal Aircraft Establishment, Farnborough.
REFERENCES 1 C. B. Scruby, H. N. G. Wadley and J. E. Sinclair, Philos. Mag. A, 44 (2) (1981) 249. 2 M. A. Hamstad, R. Bianchetti and A. K. Mukherjee, Eng. Fract. Mech., 9 (1977) 663. 3 K. Ono, G. Huang and H. Hatano, 8th World Conf. on Non-Destructive Testing, Cannes, September 1976. 4 C. M. Scala and S. MeK. Cousland, Defense Adva nced Research Projec ts A g e n c y - U.S. Air Force Conf. on Quantitative Nondestructive Evaluation, Boulder, CO, August 1981. 5 S. McK. Cousland and C. M. Scala, Met. Sc£, 15 (1981) 609. 6 K. L. Rusbridge, C. B. Scruby and H. N. G. Wadley, Mater. ScL Eng., 59 (1983) 151. 7 C. J. Peel, D. Clark, P. Poole and R. A. Farra, R A E Tech. Rep. TR 78110, 1978 (Royal Aircraft Establishment, Farnborough). 8 C. B. Scruby, C. Jones, J. M. Titchmarsh and H. N. G. Wadley, Met. Sci., 15 (1981) 241. 9 R. H. van Stone, R. H. Merchant and J. R. Low, A S T M Spec. Tech. Publ. 566, 1974, p. 93.
169
Origin of Acoustic Emission in Al-Zn-Mg Alloys II: Copper-containing Quaternary Alloys C. B. SCRUBY, H. N. G. WADLEY* and K. L. RUSBRIDGE
Atomic Energy Research Establishment, Harwell, Oxon. (Gt. Britain) (Received October 2, 1982)
SUMMARY
Acoustic emission has been measured during the deformation and fracture o f t w o quaternary A I - 5 . 5 w t. % Z n - 2 . 5 w t. %Mg- 1.6w t. %Cu alloys with different grain sizes as a function o f aging. The results have been compared with a previous study o f the ternary alloy (Part I). It has been f o u n d that in the larger-grain-size alloy the m o s t energetic acoustic emission source is the formation o f intense slip bands in the peak-aged condition similar to the ternary alloy. However, reducing the grain size (by the addition o f zirconium to the quaternary alloy) suppressed both this emission source and weaker sources in overaged states. The effects are interpreted in terms o f a reduction in the area o f slip events. S phase inclusion fracture was observed to have occurred in the quaternary alloys and may have generated small-amplitude emissions in peak-aged material. 1. INTRODUCTION
When metals are deformed, they emit elastic waves. In pure metals these elastic waves are generated by moving dislocations. The properties (amplitude and frequency) of the elastic waves which determine their detectability axe controlled by the velocity and propagation distance of individual or coordinated groups of dislocations [1]. Impure metals contain inclusions and the brittle fracture of these has been suggested as an additional possible source of acoustic emission [2-5]. Previous work [6] (Part I) has shown that in inclusion-free A1-5.5wt.%Zn2.5wt.%Mg the generation of elastic waves *Present address: National Bureau of Standards, Washington, DC 20234, U.S.A. 0025-5416/83/0000-0000/$03.00
(acoustic emission) by moving dislocations was critically dependent on the precipitate distribution, which in turn was determined by the aging treatment. In particular, it was found that in the peak-aged condition where dislocations sheared 77' precipitates, resulting in the unstable formation of intense slip bands, very energetic acoustic emission signals were generated whilst other conditions generated, in the main, much weaker emission. In order to investigate further the influence of metallurgical variables and the presence of inclusions, t w o further alloys based on the A1-5.5wt.%Zn-2.5wt.%Mg system (F68) were prepared. In the first of these, referred to as F69, 1.6 wt.% Cu was added. This resulted in the formation of insoluble S phase inclusions, a slight modification to precipitate composition and distribution on aging, a coppersaturated matrix and a reduced grain size. The second alloy, designated F70, was further modified by the addition of 0.16 wt.% Zr, an alloy otherwise identical in composition with F69. The zirconium forms small insoluble particles which pin grain boundaries during solution treatments restricting grain growth, thus enabling the effect of grain size to be determined independently. In this paper we report the effects of these additional metallurgical variables on acoustic emission from alloys in the same aged states as in Part I. Using mechanical property measurements and microstructural characterization techniques, we attempt to relate the emission sensitivity to these variables to the changes in deformation mechanism. 2. MATERIALS
Two ingots, F69 and F70, were prepared using the technique described in Part I for © Elsevier Sequoia/Printed in The Netherlands
170 TABLE 1 Bulk chemical compositions of A1-Zn-Mg alloys Alloy
Zn (wt.%)
Mg (wt.%)
Cu (wt.%)
Zr (wt.%)
Fe (wt.%)
AI (wt.%)
F69 Top F69 Bottom
5.51 5.48
2.49 2.47
1.59 1.59
---
0.02 0.02
Balance Balance
F70 Top F70 Bottom
5.36 5.54
2.36 2.47
1.48 1.60
0.16 0.15
0.03 0.03
Balance Balance
All other elements were below the detection level of 0.01 wt.%.
alloy F68 [6]. The ingot compositions measured at the top and bottom of each ingot are given in Table 1. The ingots were rolled to 50 mm plate, using procedures identical with those for F68. An electron microprobe was used to seek evidence of microsegregation in both the as-rolled condition and following solution treatment for 2 h at 465 °Co Insignificant levels of segregation were observed in these two alloys, unlike F68. Inclusions were present in alloys F69 and F70 (Fig. 1). Electron microprobe chemical analysis of the inclusions in alloy F69 indicated the presence of two inclusion types: Cu0.4Alo.05Mg0.1vZn0.03, type I; Cu0.21A10.33Mg0.2sZn0.3, type II. The type II inclusions were of similar composition to the S phase, A12MgCu, but with zinc substitution of some aluminium. The inclusions were up to about 10 pm in size and were distributed
I
20pm
!
Fig. 1. Inclusions in undeformed AI-5.5wt.%Zn2.5wt.%Mg-l.6wt.%Cu alloy (F69).
in a dendritic structure with an interdendrite spacing of 50-100 ~zm. Tensile samples of dumb-beU geometry [6] were machined from the rolled plates and solution treated at 465 °C for 2 h and quenched into warm water. In alloy F69 this resulted in a fully recrystallized microstructure with equiaxed grains whose average size (by the line intercept method) was about 230 pm (compared with 420 #m in alloy F68 studied previously). In alloy FT0, recrystallization had been inhibited, as anticipated, by the addition of zirconium and the microstructure was that of a recovered rolled structure, i.e. elongated grains up to 200 pm in length and 50 pm across, containing a subgrain structure with an average subgrain size of about I0 ~m. 3. EXPERIMENTAL METHODS
The solution-treated samples were divided into three groups, and pairs of samples were given one of three types of aging treatment: (a) isothermal aging at 20 °C (natural aging) for up to 30 days; (b) isothermal aging at 120 °C (120 °C aging) for up to 41 days; (c) 24 h at 120 °C plus isothermal aging at 180 °C (double aging) for up to 12 h. The samples were electropolished to remove any oxide layer formed during heat treatment. One of each pair was used for acoustic emission measurement and the other for detailed mechanical property studies, using the techniques and procedures described in Part I. After testing, the fracture surfaces of the samples were examined to characterize the fracture mode. Additionally, transmission electron microscopy samples were cut from the gauge regions to observe the precipitate distribution and the dislocation structures of the deformed material.
171 4. ACOUSTIC EMISSION RESULTS
4.1. Alloy F68 The detailed acoustic emission results for the ternary alloy (F68) are given in Part I [6] but are summarized here for completeness. The behaviour of the quenched material was characterized by the generation of many small emission signals at yield and during the load drops accompanying serrated yielding (the Portevin-Le Chatelier effect) and by the generation of a few large-amplitude signals randomly throughout the test. In contrast, material that had been naturally aged or aged at 120 °C emitted isolated but individually energetic signals near yield and many very energetic signals during plastic deformation. These energetic signals were most intensely emitted from samples that had been aged for 24 h or longer at 120 °C. A second aging treatment at 180 °C caused the disappearance of energetic signals during plastic deformation and the reappearance of small emission signals at yield. 4.2. Alloy F69 The quenched condition generated very similar acoustic emission signals to the ternary alloy. A broad peak in activity composed of many small-amplitude signals was observed at yield (Fig. 2(a)). This was followed by further weak emission associated with serrated yielding during post-yield deformation. In this
region, isolated large-amplitude signals were also observed. The effect of natural aging is shown in Figs. 2(b)-2(d) and again was similar in many ways to that of alloy F68 in that the yield region activity quickly disappeared and highly energetic signals were emitted during plastic deformation. However, in contrast with F68, the naturally aged copper-containing alloy copiously emitted small-amplitude signals during serrated yielding (Fig. 3). The emission tended to be linked with the drops in load accompanying each serration. Aging at 120 °C again resulted in the generation of many extremely energetic signals (Fig. 4). This activity was most intense in the vicinity of peak hardness, similar to alloy F68. Double aging at 180 °C (Fig. 5) resulted in a gradual disappearance of the highly energetic burst emission, the reappearance of yield region emission and, after long aging times, continuous acoustic emission during postyield deformation. A cluster of events was observed during deformation as fracture was initiated and the nominal stress fell. To summarize, the effect of the copper addition seems to be rather small, the most noticeable effect being the more prominent so-called "continuous" emission during postyield deformation in underaged and overaged conditions.
600~-
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Fig. 2. Effect of natural aging (20 °C) on acoustic emission and mechanical properties for F69: (a) quenched; (b) aged for 1 day; (c) aged for 7 days; (d) aged for 30 days.
172 400
~ '~ Q.
.
_ I I
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i
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Fig. 3. A portion of the stress-strain, acoustic emission (AE)-strain data showing correlation between dynamic strain aging and acoustic emission for F69 (aged for 7 days at 20 °C).
400
./Acoustic
Emission Power
20
(a)
1
2
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(b)
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400 (f)
2oo~/j~m O0
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Extens,on (rnm)
Fig. 4. Effect of aging at 120 °C on acoustic emission and mechanical properties for F69: (a) quenched; (b) aged for 6 rain; (c) aged for 24 min; (d) aged for 120 min; (e) aged for 1440 rain; (f) aged for 60 000 min.
4.3. Alloy F70 The addition of 0.16 wt.% Zr to F69 had a great effect on the acoustic emission activity. Whilst a yield peak was still emitted in the quenched condition (Fig. 6(a)), there was very little post-yield activity during serrated yielding. The yield activity decreased with natural aging (Figs. 6(b)-6(d)) but, in marked contrast with alloys F68 and F69, there was
only a small increase in activity during postyield deformation in naturally aged samples. Similarly, for samples aged at 120 °C, the highly energetic signals observed in F68 and F69 near peak hardness were almost totally absent apart from a small cluster of events associated with fracture (Figs. 7(e) and 7(f)). Overaging also had a different effect in F70 from in F68 and F69 (Fig. 8). No yield
173
60¢
¢0
region peak was observed, and the continuous acoustic emission during post-yield deformation of F69 was absent. The sample aged for 24 h at 120 °C plus 12 h at 180 °C was noteworthy because it generated no detectable elastic waves other than from the stress drop of final fracture.
400J" ~ - l~--Acoustlc Emtss,on I 20OF/ / Power I Oo ~ 2 3 O0
(a)
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5. MECHANICAL PROPERTIES
"
20
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.,
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Load-extension data were digitally recorded for every sample condition and analysed to give the nominal stress-strain behaviour. This was then characterized by parameters such as the 0.1% proof stress o0.1, the ultimate tensile stress OUTS, the strain to OUTSand the work-hardening exponent and rate. These data are given in Tables 2 and 3 for each alloy.
3
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0~"- -
(f)
-
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'1 ''~[-'L'J`2 "=
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The effect of natural aging for all three alloys is shown in Fig. 9. It can be seen that b o t h Oo.x and OUTSwere proportional to log(aging time) for all three alloys. In part the effect of the composition difference was to change the grain size and thus to displace the curves so that the strength of F70 was greater than that of F69 which in turn was greater
0
Extension(ram) Fig. 5. Effect of secondary aging at 180 °C on the acoustic emission and mechanical properties for F69 : (a) 24 h at 120 °C only; (b) secondary aging for 1 h; (c) secondary aging for 2 h; (d) secondary aging for 4 h; (e) secondary aging for 8 h; (f) secondary aging for 1 2 h .
600 400 ~/Acoustic 200 ~ . , ....
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00
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7
;
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Extension (mm) Fig. 6. Effect of natural aging (20 °C) on acoustic emission and mechanical properties for F70: (a) quenched; (b) aged for 1 day; (c) aged for 7 days; (d) aged for 30 days.
174 t.O 400 f200 ~ i hA~,i "AcOustic Emission Power
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6
Fig. 7. Effect of aging at 120 ~C on acoustic emission and mechanical properties for F70: (a) quenched; (b) aged for 6 rain; (c) aged for 24 min; (d) aged for 120 rain; (e) aged for 1440 rain; (f) aged for 60 000 rain. than that of F68. Similarly the copper addition did n o t greatly affect the kinetics of hardening at 120 °C (Fig. 10); all three alloys had a similar time dependence, but at each time the strength decreased from F70 to F69 to F68. Double aging caused a loss of strength (Fig. 11). The rate of decrease in alloy F69 was lower, whilst in F70 it was greater, than that of F68.
5.2. Ductility We used the plastic strain to maximum true stress as the ductility indicator. From Tables 2, 3 and Part I it is seen that natural aging in all three alloys tended to reduce ductility, although this effect was less evident in F70. Isothermal aging at 120 °C caused an almost m o n o t o n i c loss of ductility with increasing strength in all three alloys. Double aging at 180 °C had surprisingly little systematic effect on the ductility. 5.3. Work-hardening rates The stress-strain curves were numerically differentiated with respect to strain to determine the work-hardening rate as a function
of plastic strain for all the tests and these are s h o w n in Figs. 12-14. Several effects of aging were observed. The work-hardening rate decreased with plastic strain in all cases to a value of about 1.2 GPa at 5% strain, for all heat treatments and compositions. At small strains (about 1%), natural and 120 °C aging tended to increase the work-hardening rate in all three alloys. At small strains, double aging at 180 °C causes a decrease in work-hardening rate that was most marked in F70 (Fig. 14). At larger strains (about 5%) the work-hardening rate tended to decrease as the peak hardness was approached. The work-hardening exponent n in the relation O"T ---~ k C T n
(where o T is the true stress, eT the true plastic strain and k a constant) is also c o m m o n l y used to describe the work-hardening characteristics of a metal. Values of n at three strains are given in Tables 2 and 3. Whilst it is usually considered that n is independent of strain, the precise recording techniques used here show that the n values usually increase with strain.
175
i
60C
40C
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Ern,ssion Power
(a) 1
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o ~E
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200~ Extension(ram)
-
./
y><.
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Fig. 8. Effect of secondary aging at 180 °C on acoustic emission and mechanical properties for F70: (a) 24 h at 120 °C only; (b) secondary aging for 1 h; (c) sec()ndary aging for 2 h; (d) secondary aging for 4 h; (e) secondary aging for 8 h; (f) secondary aging for 12 h.
lOO
1101
1102
103
1104
105
Ag ing Time (rain)
The values at a strain of about 5% are probably the best to use for comparison with other published data. In all three alloys, natural aging in general decreased n, aging at 120 °C resulted in larger decreases whilst double aging led t o a recovery in the n value. For a given aging condition the
(b)
Fig. 10. Effect of aging at 120 °C on (a) the 0.1% proof stresses and (b) the ultimate tensile stresses of alloy F68 (o), F69 (A) and F70 (u). [ ] indicates premature failure.
+oo[ c~
400
400
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c~ 'E z ~E
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2OO
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100 (b)
u
I I
I A9,ng
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10 (days)
I
100
Fig. 9. Effect of natural aging on (a) the 0.1% proof stresses and (b) the ultimate tensile stresses of alloys F68 (I), F69 (A) and F70 ( i ) .
176 TABLE 2 Mechanical properties: F69
Heat treatment
Stress (MPa) O0.1
Isochronai aging at 20 °C Quenched 119 1 day 155 30 days 249 Isochronal aging Quenched 6min 12min 24min 36 min 60min 120 min 1440mm 6000 min 60000 min
O0.2
Ultimate strain a
Work-hardening rate (GPa) Work-hardening exponent for the following strains for the following strains
OUTS
60.1
61
e5
eO.1
61
65
129 166 259
274 312 385
0.140 0.164 0.110
9.2 12 13.5
2.5 2.7 4.1
1.7 1.1 1.2
0.07 0.07 0.06
0.12 0.10 0.08
0.38 0.21 0.17
at 120 °C 119 129 209 220 211 220 233 246 272 285 293 303 298 308 398 415 428 445 423 445
274 356 344 362 383 383 414 487 485 476
0.140 0.152 0.147 0.152 0.074 0.047 0.119 0.058 0.010 0.026
9.2 16.7 14.2 17.5 17.5 20.0 19.0 25.9 31.0 28.9
2.5 3.2 3.1 3.3 4.3 4.5 4.4 5.0 5.1 3.2
1.7 1.5 1.2 1.3 1.3 1.2 1.1 0.9 --
0.07 0.09 0.07 0.08 0.07 0.09 0.08 0.06 0.05
0.38 0.22 0.19 0.18 0.16 0.13 0.12 0.09 --
--
--
0.12 0.08 0.08 0.08 0.08 0.07 0.07 0.06 0.05 0.04
487 436 422 417 402
0.058 0.031 0.041 0.056 0.045
25.9 18.6 14.7 11.0 13.0
5.0 3.3 3.0 3.0 3.2
0.9 --1.0 1.0
0.06 0.04 0.04 0.03 0.04
0.06 0.04 0.04 0.04 0.05
0.09 --0.10 0.10
Doubleagingat180°C 0h 398 2h 376 4h 358 8h 341 12 h 321
415 382 363 345 329
--
aplastic strain corresponding to maximum true stress. 500
i[:: Z
z ,ooL
•
300
-,---%,_
300
• ..,.
Z 200
g zoo
100
_~ 100
n E
I~ (a)
I 1
[ 10
] 20
Aging Time (b)
0
(b)
I~
L 1
l 10
I 20
Aging Time (h)
Fig. 11. Effect of secondary aging at 180 °C on (a) the 0.1% proof stresses and (b) the ultimate tensile stresses of alloy F68 (e), F69 (A) and F70 (m) in the 120 °C, 24 h aged condition. n v a l u e a t 5% s t r a i n d e c r e a s e d f r o m F 6 8 t o F69 to F70. 6. MICROSTRUCTURAL CHARACTERIZATION
6.1. P r e c i p i t a t e s Transmission electron microscopy examinat i o n o f s a m p l e s o f F 6 9 a n d F 7 0 a g e d a t 1 2 0 °C
s h o w e d t h a t t h e p r e c i p i t a t e size a n d i n t e r particle spacing distributions in these alloys were similar to those in F68. A high density of small precipitates formed within the grains a f t e r 2 4 h a t 1 2 0 °C w i t h a n a r r o w p r e c i p i t a t e ~ free z o n e a d j a c e n t to grain b o u n d a r i e s . The boundaries themselves, including the subgrain b o u n d a r i e s in F70, were d e c o r a t e d with larger
177 TABLE 3 M e c h a n i c a l p r o p e r t i e s : F 70
Heat treatment
Stress (MPa) O0.1
Isochronal aging at 20 °C Quenched 160 1 day 211 7 days 246 30 d a y s 275 Isochronal aging Quenched 6 rain 12 rain 24 r a i n 36 m i n 6 0 rain 1 2 0 rain 1 4 4 0 rain 6000 min 6 0 0 0 0 rain
Ultimate strain a
00.2
Work-hardening rate ( G P a ) Work-hardening exponent for the following strains for the following strains
OUTS eO.1
61
65
60,1
61
65
170 216 260 288
330 358 401 438
0.165 0.145 0.114 0.163
12 7 17 15
2.6 3.1 4.8 4.7
1.8 1.4 1.4 1.3
0.08 0.03 0.08 0.06
0.09 0.09 0.10 0.09
0.32 0.22 0.18 0.16
at 120 °C 160 170 220 232 239 249 263 277 265 297 316 323 336 340 407 424 419 456 426 453
330 382 385 397 392 444 454 508 505 502
0.165 0.176 0.149 0.109 0.099 0.109 0.098 0.081 0.073 0.077
12 15 17 16 19 19 18 28 31 32
2.6 3.7 4.0 4.1 4.2 4.6 5.0 5.3 3.7 3.6
1.8 1.5 1.2 1.2 1.2 1.3 1,2 1,0 0.9 0.7
0.08 0.07 0.08 0.07 0.08 0.07 0.06 0.06 -0.06
0.09 0.09 0.08 0.08 0.08 0.08 0.07 0.06 0.04 0.04
0.32 0.21 0.17 0.16 0.14 0.14 0.13 0.08 0.08 0.06
508 485 447 434 373
0.081 0.061 0.081 0.073 0.057
28 27 15 11 13
5.3 3.8 3.0 3.2 2.7
1,0 0,9 1.0 1.2 1,3
0.06 0.05 0.03 0.03 0.05
0.06 0.05 0.04 0.04 0.05
0.08 0.09 0.10 0.13 0.16
Double aging at 180°C 0 h 407 2 h 406 4 h 358 8 h 345 12 h 280
424 421 366 353 288
aplastic s t r a i n c o r r e s p o n d i n g t o m a x i m u m t r u e stress. 40 35
3o
3O
Q.
~D
~25
"~ 25
o tr
~2o
~ 2o
~v
"o a,
2.15
15
i
10
0 ~uencnea ~
(a)
days
~10 1 day
~_~
I
0.01 Plastic Strain
0.02
(b)
0.01 Plastic Strain
30 days
Quenchedl
1 0.02
(c)
0.01 Plastic Strain
0.02
F i g . 12. E f f e c t o f n a t u r a l aging ( 2 0 ° C ) o n t h e yield region w o r k - h a r d e n i n g r a t e : (a) F 6 8 ; (b) F 6 9 ; ( c ) F 7 0 .
precipitates. Prolonged aging at 120 °C gave precipitate coarsening (Fig. 15). Peel et al. [7] have found that aging at 120 °C for 24 h gives almost entirely 7' precipitates in an alloy
of the F69 composition, rather than a mixture of 7' and 77 as in F68. In contrast, aging at 180 °C for 8 h did produce a different precipitate size distribution
178
~. 151\ \n |1 _ \ \ ~ d o ~
@h
2
d~ys
~oy5 Ib h
0.01 Ptast,¢ Strain
002
24rain
0
0.01 Plastic Strain
0.02
(a)
0
0.01 PLastic Strain
0.02
(b)
0
(c)
Fig. 13. Effect of aging at 120 °C on the yield region work-hardening rate: (a) F68; (b) F69; (c) F70.
30 o Q_ ®
"-" 25
25
0 n~
¢F
~, 20
~ zo
2O n 0
"u
20 15
er
0h
i
Os.
g 0
"
~ (a)
~10
10
5
0
Sh 1 '~
4h
I 0.01 Plastic Stra,n
--
I 002
~
0
I
I
0.01 Plash¢ Strain
(b)
0.02
001 Plastic Strain
I
002
(c)
Fig. 14. Effect of double aging at 180°C on the yield region work-hardening rate: (a) F68; (b) F69; (c) F70.
!OOnm i
Fig. 15. Transmission electron micrograph of fine precipitates in alloy F69 aged at 120 °C for 41 days.
i
Fig. 16. Transmission electron micrograph showing coarse matrix precipitates and precipitate-free zone at grain boundary in alloy F69 aged at 120 °C for 24 h and at 180°C for 8 h.
179
(a)
(b) Fig. 17. (a) Transmission electron micrograph of precipitates in alloy F69 aged at 120 °C for 24 h and at 180 °C for 8 h; (b) transmission electron micrograph of precipitates in alloy F70 aged at 120 °C for 24 h and at 180°C for 8 h. in F69 and F70 compared with that in F68. In F68, rod-shaped precipitates with a precipitate density of 1016 cm -3 were observed. In F69 and F70 there was a similar density of the rod-shaped precipitates but a distribution of smaller spherical precipitates was also present, giving a total precipitate density of 3 × 1016 cm -3 (Figs. 16 and 17).
6.2. Dislocation structures The dislocation structures in F69 and F70 samples were also examined using transmission electron microscopy. F69 was found to be similar to F68, showing uniform dislocation structures in quenched and underaged conditions with a transition to localized slip in bands around peak hardness. The dislocation density within the bands was less in F69 than in F68 (Fig. 18). There was a reversion to uniform deformation after a secondary aging treatment at 180 °C. In contrast, F70 exhibited little heterogeneous slip.
Fig. 18. Transmission electron micrograph of lightly deformed bands in alloy F69 aged at 120 °C for 4 days.
Fig. 19. Scanning electron micrograph of fracture surface of alloy F69 aged at 120 °C for 6 min, showing mainly ductile fracture.
6.3. Fractography Fractography indicated that alloy F69 exhibited a tendency to undergo a transition from ductile dimple fracture in the underaged state (Fig. 19) to a mixed fracture consisting of small areas of intergranular failure linked by regions of ductile dimples when peak aged. The addition of copper reduced the extent of intergranular cracking in peak-aged material but did not totally eliminate it. A combination of intergranular and ductile fracture was also found in F70 alloys aged at 120 °C to peak hardness, with the proportion of intergranular fracture decreasing away from peak hardness. In F70 aged for 24 h at 120 °C, the fracture path appeared to follow the boundaries of the elongated grains. Fracture along strings of inclusions at grain boundaries and fracture along subgrain boundaries were also seen (Fig. 20).
180 that considerable plastic relaxation had accompanied the cracking.
7. DISCUSSION
Fig. 20. Scanning electron micrograph of fracture surface of alloy F70 aged at 120 °C for 24 h, showing intergranular fracture, fracture at inclusions A and subgrain cracking B.
Fig. 21. Scanning electron rnicrograph of a polished gauge region of alloy F69, showing cracks across S phase inclusions. The tensile axis was horizontal.
6.4. Inclusions
Scanning electron microscopy of fracture surface indicated that the majority of inclusions present on the surface had fractured in both F69 and F70. The gauge regions of fractured samples were then sectioned and polished to reveal the inclusions below the fracture surface. It was found (Fig. 21) that almost every inclusion had fractured, the larger ones containing several cracks. The fractures tended to be normal to the tensile axis, indicating mode I failure. The large crackopening displacements (about 1 pm) suggested
It was shown in Part I that during deformation the product of the propagation distance a and velocity v of individual (or a group of n) dislocations must exceed a threshold value (determined by experimental sensitivity) for detectable elastic waves to be emitted. For aluminium alloys, nay >/3.6 × 1 0 - 2 m 2 s -1 for detectable emission to be generated, where n is the number of dislocations moving together. The smallest detectable motion of a single dislocation might thus correspond to the creation of a loop 180 ~zm in radius whose mean radial velocity was 200 m s-1 [1]. Since b o t h the distance of dislocation propagation and the velocity are distributed quantities, only a fraction of the deformation processes of a sample are detectable, and the acoustic emission measurements indicate this to be very dependent on metallurgical variables such as precipitate t y p e and distribution and grain size. In a detailed analysis of results from the ternary alloy in Part I, it was shown that three types of deformation mechanism were able to generate detectable acoustic emission and each depended differently on microstructure. (i) In quenched or very lightly aged conditions, dislocation unpinning from solute atoms either at yield or during plastic flow (dynamic strain aging) generated, in the main, weak signals. (ii) In peak-aged samples the formation of intense slip bands associated with the shearing of fine precipitates was an intense source of individually energetic elastic waves. (iii) In overaged samples, with a wider precipitate spacing, the cooperative motion of groups of dislocations from one set of precipitates to the next could generate many weak signals. The copper-containing alloys could potentially have a fourth emission source: the fracture of S phase inclusions. In the following the acoustic emission characteristics of the basic ternary, coppercontaining quaternary and zirconium-doped quaternary alloys are compared and discussed in terms of the effect of microstructural differences on the ability of these deforma-
181
tion and fracture mechanisms to generate detectable acoustic emission. We shall consider each mechanism in turn.
7.1. Dislocation unpinning Similar yield region acoustic emission activity in the quenched state was observed in all three alloys. In the ternary alloy it was shown that, with the aid of internal stress concentrations (possibly ahead of dislocation pile-ups), a sufficient stress could be reached for individual dislocations to unpin themselves from impurity atoms segregated at dislocation cores. After escape the dislocations could multiply and propagate at a high speed, resulting in a sudden local stress relaxation. It seems reasonable to expect the magnitude of this stress relaxation to be determined by the distance of dislocation propagation, i.e. by the grain size at small strains. In the ternary alloy the grain size was a b o u t 430 pm and signals from single dislocations were predicted to be just detectable. In F69 the grain size was less (about 230 pm) b u t the flow stress and possibly the dislocation velocity were higher, and so we might again expect similar levels of acoustic emission. In alloy F70, however, whilst the grain size was still a b o u t 200 pm, there was also a well-developed subgrain structure with a spacing of a b o u t 10 pm which would severely limit the distance of propagation. In this case a single dislocation, even if it propagated at the unlikely limiting (transverse) wave speed of 3200 m s-1, would have an av product of only 3.2 × 10 -2 m 2 s-1 and so be barely detectable. The generation of the observed levels of acoustic emission in F70 could have three possible explanations. (i) There could be coordinated escape of many dislocations, each propagating over a mean subgrain diameter. (ii) Single dislocations could move in the largest subgrains at the top of the grain size distribution so that the propagation distance is greater than the mean subgrain size. (iii) The subgrains were not strong enough to arrest high speed dislocations, and so they propagated to the grain boundaries. It is n o t possible to determine the contribution of these various mechanisms of slip. It is possible that to some extent all three occurred.
During post-yield deformation of b o t h quenched and aged samples, more serrated yielding and its associated acoustic emission were observed in alloys F69 and F70 than in F68. The generation of acoustic emission after large plastic strains when the mean free path for individual dislocations has been reduced to the dislocation cell size (a few micrometres) can only come from the coordinated motion of many thousands of dislocations. Evidence for such widespread sudden motion is given by the stress-strain curves. Here, serrated yielding is analogous to very low frequency (less than 10 Hz) acoustic emission. Each time that sufficient unpinning occurred to give a detectable stress drop on the stress-strain curve, groups of acoustic emissions were observed. The load cell detected the static change in stress whilst the emission signals were from the micromechanisms of slip that a c c o m m o d a t e d the change. The enhanced activity in F69 and F70 could be related to the higher flow stress of these alloys which would aid the escape of dislocations and cause them to propagate at a high velocity. Whilst this flow stress increase was achieved in part b y a decrease in grain size, this did not limit the acoustic emission activity directly since this was determined by the much smaller dislocation cell size.
7.2. Precipitate shear It was argued in Part I for F68 that, in peak-aged material with a fine distribution of small precipitates, dislocation propagation involved precipitate shearing which in turn led to unstable slip localization in bands and the emission of energetic signals. The results from F69 and F70 support this conclusion as the energetic signals were only detected in peak-aged F69 in which slip bands were found using transmission electron microscopy, whereas F70 always showed a uniform dislocation structure and deformed quietly. The peak-aged F69 alloy showed a somewhat reduced level of activity compared with F68. This change could be brought about if the precipitate strength was greater in the copper-containing quaternary alloy, reducing the tendency for precipitate shearing, or it could be related to the smaller grain size which would reduce the length of individual dislocation pile-ups and the local stress on
182 precipitates. The loss of acoustic emission in FT0 might also be the result of the small grain size. A few signals were detected towards the end of deformation in this alloy, suggesting that only then could sufficient stress be applied to the leading dislocation in a pile-up to shear precipitates. However, these signals could also be related to fracture.
7.3. Orowan looping The results from double-aged F68 reported in Part I were surprising in that detectable acoustic emission was recorded during yield although the slip distance might be expected to be limited to the interparticle spacing kp. It was postulated that the acoustic emission was due to the sudden growth of dislocation loops that had been pinned b y an inner set of precipitates lying at a radius r to a set lying at a radius r + kp. The detection of loop growth may then be possible provided that the loop radius (which is limited by the grain size) and interparticle spacing are large. Both the grain size and the interparticle spacing were less in alloy F69 than those of F68 b u t this appears to have had only a small effect on elastic wave generation at yield. This could perhaps result from the larger yield Stress of F69 which could drive free dislocations at a higher velocity. Alloy F70, whilst having a similar interparticle spacing to F69, had a much smaller effective grain size (10 #m compared with 230 pm). Thus, the area swept out when a loop expanded from one set of particles to the next would have been reduced b y a factor of a b o u t 20, and this seems sufficient to ensure no emission of detectable elastic waves.
7.4. Inclusions There is good evidence that inclusion fracture (and possibly decohesion) is an important source of acoustic emission in commercial grade aluminium alloys [2, 4, 5]. Alloys F69 and F70 b o t h contained large numbers of S phase inclusions, but the acoustic emission results provide little evidence that they are an important source of detectable acoustic emission. Indeed, the overaged samples of alloy F70 generated little or no acoustic emission at all, suggesting that inclusion cracking is a slower process in the S phase inclusions. This is supported by recent results of Cousland and Scala [ 5 ] who have shown that the acoustic
emission from inclusion cracking in aluminium alloys is predominantly from brittle iron-containing inclusions. Let us try to estimate the emission amplitudes that might be generated by S phase inclusion fracture. It has been shown [8] that, for a detectable signal to be emitted by a mode I microcrack,
t 3C1Eh Uminl ~2
acVcl i 2 ~ ~
J
where ac is the crack radius (assuming a circular microcrack), vc the radial crack speed, C1 the longitudinal wave speed, E Young's modulus, h the crack-to-detector separation, UmL~ the smallest detectable displacement, v Poisson's ratio and o the applied stress of the instant of fracture. Using, for aluminium, C1 = 6400 m s-1, E = 80 GPa and v = 0.3 and assuming Umm = 10 -14 m, h = 40 m m and a = 400 MPa (for peak-aged material) gives the condition that acv¢~2 > 1.45 × 10 -~ m 3/2 s-1/2. Using a value for ac of a b o u t 3 #m (representing one of the larger microcracks), then the crack growth speed would only need to be more than 25 m s-1 for a detectable signal to be released. Let us assume a range of inclusion sizes with 0.5 pm < ac < 5 pm and a range of speeds with 10 m s-1 < v~ < 1000 m s-1. Then the model predicts a range of acoustic emission amplitudes given by 10-16 m ~ Umm < 10-12 m About half this range exceeds the detection threshold of 10-14 m. Thus it is concluded that S phase inclusion fracture is a possible weak emission source, but only when the
larger inclusions crack at moderate to high velocities. The absence of emission during the period of deformation when cracking of the majority of the inclusions was expected to be occurring [9] (at large plastic strains) suggests that crack growth speeds are generally well below 1000 m s -1. The stronger emission from iron-rich inclusions in the materials studied by Scala and Cousland [4] and by Cousland and Scala [5] is consistent with a larger inclusion size (1-14 pm) and possibly a higher fracture speed. Their work also supports the notion of the critical a~v~~2 value for detectable signals of a b o u t 3 × 10 -4 m 3/2 s -1/2. When material
183
with a smaller inclusion size (1-6 pm) was tested, very little emission was observed.
iron-rich inclusions in commercial alloys could well be a more important emission source.
8. CONCLUSIONS
ACKNOWLEDGMENTS
The acoustic emission activity of quaternary A1-5.5wt.%Zn-2.5wt.%Mg-l.6wt.%Cu alloy has been shown to be critically dependent on composition and heat treatment. The origin of these effects has been attributed to changes in deformation mechanism. The following facts have emerged from this study. (1) The addition of copper to the ternary alloy increases the weak emission associated with dynamic strain aging. The effect is greatest for small-grain-size (high flow stress) samples. (2) The most energetic emission source in the base ternary alloy, that associated with strain localization in the peak-aged state, is only slightly reduced by the addition of copper alone. (3) The addition of zirconium which greatly reduces the grain size causes almost complete disappearance of strain localization and the associated emission. It also leads to the disappearance of the emission associated with Orowan looping. Both effects are consistent with the restriction of the movement of dislocations by grain boundaries. (4) Fracture of the S phase inclusions appears to be a small-amplitude emission source, although other work has shown that
We wish to express our gratitude to D. Stockham-Jones for experimental assistance and Dr. J. Hudson for discussion of this work. The study was funded by the Royal Aircraft Establishment, Farnborough.
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