Stress corrosion susceptibility of AlZnMg weldments: Microstructural effects

Scripta

METAI,LURGICA

Vol.

23, pp. 2 0 9 ] - 2 0 9 6 , 1089 P r i n t e d in t h e U . S . A .

Pergamon P r e s s p [ c All rights reserved

STRESS CORROSION SUSCEPTIBILITY OF Al-Zn-Mg WELDMENTS: MICROSTRUCTURAL EFFECTS

C. GARCIA-CORDOVILLA 1 , E. LOUIS 1,2, A. PAMIES1 , L. CABALLERO3 , V. SANCHEZ- GALVEZ3 and M. ELICES3. 1. Industria Espal~ola del Aluminio. Centro de Investigaci6n y Desarrollo. Apartado 25. 03080-Alicante. Spain. 2. Departamento de Fisica Aplicada. Universidad de Alicante. Apartado 99. 03080-Alicante. Spain. 3. Departamento de Ciencia de Materiales. E.T.S. Ingenieros de Caminos. Ciudad Universilaria. 28040-Madrid. Spain.

(Received July 31, ]989) (Revised September 29, 1989) Introduction The most severe limitation to a full exploitation of weldable AI-Zn-Mg alloys is the occurrence of Iocalised corrosion in weldments [1-3]. Cracking is usually initiated at the weld-toe and propagales by stress corrosion through a zone, adjacent to the weld bead, commonly called "white zone" (WZ) [4-6], due to its response to etching in nitric acid [6]. Besides the current efforts addressed to avoid (or delay) crack initiation [4,7], a considerable amount of work is being devoted to identify and improve the understanding of the properties of the WZ, responsible for its higher sensitivity to stress corrosion cracking [6, 8-11]. In this context, the work of Schmiedel and Gruhl [8] is particularly revealing. These authors produced massive samples having a microstructure very similar to that of the WZ. This approach to the problem is particularly useful because, although other revelant characteristics of the WZ (such as compositional gradients) are left out, it allows to investigate the macroscopic properties of such an small region (0.41.0 mm wide) [6]. The emphasis of their work [8] was placed on the study of the influence of the conditions of the thermal cycle used to reproduce the microstructure of the WZ, on its susceptibility to stress corrosion cracking, while the latter was evaluated onl~ by means of standard constant load tests. In this work we follow the Schmiedel and Gruhl procedure [8] to investigate stress corrosion cracking in the WZ. We did not vary the conditions of the thermal cycle: instead we tried to approach as much as possible the heating and cooling rates attained during welding, in the region where the WZ is produced. On the other hand, attention was concentrated upon the evaluation of the stress corrosion susceptibility. To this end a method not suilable to routine testing, but very powerful in investigating the nature and mechanisms of failure, was used: namely, the Slow Strain Rate Test (SSRT) [12]. Moreover, crack propagation will be studied by means of double beam wedge-loaded specimens (DBWL) testing. Our results confirm those obtained by Schmiedel and Gruhl [8] in the sense that WZ-like (WZL) samples are much more sensitive to stress corrosion cracking than the parent alloy in the T651 temper. In particular, crack growth rates were found to be around 2x10 ~6 m s-1, a figure at least one thousand times that for T651 material. Materials and Experimental

Procedures

The material used in this investigation was supplied in the form of plates 10 and 30 mm thick and in the T651 temper. Atomic absorption gave the following composition in weight per cent: Zn 5.01, Mg 2.44, Cu 0.12, Cr 0.17, Zr 0.13, Mn 0.29, Fe 0.23, Si 0.11, Ti 0.05. MIG welding was carried out by means of an automated equipmenl on plate samples with V-shaped joining edges. The main welding parameters were: 240 A, 30 Volts, 47 cm min -1 traverse rate, and welding wire (1.6 mm of diameter) of alloy AA 5356; the number of passes was 3+1. To determine the characteristics of the thermal cycle to which the heat affected zone was subjected along the welding process, several thermocouples of the type ChromeI-Alumel (0.1 mm of diameter) were placed at different distances from the edges of the samples. The exact position of the thermocouples was checked, after welding, by means of optical microscopy. The maximum temperatures attained in the region of the weld bead adjacent to the interphase were in the range 640-690 C, whereas in the other side of the interphase they were in the range 580-625 C. The temperature in the region close to the so called white zone was determined by averaging values oblained in different welding processes, Io be around 600 C. It is noted that the heating rate is very high, around 100 C s 1, whereas lhe cooling rate is much smaller, being lower than 20 C s -1 in the range 600-300 C. To reproduce the m~crostructure of the white zone in weldmenls, we have followed the procedure of Schmiedel and Gruhl [8]. This consists in introducing the samples in a massive copper block with an orifice, previously heated to a

2091 0036-9748/89 $3.00 + .00 Copyright (c) 1989 Pergamon Press plc

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given lemperature; once a chosen temperature was attained, samples were taken out and quenched in different media. In this work we have used samples of dimensions 150x30x15 mm (LxSTxLT directions). Heating and cooling rates were around 3 C s"1 and 7.5 C s"1 (in the range 600-300 C) respectively; the latter was achieved by placing the hot samples among massive aluminium blocks. In comparing these figures with those for the MIG process, we note that there is a more important difference for the heating rate than for the cooling rate. The effect of the heating rate on the microstructure was investigated by varying the size of the samples and the furnace temperature. In this way, heating rates in the range 0.5-10 C s"1 were attained; higher heating rates can only be achieved by using a different technique, such as one based on the Joule effect. Testing and microstructure examination were carried out either after a period of natural ageing longer than 30 days. Stress corrosion susceptibility was evaluated by means of SSRT and DBWL tests. The former was carried out on cylindrical samples with ST like axial direction (gauge length 12.5 mm and diameter 5.00 mm), in naturally aerated aqueous solutions of NaCI (3.5%) at a pH of 6.5 and a temperature around 23 C; the potential was varied in the range -700, -1700 mV (referred to the standard calomel electrode) and the crosshead speed between 1.7x10-g and 10-6 m.s -1. The results obtained in dry air (samples immersed in magnesium perchlorate), at a crosshead speed of 5x10-9 ms "1 were used as reference; thus the UTS results obtained in aggressive media are presented as a percentage of those obtained in dry air. DCB samples were machined in such a way that crack propagation occurred in the rolling direction, with the load applied in the ST direction, and had a half height of 12.5 mm, a total length of 135 mm, and a thickness of 12.5 mm in the WZL samples and 25 mm in the T651 samples. It should be noted thai DCB while zone-like samples are appreciably thinner by standard ASTM E-399 than required to achieve plane strain conditions; wider samples cannot be produced without an important loss in the heating rale, and therefore, significant changes in the microstructure of WSL samples. The samples were pop-in precracked and kept at constant slrain in aqueous solutions identical to those used in SSR testing. The microstructure of the samples was examined by means of optical Microscopy, whereas Scanning Electron Microscopy (SEM) was used to examine the fraclure surfaces; the latter were carried out in a Phillips microscope model 500. Results and Discussion

Microstructure of the White Zone The microstructure of the white zone in welded material has been investigated by many authors ]5, 6, 8-11]. Its salient features, as compared to the parent metal, have been outlined by Cordier et al [5] to be, i) Grains larger and more equiaxed, ii) a reduced amount of sub-micron intermetallic compounds within the grains, iii) large and finely dispersed intermetallic compounds and precipitate particles along the grain boundaries and iv) narrower precipitatefree zones. Some of the features described above are clearly visible in the micrographs of welded samples studied in this work (Fig. 1). First, it should be mentioned that the width of the white zone in our samples varied in the range 200500 pm, depending upon the welding conditions. The general aspect of the microslructure of the weldments around lhe white zone, is shown in Fig. 1. The enlarged micrograph (Fig. lb) illustrales the microslructural features of the white zone mentioned above (i to iii). The chain-like distribution of particles along the grain boundaries is specially remarkable. Etching reveals a very low amount of particles within the grains; Ihis could be mainly due to two causes, a) absence of intermetallic particles (such as ZrAI3), and b) enrichment of boundaries with main alloying elements (Zn and Mg). Both can be consequences of lhe..sweeping of particles and solute elements by migrating boundaries [6]. It is also worth remarking that the grains are highly irregular (both in size and shape) and that there are boundaries where partial melting is clearly visible. The microstructure of WZL samples is illustrated in Fig. lc. As compared to the actual white zone, these samples show the following differences: a) larger grains, b) higher amount of particles wilhin the grains. Bolh features should be a consequence of the lower heating rates achieved in the fabrication of WZL samples, as compared to those in the welding process. This was checked by investigating the evolution of the microstructure with heating rate in the range 0.5-10 C s 1 . The results do in fact point in this direction, namely, increasing the heating rate from 3 to 10 C s"1 decreases by a factor of two the grain size. The amount of particles within the grains do also appreciably decrease. Despite these differences we suspect that WZL samples do reproduce most of the salient microstructural features of the white zone in actual weldments. Slow Strain Rate Test Results The mechanical and electrochemical properties of T651 and WZL material, relevanl to the present discussion are reported in Table I.

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TABLE 1 Characteristics of Alloy AA 7017 in the T651 and WZL Tempers. TEMPER

~0.2 (MPa)

~UTS (MPa)

T651

397

461

WZL

265

329

Fract. Strain % (over 10 ram)

KIC (MPa m 1/2)

FCP (mVsc E)

8.5

27.6

-91 5

2.5

28.3

-9 26

W'Z LIKE

C

FIG. 1 Microstructure of: a) white zone in welded samples, general view, b) detail of a) and c) white zone-like samples. The KIC values were determined following the ASTM E-399 standard, in the short transverse (S-L) direction. On the other hand, the free corrosion potential (referred Io the standard calomel electrode) was measured in a naturally aerated aqueous solution of NaCI (3.5%) with a pH of 6.5. Fig. 2 shows SSR results for WZL samples at a strain rate of 5x10 9 m.s -1 and several potenlials in lhe range -800, -1500 mVSCE; results for samples in the T651 temper, to be considered as a reference, are also included. In the vertical axis appears the percentage of rupture load with respect to the corresponding rupture load in dry air. It is noted a very strong dependence of the ultimate tensile strength on the testing potential. A maximum in UTS, around -1200 mVsc E, is clearly noticed, indicating that there is a potential window within which the WZL material is safer, as also observed in actual weldments [4].

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AI-Zn-Mg WELDMENTS

100

/f%

121

z

'°

60

j

/

o o.i !

50

W

0

,,

,_

30

~ a l ( ) 9 mls

20

~

-1800

/0

A T

-1600

I

Crack Propagation Studies: DBWL Results

651

t

I II

I

- 800 -600 /| POTENTIAL ,(mVsCE) IL--T 6St / L---WZL .I FCP -1400

-1200

-lO00~

FIG. 2 Slow strain rate results for the AA 7017 alloy in the T651 temper (O), and white zone-like

samples WZL (A).

a

The most appealing features of the fractured surfaces of samples tested at -1200 mVsc E are illustrated in Fig. 3 f o r T651 material and WZL samples. Two zones are clearly differentiated in T651 material (Fig. 3a); a peripherical ring where fracture surface is flat, corresponding to stress corrosion, and a cenlral zone showing the dimples characteristic of the ductile fracture promoted by overstress; these features are among the typical of stress corrosion failures in most materials. On the other hand, WZL samples show a more uniform fracture surface (Fig. 3b); no region can be differentiated from the rest of the surface. Fraclure is intergranular and reveals the recrystallized coarse grained microstructure characteristic of WZL samples; it is interesting to note that fracture surfaces in actual weldments [4, 9] are very similar to those shown in Fig. 3c.

O WZL

10 0

/

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SSRT resulls for WZL samples indicate that they are much more susceptible to stress corrosion cracking than the parent material (in the T651 temper). These results are in line with those obtained by Holroyd et al [4] in actual weldments. Now, the question is which is the determining factor in this behaviour, namely, crack initiation or crack propagation, To elucidate this point we have measured the crack growth speed in our simulated samples (WZL) by means of double beam wedge-loaded specimens. The results are summarized in Fig. 4. Crack growth occurred at a very high pace during

T-651

C FIG. 3 Scanning electron micrographs of fracture surfaces of alloy AA 7017 after slow strain rate testing (-1200 mVscE, and strain rate 0.3 p_m min-1): a) General aspect of T651 material, b) same as a) for white zone-like samples, and c) enlarged micrograph of b).

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]2

A1-Zn-Mg W E L D M E N T S

2fl95

iGs the first 24 hours of testing. In order to ensure that overloading was not obscuring the results, some samples were kept in dry air for around a week, before they were introduced in the aggressive medium. No appreciable advance of the crack was observed (less than 3 mm) during that period. Instead, right after the introduclion in the corrosive medium, a very fasl growth of lhe crack occurred, leading to speed values identical (within our experimental accuracy) to those observed in samples directly introduced in the sodium chloride aqueous solution. This indicates that crack growth was in fact a consequence of stress corrosion cracking.

~ o°

167

£

o

5°

/; 0

'u~

In comparing these results with those for T65! material tested in the same conditions (see Fig. 4) we note that crack growth speeds for WZL samples are much higher. This result indicates that crack growlh might well be the factor determining the higher sensitivity of WZL samples to stress corrosion, and gives support to models of stress corrosion failure in actual weldmenls [4] which ascribed a major role to the while zone in the crack propagation process. The results presenled here are the first experimental evidence of anomalously high growth speeds in heat affected zones of welded aluminium alloys.

10

o WZL • T 651

10

20-K,(MPa

30

&0

Discussion

m 1/2)

FIG. 4 DBWL test results for the T651 temper and white zone-like samples of alloy AA 7017.

The present results for stress corrosion cracking in white zone-like samples can be qualitalively underslood in terms of a rationale recently proposed to explain the slress corrosion susceptibility of aluminium alloys [14], which stresses the relevance of the microstructure in hydrogen embrittlement. The argument, based on detailed microstructural studies, ascribes a major role to planar slip: extensive planar slip favours hydrogen transport, whereas non-planar slip difficults that process. Thus, a microstructure leading to planar slip would be much more sensitive to stress corrosion cracking than one producing non-planar slips. Then, nalurally aged samples having a microstructure formed by precipitates (G.P. Zones and 11') which can be cut by slip bands, and therefore build up a planar slip structure, would be much more sensitive to cracking than artificially aged samples, whose microstructure contains particles (q and 11') which force slip bands to bend; this gives a sound interpretation of experimental data for aluminium alloys of the 7000 series. In the present case an explanation along these lines is much simpler. White zonelike samples have a much lower density of particles within the grains than standard material; therefore, having no obstacles to bend the slip bands, it will be much more sensitive to stress corrosion cracking, no matter what the nature of the precipitate particles. Concluding Remarks The following conclusions can be drawn from the results presented in this work: i)

The main microstructural features which characterize the so called white zone in actual weldments, namely, recrystallized grains and a reduced amount of particles (both precipitate and intermetallic) within the grains, can be satisfactorily reproduced by the procedures utilized in this work.

it)

SSRT resulls indicate that white zone-like samples are much more susceptible to stress corrosion than T651 material, in agreement with results obtained by Schmiedel and Gruhl [8] from standard constant tensile load tests.

i i i)

Both the SSRT results and the features of the fracture surfaces, reported in this work for WZL samples, are remarkably similar to those observed in actual weldments by Holroyd et al [4]. In particular, a maximum was observed in the UTS around -1200 mVsc E.

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i v) DBWL tests indicate that crack speeds in WZL samples are between three and four orders of magnitude higher than those for T651 material. This result suggests that the relevant factor in determining the stress corrosion susceptibility of white zone-like samples, and very likely of actual weldments [4], is the crack growth speed in the heat affected zone. To further support this conclusion, c~ack growth speeds in the white zone of actual weldments should be measured.

Acknowledgements Partial support of ~his research by the Spanish Comisi6n Interminislerial de Ciencia y Tecnologia (CICYT) under Grant No. PC85-0096-C01-01, is gratefully acknowledged by the authors. Thanks are also due to the Industria Espafiola del Aluminio, S.A. for permission to publish this work.

References 1. 2. 3. 4.

S. Abis and E.D. Russo, !n Proc. Conf. L'Aluminio nei Transporti (1987). G.M. Scamans, J.A. FiuP.t~r ~,nd N.J.H. Holroyd, In Proc. 8th. International Light Metal Conference, I, 699 (1987). K.G. Kenl, Metals ~:~c; k4atefials, ;~.~atallurgical Reviews, Review 147, 135 (1970). N.J.H. Holroyd, W. Hepples &rid G.M. Scamans, In Proc. Conf. Fatigue Corrosion Cracking, Fracture Mechanics and Failure Analysis, ASM, 291 ('i'~85). 5. H. Cordier, M. Shippers and !.J. PoZmear, Z. Metallkunde, 68, 280 (1977). 6. I.J. Polmear, In Recrystallization ~nd Grain Growth of Multi-Phase and Particle containing Materials, RISO, 177 (1980). 7. E.D. Russo and S. Abis, In Advanced Materials Research and Developments for Transport, Light Metals, 227 (1985). 8. H. Schmiedel and W. Gruhl, Melall., 38, 32 (1984). 9. B. Grzemba, H. Codier and W. Gruhl, Aluminium, 63, 496 (1987). 10. M.S. Rahman and I.J. Polmear, Z. Melatlkde, 74, 733 (1983). 11. M.C. Reboul, B. Dubosl and M. Lashemeres, Corrosion Science, 25, 999 (1986). 12. R.N. Parkins, ASTM-STP 665, 5 (1979). 13. N.J.H. Holroyd and G.M. Scamans, ASTM-STP 821,202 (1984). 14. D. Nguyen, A.W. Thompson and I.M. Bernstein, Acla Metall., 35,2417 (1987).