microanalysis of age-hardening precipitates in aluminium alloys

km merull. Vol. 30. pp. 1389 to t 395. I982 Oooi-6160 92 071389-07503.00 0

Printed in Great Britain. -\tl n&s resrrwd Copyright Q 1952 Pergamon Press Ltd

MICROANALYSIS OF AGE-HARDENING PRECIPITATES IN ALU~INIU~ ALLOYS

J. P. LYNCH’?, L >I. BROWN’ and $1. H. JACOBS’

‘Cavendish Laboratory, Madingley Road, Cambridge and ‘Tube Investments Research Laboratories. Hinxton Hall, Saffron Walden. Essex. England

(Received 4 August 1981: in rr&edform I1 iVoremher)

Abstract-The microdiffraction and X-ray emission spectroscopy facilities of a field emission gun STEM have been used to study precipitates formed on aging Al-l.16 wt?; (MgaSi) and Al-l.16 wt”,; (Mg,Ge) at 25O’C after solution treatment. In the former. a new intermediate precipitate structure, ,9” (a = 0.30 nm, b = 0.33 nm, c = 0.40nm; z = /I = 90’, ; = 71’: (oO1],..i,‘[oO1],,,,i,, [loo], A [lOO],,,,i, = 5’). formed after 3 h, has been observed and a previously suggested structure of hexagonal symmetry, formed after 5 h, has been confirmed. Magnesium concentrations in precipitates formed after 5 h in both alloys have been found to be substantially tower, relative to the other alloying element, than in the equilibrium precipitate (the atomic ratios observed being Mgc,&i and Mg,.,,Gef. It is suggested that. in the very early stages, precipitation is controlled by only one, the least soluble, of the two alloying elements.

Rkumk-Nous avons etudii les precipites formis lors du vieillissement a 25O’C d’alliages Al-l, 169; (Mg,Si) et Al-l, 16% (Mg,Ge) en poids. aprbs mise en solution, grace aux possibilitts de microdiffraction et de spectroscopic d&mission de rayons X dun microscope Q balayage en transmission equip6 d’un canon ii emission de champ. Dans le premier alliage, nous avons observe la formation dun nouveau precipite intermidiaire F (a = 0.3Onm, b = 0.33 nm, c = 0.40nm; z = fi = 90’. 7 = 71’: C~llr//EPollm,t,ice~ EWi7” n [lOO],,,ri,, = 5’) apris trois heures, et nous avons confirme la formation dune structure haxagonale aprts cinq haures. Dans les deux alliages, la concentration en magnesium des prtcipites form& aprts cinq heures ttait nettement inferieure. par rapport a l’autre Clement d’alliage, a celle des precipites a I’iquilibte (les rapports atomiques observes &ant tgaux a Mg,,,,Si et Mg,.,,Ge). Nous pensons que, dans les tout premiers stades. la precipitation n’est controlee que par un seul (le moins soluble) des deux tliments d’alliages.

Zusammenfassung-Mit den Moglichkeiten eines Rasterdurchstrahlungselektronenmikrosko~s mit Fel- demissionskatode in der Mikrobeugung und der Rontgenemissionsspektroskopie wurden Ausscheidun- gen untersucht, die sich bei der Auslagerung von AI-l.16 Gew.-9; (MgaSi) und AI-l,16 Gew.-p;, (Mg,Ge) bei 250 C nach einer Losungsbehandlung bilden. In der ersten Legierung bildeten sich nach 3 Stunden Ausscheidungen mit einer neuen Zwischenstruktur, j3” (a = 0,30 nm, b = 0,33 nm, c = 0,40 nm; 2 = /I = 90’. 7 = 71”: [~l~~*~/[~l]~~~i~, [lOO]a” A [lOO],,z,i, = 5’). AuBerdem konnte das Auftreten einer frtiher vorg~ch~agenen hexagonalen Struktur, die sich nach 5 Stunden bildete, bestgtigt werden. Es wurde gefunden, da13 die Magn~iumkonzentrationen in Ausscheidungen, die sich nach 5 Stunden bildeten, in beiden Legierungen betrlchtlich geringer war, vergiichen mit den anderen Legierungeselementen. als in der Gleichgewichtsausscheidung (die beobachteten Zusammensetzungen waren Mg0.44Si und Mg,.,,Ge). Es wird vorgeschlagen, daB die Ausscheidungsbildung zu Anfang durch nut eines, nlmlich das von beiden am geringsten ltisliche Legierungselement bestimmt wird.

1. INTRODUCTION

Alloys in the Al-Mg-Si system possess good mechan- ical properties, based on their age hardening at ele- vated temperatures after solution treatment. The microstructures developed during the aging process have been widely studied in attempts to gain insights into the hardening process.

The morphology of the precipitation sequence during aging of alloys near the ‘pseudo-binary’ Al-(Mg,Si) section is known from X-ray diffraction [l-7] and electron microscopy [8-l 11 studies. Needle-like G.P. zones are first formed which grow into rod-shaped precipitates, a few nm in diameter

- t Present address: Institut Francais du Petrole, BP.31 1,

92 Rueil IMalmaison. France.

and a few tens of nm in length, oriented parallel to (100) matrix directions. In the final stages, square platelets are formed about 10 nm in thickness and several hundred nm in width. Diffraction studies have shown these platelets to be of the equiIibrium fi phase

(Mg$Si) C121. Early X-ray diffraction studies [4,6] showed that

the intermediate, rod-like, precipitates have the periodicity of the matrix along the rod length. Selected area electron diffraction (SAED) studies confirmed that their structure is not that of Mg,Si. The suggested structures for this p phase have been many. Thomas [ 1 l] obtained SAED patterns from precipitates in high purity Al-l.53 wto/:, (Mg,Si) which he attributed to an f.c.c. lattice with parameter 0.642 nm. Wahi and von Heimendahl[13], using X-ray difirac-

1389

1395 LYNCH et al.: PRECIPITATES IN Al ALLOYS

tion. reported the occurrence of a second intermediate phase. /3”, with a monoclinic structure a = 6 = 0.616 nm, c = 0.71 nm, x = p = 90’. 7 = 82”. Pashley et a[. [S], in a study of Al-l.2 wt?: (Mg,Si), also obtained discrete diffraction spots from precipitates at this stage and, although unable to give a full interpretation, concluded that the p’ phase was not a cubic structure.

Jacobs [t2] employing the same alloy, proposed a hexagonal j3’ structure, coherent with the lattice along its c-axis (the long axis of the rod): u = b = 0.705 nm, f = 0.405 nm, x I p = 90”, y = 125”;

(~l)~‘//(~l)~~,i~, Ciool!F//tr fOlm~tiir- This model for the structure was built up from a

large number of SAED patterns taken from rods lying perpendicular to the electron beam. No confirmation of the pattern expected from the (001) zone axis could be obtained.

Recently, the composition of the /3’ precipitates has also been brought into question. An electron microscope study of precipitation in a commercial Al-9%Si alloy [ 141 showed similar rod-like precipitates to occur, although the magnesium content of this alloy was less than 0.4%. It was suggested that a composition Mg$i for these precipitates is unlikely.

Problems in studying the structure and composition of the intermediate precipitates arise from their size and shape. SAED is limited in spatial resolution to about 1 pm [lj], so that a large volume of matrix, and many particles of several orientations. must be analysed SimultaneousIy. The STEM [16], providing such information from areas a few nm in diameter, seems eminently suited for such studies. This paper presents results from microdiffraction and X-ray emission spectroscopy of aged Al-l.16 wt% (Mg,Si) and the apparently analogous alloy Al-l.16 wt% (MglGe). The STEM employed was a VG (Micro- scopes) HB5 equipped with field emission gun.

2. SPECIMEN PREPARATION

Microscope specimens were obtained from 10pm thick foils of the alloys by punching out 3 mm discs. After heat treatment the specimens were jet polished electrochemically. The thinning solution used was 20% perchloric acid/go% ethanol. A current density of lOAcm-‘, potential of 60 V, and temperature of -2O’C to - 15°C were maintained AI-(Mg,Ge) specimens were argon-ion milled to remove electron opaque surface deposits.

The work of Chatterjee and Entwistle[17] suggested the loss of magnesium to the surface of such alloys on heating in air. In order to examine the extent of this effect, 0.5 mm thick specimens of the Al-IMg-Si alloy were heated to 560°C for OS, 1 and 2.5 h in air and for 2.5 h in argon at low pressure. Cross-sections cut through the slabs were polished flat to within 0.5 pm using diamond paste, and exam- ined using a scanning microprobe analyser [18]. Traces were taken of the Mg (K,) X-ray intensity pro-

0.75

I”

32 0.5 tonter

f sample

0.2s -*- 2.5 tlrs f

--- bulk v0lua

ij 0.6F----_____.__ ______

Fig. 1. Mg concentration w distance from surface for samples treated (a) in air for 0.1, 1 and 2.5 h, (b) in low

pressure argon for 2.5 h.

duced along a line scan perpendicular to the surface exposed during heating. Calibration was performed using bulk, pure element samples and the ZAF correction procedure [19] apphed to produce plots of the variation of magnesium concentration with distance (Fig. 1).

Although the ZAF corrections must be treated with care closer than about 5 pm to the edge of a specimen. a large Mg depletion towards the surface is evident in the air treated samples, extending over more than 100 /cm after 1 h. The variation of concentration in the sample treated in argon is weaker and may be due to errors introduced by the correction technique or changes in the collection efficiency due to variation of orientation of the specimen, with respect to the detector, during the scan. Electron diffraction studies of the stripped oxide showed it to consist of small crystals of MgO-A1203 spinel. The loss of magnesium is almost certainly due to a preferential oxidation [20].

Heat treatment in air will evidently seriously de- plete a lo@rn thick specimen of its magnesium. The solution treatment was therefore carried out in sealed glass ampules at a pressure of lo-‘torr of argon. After heating for 1 h at 560°C the containers were broken on water quenching. Specimens were then aged at 250°C in a furnace which was evacuated to better than 0.1 torr. Aging times of 3 h and 5 h were used.

3. MICRODIFFRACTION RESULTS

Precipitates in Al-l.16 wtyti (Mg,Si) aged for 3 h are rod-like, up to 3OOnm in length and about 4nm in diameter (Fig. 2a). It is not possible to obtain SAED patterns showing distinct precipitate reffec- tions, although streaks parallel to (100) matrix directions are observed (Fig. 2b). If these streaks are inter- preted as due to sheets of intensity in reciprocal space, they show the periodicity of the rod structure to be

LYKCH er &I PRECIPfTATES

(a)

Fig. 2. (a) Image and (b) SAED pattern of A-1.16wtY; (Mg,Si) after 3 h aging.

well defined, and equal to that of the matrix. parallel to the rod direction, but poorly defined perpendicular to it.

Microdiffraction patterns in (001) matrix orientation from rods lying perpendicular to the beam confirm that this is the case [Fig. 3(a)]. Diffraction discs may be seen lying close to the forbidden ma&i.. reflec- tions (010). (030), etc. Because of the large variation of intensity in the diffraction discs (presumably a result of strain) their positions cannot be measured to great accuracy in any individual pattern. Averaging over many patterns improves the accuracy and is a technique discussed below.

Precipitates oriented parallel to the beam (Fig. 3(b)] give rise to microdiffraction patterns with distorted hexagonal symmetry. Again intensity variations in the diffraction discs present problems in ac- curately measuring the lattice parameters. In order to overcome this difficulty, two basis vectors a and b were defined having angles 6, and Bb were calculated for all discs separated by one basis vector. and the

process repeated for all patterns. Arithmetical means and standard deviations were then calculated from the sets of values for each of the four parameters. This is simply the two-dimensional equivalent of measur-

ing the spacing of diffraction spots by taking rhe tength of a row of spots and dividing by the number of spots in the row. In this case several rows in several diffraction patterns were used for both directions. Weak, or badly distorted. discs were ignored. so that the ends of rows are well defined, in order to reduce errors. The standard deviations showed the errors in f a [ and /b [ to be less than 3”: and those in 0, and 0, to be I*-,.

The structure /I”; derived from this analysis. is monoclinic, although differing from that reported by Wahi and von Heimendahl. The values calculated were: u = 0.30 5 0.008 nm, b = 0.33 5 0.008 nm, c = 0.40 2 O.@Mnm; x = fi = 90” + 1’. 7 = 71” * 2”. with the orientation relationship [OOlld [OO1]m.r,,X; [lOOIT A [lOO],,,,i, = 5’ k 1” (Fig. 4).

In specimens aged for 5 h the precipitates are again rod-like. although longer ( - 500 nm) and more widely spaced. Figure 5 shows microdiffraction patterns from these precipitates which confirm the p’ structure suggested by Jacobs. Those from rods lying perpendicular to the beam are similar to the SAED patterns he obtained in the same orientations. In addition, it is possible to obtain microdiffraction patterns from rods lying parallel to the beam, showing clearly the hexagonal symmetry. A similar computational analysis to that described above , gtves : a = b = 0.708 2 0.005 nm, c = 0.405 f 0.004 nm; z = B = 90” * 1’.

(b) Fig. 3. Microdiffraction patterns of rod-like precipitates formed by 3 h aging, (a) beam perpendicular to rods, (b) beam parallel to rods. Objective semi-angle. r, = mrad,

collector semi-angle. Y, = 1.7 mrad.

LYXCH rt of.: PRECiPiTAfES IX ;ti ALLOYS

0 t

n 0 0

.-

"'A I

0 0

0 3

I 0 , 0

3 ’ 0 I

*O"Al

0

3 l

0

3

. 0

3 l

Fig. 4. Model of (OOI),,.. reciprocal lattice section.

7 = 60” f 22 with [CQl],&[llO],,,. [lCC&.,” [tiOJ,,,Si.i.~ which agrees within experimental error

with Jacobs’ /3’ structure. It is possible to calculate the strain involved in a

tr~sformation 8” -L j3’, and since the c-axis is undis- turbed, this is a two dimensional pro&m. ff the point

positions (Xi . Xf of an array are to be superimposed on those of a second. (XbYi), the transformation matrix may be calculated by selecting any two pairs of points and solving the matrix equation:

The result of the cahzulation is the strain matrix Cij

-0.012 + 0.012 0.168 +- 0.012

0.197 i 0.012 0.009 + 0.012

which transforms the p[OOl] pattern (marked in Fig. 6). Within the limits of accuracy of the method, the diagonal components may be negfected so the matrix represents a pure shear of -c 18% with axes of princi- pal strain at 45‘ to the [lOO] matrix direction. An ordering reaction in the (001) plane is then required to complete the transformation j?“-/?‘. It should be noted that this calculation assumes such a transformation takes place. It is also possibte that nuclei of the /$’ (and the /3) phase co-exist and compete with the fi phase [21]. The misfit of lg% is weif above the estimated critical value of 5:*; required to produce dislo- cations [24]. The strain required is a pure shear, so it can be accomplished by a mixture of dislocation nu- cleation and short-range diffusion, hence it should be easy at the aging temperature. No external dislocation loops will be Gsible after the transformation, because vacancy loops and interstitial loops will shrink by short-range diffusion of vacancies from the former to the latter.

.

*?A1 . l

l * .

.

(bt

Fig. 5. Microdiffraction patterns of rod-like precipitates formed by 5 h aging. ia) beam parallel to rod: z, = 1.7 mrad. z, = 0.7 mrad. (b) beam perpendicutar to rod. iO31) matrix orientation (to be compared with Fig. 6(b)

of Rrf [ 131~: x0 = 5 mrad. zC = 0.7 mrad.

. . . . ‘ I .

I . .

. l . :"A1 l

.

l * . I

l *

. f

l

-- .

l *

.

Fig. 6. Possible transformation of (OOI), onto (OOl),.

LYNCH v! of.: PRECIPTTATES 13 .4i ;\LLOYS 133

X-ta); spectra were obtained from both ailoys. aged under identical vacuum conditions. Because of the surrounding matrix, the possible incorporation of aluminium into the precipitate structure cannot be unambiguously evaluated by this method, as is the case also for the vacancy content. ft was felt that the most reliable procedure was to compare spectra from overaged precipitates. whose composition was assumed to be Mg:Si. with spectra from intermediate precipitates.

Figure 7fa) shows 3 spectrum from an MgzSi precipitate in an overaged specimen. Figure 7(b) shows a spectrum from a ,Y precipitate. It is immediately obvious that the magnesium:siIicon ratio is much less for the latter than it is for the former. The small peak due to chlorine in Fig. 7(b) is attributable to products of the polishing process.

Quantitative interpretation of the spectra can be made as follows: a reference spectrum from the matrix for which the Mg and Si peaks are absent provides the shape oi the background near the peaks. This is subtracted. and the ratio of counts I,, Is, for the overaged precipitate in Fig. 7ia) is found to be 1.75 2 0.05. The mass ratio C,,‘Cs, in z/fg,Si is 1.73. Thus to within the experimental error

C,,, C,, = k ~,,‘I,, i’)

where the constant k is not appreciably different from unity.

The count ratio similarly calculated for the precipitate spectrum in Fig. 7tb) is 0.39 k 0.01: this gives an equaf mass ratio and an atomic ratio of 0.44. Thus the composition of the precipitate at this stage is

Wg0.&. A feature of these calculations is that they differ

appreciably from the results which would be obtained by the use of mineral standards. The microscope and detector have been calibrated and the resultant k in equation (2) is 1.55 [Xl. This would yield a composition Mg,.,Si for the averaged precipitate. and M&&i for the intermediate precipitate. However. the use of mineral standards in this case is question- able because of the very close proximity of the three Kr lines from Mg, Ai. and Si. Reference to the paper by Goldstein et ul. [.?S] shows that the absorption of the Si radiation by Mg and Al cannot be neglected except for films thinner than a few tens of nanometres. Likewise, secondary fluorescence of the Mg radiation will be considerable. Both effects will increase the :ifg:Si count-rate ratio. It is extremely difficult to make a quantitative estimate of the corrections because of the highly inhomogeneous structure of the precipitate-containing thin film. If one assumes a homogeneous distribution of Si in Al. the absorption correction factor for the Si:At ratio may be estimated to be 0.7 for a film 500 nm thick.? Likewise. the

t Based on Fig. 6 of Goldstein cr uf. f3], with attenu- ation from Reed’s book [35]: X-ray take-off angle 2 = 15

Fig. 7. X-ray emission spectra of prrcipitaws prs%ted by (a) overagins Ai_iMg$it. fbt 5 h aping of Al -(Mg:Sir. IS)

5 h aging of Al-(SlplGek

139-t LY%CH rr ~1.: PRECIPITATES IN .U ALLOYS

fluorescence of the Mg by the AI radiation may be estimated according to Nockolds et &[27] and the tIuorescence enhancement ratio is 0.01, much tess sig- nificant than the absorption correction because of the Row fluorescent yield of these light elements. These factors will affect the derived Mg:Si mass ratio, reducing it from 3.1 to -2.2 It seems clear that the corrections are very important, and the use of an overaged precipitate as a standard is tikely to produce more accurate results. However, even in this case the derived composition will depend somewhat upon precipitate size and geometry, so the numerical compo- sitions quoted here must be treated with caution.

Parallel experiments were carried out in AI- Mg-Ge, which produces a similar precipitate morphology. A spectrum from one of the precipitates [Fig. 7(c)] shows an atomic ratio tclg:Ge of 0.29 & 0.07 (the larger errors arising from low count- ing rates in this case). The copper peak here is due to fluorescence from parts of the microscope and specimen holder and is typical of such spectra.

Thus in both Al-Mg-Si and Al-Mg-Ge the intermediate precipitates are magnesium-poor. One possi- bility which cannot be allowed for is that the precipitate is surrounded by a ‘cloud’ of soiute. In that case the elemental ratio detected will include some contri- bution from the surrounding volume (depending on the spatial resolution with which spectra may be obtained). Because the probe is much smaller than the precipitates. the effect of the external solute distribution is small, and the composition ratios quoted here are likely to be quite accurate.

5. DISCUSSION

It does not seem possible to relate the X-ray spectroscopy and microdiffraction results by proposing a /?’ unit cell. As has been noted, the observed atomic ratios may be characteristic not of the /?’ phase itself but of the region containing the precipitate. In addition, the specimens were sufficiently thick for the diffracted intensity variations to be affected by dyna- mical effects [lj]. The result show. however. that both the structure and the composition of the /?’ precipitates are not of the equilibrium form. The alloy. in the early stages of precipitation, does not behave in a pseudo-binary manner: silicon. or germanium, atoms migrate to the precipitation site at an earlier stage in aging than the magnesium atoms. This is in agree- ment with the result of Kovacs er al. [23] who suggested, from resistivity measurements. that in the very early stages of precipitation, clusters are formed almost entirely of vacancies and silicon atoms.

It is interesting to note that the hardness maximum occurs at this early stage in the precipitation sequence, and that addition of silicon in excess of the 2Mg:Si ratio is known to increase the maximum hardness [ 131.

Evidently, an extended study of the compositional variations in both alloys. covering the whole of the

precipitation sequence, would be of interest. In par- ticular. elemental mapping techniques [lg] might reveal the extent and rates of segregation of the two aIIo_ying elements. Further diffraction studies. of the Al-%fg-Gs system and of the transformations involved in precipitate growth in Al+lg2Si) would also be of value.

The main conclusion to be drawn from these obser- vations is that precipitation is controlled by the supersaturation of the least soluble component of the equilibrium precipitate. The precipitation reaction is, in its early stages, similar to that in ALSi and it is not necessary to form Mg-Si bonds to nucleate precipitates.

The unit cell dimensions for the fi” and p’ structures proposed here are refinements on an earlier version of this work (ZS].

As a final comment. it should be emphasized that STEM utilising a field emission gun allows for the first time the combination of micro diffraction and compositional analysis at sufficient spatiat resolution to elucidate microstructure at the nanometre level.

Ackno~~~edgen~rnts-J.P.L. is grateful to the Science Research council and to Tube Investments Ltd. for a CASE award. The S.T.E.Xf. project is supported by the S.R.C. and by the liniversity of Cambridge. It is a pleasure to acknowledge instrumental help from Dr S. J. Pennycook and Mr L. G. P. Jones. %‘e are grateful to the referee for helpful criticism of the paper.

REFERENCES

I. A. Guinier. Solid St. Ph~s. 9, 293 (1959). 2. A. H. Geisler and J. K. Hill. .Ictn Cr!-jml/ogr. I. 283

(1948). 3. A. Guinier and H. Lambot. Comprvs Rrnnu 227. 74

(1948). 4. A. Lutts, Acrcr merdl. 9, 577 (1961). 5. H. Lambot, Rev. hferuil. 47. 709 (1950). 6. A. Lutts and H. Lambot, Rer. Merull. M. 775 (1957). 7. A. Guinier, rlcrcd Cr.wdogr. 5, 121 (1952). 8. D. W. Pashley. J. Insr. Mrruls 94, 41 11966). 9. R. Castaing and A. Guinier, Compres Rmh 228. 1146

i 1949). 10. A. Saulnier and P. Mirand, Rec. .Lferdf. 57.91 (1960). 11. G. Thomas, i. Insr. Merds 90, S7 (196162). 12. M. H. Jacobs. Phil. Mug. 26. I (19721. 13. R. P. Wahi and IM. van Heimendahl. Physica .urarus

solidi. (I 25, 607 (197-1). II. M. Kaczorowski. M. W. Grabski. J. Sawicki and P.

Murza-Mucha. J. Marrr. Sci. 14, 2X1 f1979). 15. P. B. Hirsch. A. Howie. R. B. Nicholson. D. W. Pashley

and M. J. Q’helan, Elecrron .~ficroscop of Thin Cr~s- ruls. Kreiger, New York (1977).

16. A. V. Crewe. 9th Int. Gong. on El. Mic. Toronto, 1979, 1, 197. See also L. 111. Brown. (1981 b 1. Phys. f II, 1

17. D. K. Chatterjee and K. M. Entwistle. f. Insr. Merats. 101, 53 (1973).

18. Xl. H. Jacobs. fnsr. Ph~s. ConJ Ser. 36.333 (1977). 19. D. B. Holt, ;\;1. D. Xfuir, P. R. Grant and I. M.

Boswarva. Qwnrirurice S.E..Cf. Academic Press. New York (197-t).

20. M. H. Jacobs. J. Microsc. 99. 165 ( 197% 21. P. G. Shewmon, Tran.~~~;rrrtarioris in Jfrrak. McGraw

Hill. New York (19691.

LYXH er ~1.: PRECIPITATES fN .A1 ALLOYS 1395

” S 4. Collett. L. hf. Brown and M. H. Jacobs. in Quan- __. t. S. J. B. Reed, S. E. .U. Confi. [ i.TR.1. Chici;yo. 1979. rirurire Muzroanalysis ti,ith Hrgh Spurial Resolurion, Vol. I. p. 315. .CfUnchrjfU 1981. Metals Society. 26. S. J. 8. Reed, Ekcrron Microprobe .-tnulysis. Cambridge

23. I. Kovacs. Acru metali. 20, 975 (1972). University Press 1193L 34. L. ZiI. Brown and G. R. Woolhousc, Phil. Msg. ti, 319 27. C. Nockolds. M. J. Sasir. G. Cliff. and G. \Y Lorimer.

(1970). EMAG 1979 fnsr. Ph.kcs Co$ Series 52. p. 417. 25. J. I. Goldstein. J. L. Costtey. G. W. Lorimcr and 28. J. P. Lynch. L. iv%. Brown and M. H. Jacobs. EMAG

1979. fnsr. Physics Co@ Series 52. p. 539.

microanalysis of age-hardening precipitates in aluminium alloys

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