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Fracture of bead-on-plate CO2 laser welds in the AlLi alloy 8090
Pergamon
Scripta Metallurgicae t Materialia. Vol. 31, No. 12, pp. 1717-1722. 1994 Copyright © 1994 ElsevierScienceLid Printed in the USA. All rights reserved 0956-716X/94 $6.00 + 00
FRACTURE OF BEAD-ON-PLATE COz LASER WELDS IN THE A1-Li ALLOY 8090 I.R. Whitaker ~ and D.G. McCartney z tDepartment of Engineering Materials, University of Sheffield, Sheffield, S 1 4DU, UK ZDepartment of Materials Engineering and Materials Design, University of Nottingham, Nottingham, NG7 2RD, UK
(Received April 21, 1994) (Revised July 25, 1994) Introduction The development of aluminium alloys containing Li has progressed during the twentieth century to the present day (1) with the discovery of improved specific properties creating potential applications for these types of alloys in the aerospace industry. The density of aluminium is reduced by about 3% for each weight percent addition of Li while the Young's modulus is increased by about 6% (2). However, binary alloys have very poor fracture toughness properties and improvements have been made by alloying element additions, together with thermomechanical treatments. Furthermore, additions of Cu, Mg and Zr can enhance the strength, modulus of elasticity, fatigue, corrosion, and stress corrosion cracking resistance. Thus, AI-Li alloy development has been focused towards replacing conventional aircraft alloys such as the 2XXX and 7XXX series A1 alloys. These types of alloys, when used in airframe structures, are mechanically fastened which requires several operations with a consequently low production rate and an upper limit to the thickness of the sheet being fastened. Hence there is a strong desire to employ welding as part of the manufacturing route when introducing 8090 as a replacement alloy. Electron beam (EB) welds were made by Edwards et al.(3) on 8090. In the as-welded condition joints failed in a brittle inter-granular manner but a higher fracture strength, together with enhanced grain boundary ductility was observed after a T6 heat treatment (solution treatment plus artificial aging). In this latter heat treatment condition, the weld strength was -73% of the base metal strength. A study into the mechanical properties of EB welds in 8090 by Le Poac et al.(4) revealed that post-weld aging alone increased weld metal hardness but reduced tensile strength. CO2 laser welding has been attempted by Biermann et al.(5) and Gnanamuthu et al.(6). After a T6 heat treatment, the welds achieved 80% (5) and _>90% (6) of the parent metal strength in each case. Previous work by the present authors (7)(8)(9) has described the microstructure of bead-on-plate CO 2 laser welds produced in the AI-Li based alloy 8090 in both the as-welded and post-weld heat treated condition. Identical procedures and material were used as in the present case. In the solidified weld, columnar grains formed near the top surface, whereas equiaxed grains were present throughout the rest of the weld. Following heat treatment, the fusion zone grain size was found to be dependent on the heating rate to the solution treatment temperature (803K). After heating at 100 K/min, massive grains up to 200~ma across were observed, with an average grain size of -221am. However, after heating at rates of 10 and 1 K/min the average grain size was reduced to -11 and -101arn, respectively, and no massive grains were present in the latter case. These differences were shown (8) to be due to the formation of different dispersions of metastable, cubic 13' phase (AIjZr). At 1 K/min heating rate a sufficient volume fraction of 13' formed to pin grain boundaries whereas at 100 K/min heating rate there was insufficient time available for appreciable amounts of [3' to nucleate and grow. The main aim of the present study was to assess the effect of fusion zone microstructure on the fracture strength and fracture behaviour of the weldment.
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Experimental The nominal composition of the alloy 8090 used in this study is given in Table 1. Autogenous bead-on-plate CO2 laser welds were made using a Ferranti 5kW laser which had a continuous wave output at 10.61am. The laser beam was focused onto the sheet surface and welding was performed on a 4mm thick sheet with a beam power of -2kW and a welding speed of -70mrrds. These conditions caused full penetration of the sheet by the bead-on-plate weld. A coaxial helium gas flow was employed to provide a protective atmosphere. Post-weld heat treatments of tensile samples were carried out at a solution treatment temperature of 803 K for -40 rains and heating rates to 803K of 1 and 100 K/min were employed. Samples were air cooled after solution treatment. Aging was performed at 463K for 16 hours. Air cooling was utilised rather than water quenching in order to simulate as closely as possible a proposed manufacturing route. TABLE 1 Nominal Composition of the Alloy 8090 (wt%).
Tensile tests were performed on five notched samples in each heat treatment condition (9). The notch geometry and orientation in the weld is illustrated in Fig. 1. The notch was added in the first instance to ensure that the test piece failed in the fusion zone. A MAYES DM4 tensile testing machine was used operating at a cross-head speed of 1.5mm/min and the load vs. cross-head displacement was continuously recorded on an X-Y plotter. Fractography was carried out on a JEOL 6400 SEM with WINSEM computer software and an accelerating voltage of 15kV.
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FIG. 1. Illustration of a notch tensile test specimen cut from 4mm thick sheet, dimensions in mm. Notch tensile tests were performed on base and bead-on-plate CO2 weld metal in the following heat treatment conditions where 'S.T.' defines solution treatment at 803K for -40 minutes, with a heating rate control of 1 or 100 K/min for the weld samples followed by air cooling and 'A' defines aging for 16 hours at 463K: Heat Treatment #1 Heat Treatment #2 Heat Treatment #3 Heat Treatment #4 Heat Treatment #5
base metal, S.T. + A, weld metal, as welded, weld metal, heated at 1 K/min to the S.T. temperature, weld metal, heated at 100 K/rain to the S.T. temperature, weld metal, heated at 1 K/min to the S.T. temperature + A,
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Results Table 2 lists the fracture strengths of samples heat treated as described above. Considering the weld metal, the results show that the notched weld metal strength is lowest in the as-welded condition, with a value of 235_+18 MPa. However, solution treatment and air cooling can increase this to 312_+38 MPa and 323-+28 MPa for heating rates of 1 and 100 K/min to the solution treatment temperature, respectively. Maximum strength was achieved after solution treating and aging which resulted in a tensile strength of 415-+19 MPa. This value is comparable to the notched base metal strength in the same heat treatment condition, 426-+5 MPa. An SEM micrograph from the fracture surface of a notched base metal tensile specimen which was tested after solution treating and aging is shown in Fig. 2. The crack has propagated intergranularly along high angle grain boundaries to produce relatively large grain faces, e.g. marked T , due to the pancake structure of the base metal grains in the rolling direction of the sheet. Smaller grains can be seen where the crack has followed lower angle 2grain boundaries. Furthermore, shallow dimples are also visible on grain faces. These were probably caused by void formation around particles such as S' (A12CuMg) or T t (AI2CuLi) which can form at the grain boundaries during aging. TABLE 2 8090 Base and CO2 Weld Metal Tensile Strengths From Notched Samples Heat Treatment Fracture Strength (MPa)
#1
#2
#3
#4
#5
426+_5 235_+18 312_+38 323_+28 415_+19
The error indicated is two standard deviations about the mean. The failure mode of the notched as-welded sample is shown in Fig. 3. At a low magnifcation in Fig. 3(a) two very different regions can be seen on the fracture surface. On the left, large rupture steps are present with cracks which run the entire length of the step, e.g. marked 'S'. This region represents the fracture of the coarse columnar grain structure in the region of the weld near the top surface of the sheet and failure has occurred along solidification cell boundaries.
FIG. 2. Fracture surface of the 8090 base metal. after heat treatment #1.
FIG. 3(a). Fracture surface of the 8090 weld in the as-welded condition (heat treatment #2).
Fig. 3(b) reveals that the steps were caused by a mixed mode of failure. The crack has followed the cell boundaries but some transgranular failure has occurred as shown by the large ductile dimples separating the rupture steps, marked 'DD'. What appear to be slip bands, marked 'SB', can also be seen on the surfaces of the elongated grains. The fracture surface of the lower region of the weld near the bottom surface of the sheet is shown in Fig. 3(c) and fine equiaxed grains are clearly visible. Individual grains e.g. marked 'C' can be seen
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which suggests that failure occurred intergranularly and the dendritic morphology is visible on the grain surfaces, marked T. However, some regions of ductility can be seen. At positions marked 'T' there are obvious ductile dimples due to transgranular failure. Again, secondary cracks e.g. marked 'S' in Fig. 3(c) are present.
FIG. 3(b). Higher magnifieation of the left of Fig. 3(a).
FIG. 3(e). Higher magnification of the right of Fig. 3(a).
The fracture surfaces of the welds which were solution treated with heating rates of 1 and 100 K/min are very similar in appearance. Figs. 4(a) and (b) show regions of equiaxed grains that are comparable in that failure has occurred predominantly along grain boundaries and individual grains are visible in both micrographs. However, Fig. 4(b) shows generally larger grains and the region ringed 'A' appears to consist of a larger than average, single grain with an internal dendritic sub-structure. It is likely that the crack has propagated along the spatial distribution of undissolved intermetallic particles within this grain.
FIG. 4(a). Weld fracture surface after heat treatment #3.
HG. 4(b). Weld fracture surface after heat treatment #4.
After solution treating and aging, failure is intergranular predominandy at high angle grain boundaries, though
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transgranular regions are visible marked 'T' in Fig. 5 which contain large ductile dimples. On the intergranular region marked 'I' in Fig. 5 the dendritic structure is present together with shallow dimples. Discussion The fracture behaviour of A1-Li based alloys has been addressed generally by Quist et al.(1). More specifically, several workers have studied the fracture of 8090 type alloys (10)(11)(12). Indeed, Weft et al.(12) have shown that in an A1-Li based alloy secondary cracks formed in the short-transverse direction of the rolled sheet. However, as well as the intergranular fracture, final failure occurred transgranularly across the delaminated grains. After aging, parallel ridges were observed on the transgranular fracture surface which were possibly caused by slip bands. In an ingot-cast A1-Li based alloy Butler at al.(13) determined that fracture occurred by a mixture of intergranular and transgranular modes. However, in the solution treated condition the intergranular surface was faceted in appearance while in the over-aged condition ductile microvoiding had occurred. Thus void nucleation was made easier in aged material and Fe-Cu rich particles were associated with the microvoids. In the present study the failure of solution treated and aged base metal sheet is entirely consistent with the behaviour discussed above and limited dimpling was indeed observed, but no attempt was made at determining the mechanism of its formation. There has been much more limited work on the fractography of A1-Li alloy welds. However, Edwards et al.(3) showed that tungsten-inert-gas and electron beam welds in 8090 failed intergranularly but with enhanced grain boundary ductility after solution treatment and aging. Microvoids formed at grain boundary precipitates and the low ductility observed was associated with the grain boundary structure rather than the grain size. These findings are very similar to those in the present work on bead-on-plate laser welds in 8090. As indicated in Fig. 6, it seems likely that it is the distribution of intermetallic particles in the solution treated bead-on-plate welds, which exist at prior solidification cell boundaries, that is responsible for controlling the fracture behaviour. Such a fracture path is indicated as A-A', distinct from the high angle grain boundaries indicated as B-B'. This satisfactorily explains why the fracture strength did not vary significantly with heating rate to the solution treatment temperature (#3 and #4) whereas markedly different grain sizes were formed (8)(9). Though no attempt was made to identify the particles within the shallow microvoids, the intermetallic particles which exist at prior solidification cell boundaries were shown by transmission electron microscopy to be enriched in Fe and Cu as reported previously (9). Therefore, following the model suggested by Butler et al.(13), the sequence of events leading to fracture could be as follows. Planar slip in the grains containing 5' precipitates causes dislocation pile-ups at these intermetallic particles be they inter or intragranular. As a result of the local stress concentration, particle-matrix decohesion could occur and where several particles are grouped in stringers the voids can coalesce during decohesion resulting in a dimpled, intergranular fracture surface. That microvoid formation is enhanced after solution treating and aging is probably due to the ease of coplanar slip across larger 5' precipitates. The fracture strengths clearly show that similar properties can be obtained between base and weld metal after an appropriate post-weld heat treatment, despite the large difference in grain size which can occur between the heat treated weld fusion zone and the base metal. Moreover the fracture strength of the solution treated and aged weld is approximately equal to that of the base metal, thus being higher than that achieved in several other welding studies. Conclusions A post-weld solution heat treatment followed by aging was able to produce a good match in fracture strengths between 8090 base metal and CO 2 bead-on-plate weld metal with the weld strength being -97% of the base metal strength. In this respect the alloy can be classed as weldable. The 8090 CO2 weld failed in a brittle, predominantly intergranular manner and the fracture strength did not depend on the weld metal grain size. This was due to the fracture path being determined by the residual solidification structure i.e. the secondary dendrite arm spacing resulting from weld solidification and not the grain size resulting from post weld heat treatment.
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;,-
FIG. 5. Weld fracture surface after heat treatment #5.
" :-\ ___----- B'
FIG. 6. Schematic diagram of the fracture path in the weld metal.
A..cknowledgments The authors would like to thank N. Calder and Professor W.M. Steen for their part in providing the welded sheets and also Dr. B. Noble for elucidating aspects of the fracture surfaces. This work was carried out at Nottingham University and was supported by the Science and Engineering Research Council and British Aerospace through a Co-operative Research Grant and CASE award research studentship. References
1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13.
W.E. Quist and G.H. Narayanan, Treatise on Materials Science and Technology, vol. 31, pp. 219-254, Academic Press Inc. (1989). ISBN 0-12-341831-3. K.K. Sankaran and N.J. Grant, Mat. Sci. and Eng., 44, 213 (1980). M.R. Edwards and V.E. Stoneham, Journal de Physique, Colloque C3, Supplement to no. 9, 48 C3-293 (1987). P. LePoac, A.M. Nomine and D. Miannay, ibid, C3-301. B. Biermann, R. Dierken, R. Kupfer, A. Lang and H.W. Bergmann, Proc. of the Sixth Int. Al-Li Alloys Conf., p.1159, Garmisch-Partenkirchen, Germany, (1991). D.S. Gnanamuthu and R.J. Moore,s in Power Beam Processing, Proc. of Int. Conf. on Power Beam Processing, Ed. E.A. Metzbower and D. Hauser, ASM InL Ohio (1989) page 181. I.R. Whitaker, N. Calder, D.G. McCartney and W.M. Steen, J. Mat. Sci., 28, 5469 (1993). I.R. Whitaker, N. Calder, D.G. McCartney and W.M. Steen, Journal de Physique IV, Colloque C7, Supplement to Journal de Physique lII, Volume 3, p.1053 (1993). I.R. Whitaker, Ph.D. Thesis, University of Nottingham, UK (1994). T.S. Srivatsan and T. Alan Place, J. Mat. Sci., 24, 1543 (1989). S.P. Lynch, Mat. Sci and Eng., A136, 25 (1991). J.A. Wert and J.B. Lumsden, Scripta Met., 19, 205 (1985). E.P. Butler, N.J. Owen and D.J. Field, Mat. Sci and Tech., 1,531 (1985).
Scripta Metallurgicae t Materialia. Vol. 31, No. 12, pp. 1717-1722. 1994 Copyright © 1994 ElsevierScienceLid Printed in the USA. All rights reserved 0956-716X/94 $6.00 + 00
FRACTURE OF BEAD-ON-PLATE COz LASER WELDS IN THE A1-Li ALLOY 8090 I.R. Whitaker ~ and D.G. McCartney z tDepartment of Engineering Materials, University of Sheffield, Sheffield, S 1 4DU, UK ZDepartment of Materials Engineering and Materials Design, University of Nottingham, Nottingham, NG7 2RD, UK
(Received April 21, 1994) (Revised July 25, 1994) Introduction The development of aluminium alloys containing Li has progressed during the twentieth century to the present day (1) with the discovery of improved specific properties creating potential applications for these types of alloys in the aerospace industry. The density of aluminium is reduced by about 3% for each weight percent addition of Li while the Young's modulus is increased by about 6% (2). However, binary alloys have very poor fracture toughness properties and improvements have been made by alloying element additions, together with thermomechanical treatments. Furthermore, additions of Cu, Mg and Zr can enhance the strength, modulus of elasticity, fatigue, corrosion, and stress corrosion cracking resistance. Thus, AI-Li alloy development has been focused towards replacing conventional aircraft alloys such as the 2XXX and 7XXX series A1 alloys. These types of alloys, when used in airframe structures, are mechanically fastened which requires several operations with a consequently low production rate and an upper limit to the thickness of the sheet being fastened. Hence there is a strong desire to employ welding as part of the manufacturing route when introducing 8090 as a replacement alloy. Electron beam (EB) welds were made by Edwards et al.(3) on 8090. In the as-welded condition joints failed in a brittle inter-granular manner but a higher fracture strength, together with enhanced grain boundary ductility was observed after a T6 heat treatment (solution treatment plus artificial aging). In this latter heat treatment condition, the weld strength was -73% of the base metal strength. A study into the mechanical properties of EB welds in 8090 by Le Poac et al.(4) revealed that post-weld aging alone increased weld metal hardness but reduced tensile strength. CO2 laser welding has been attempted by Biermann et al.(5) and Gnanamuthu et al.(6). After a T6 heat treatment, the welds achieved 80% (5) and _>90% (6) of the parent metal strength in each case. Previous work by the present authors (7)(8)(9) has described the microstructure of bead-on-plate CO 2 laser welds produced in the AI-Li based alloy 8090 in both the as-welded and post-weld heat treated condition. Identical procedures and material were used as in the present case. In the solidified weld, columnar grains formed near the top surface, whereas equiaxed grains were present throughout the rest of the weld. Following heat treatment, the fusion zone grain size was found to be dependent on the heating rate to the solution treatment temperature (803K). After heating at 100 K/min, massive grains up to 200~ma across were observed, with an average grain size of -221am. However, after heating at rates of 10 and 1 K/min the average grain size was reduced to -11 and -101arn, respectively, and no massive grains were present in the latter case. These differences were shown (8) to be due to the formation of different dispersions of metastable, cubic 13' phase (AIjZr). At 1 K/min heating rate a sufficient volume fraction of 13' formed to pin grain boundaries whereas at 100 K/min heating rate there was insufficient time available for appreciable amounts of [3' to nucleate and grow. The main aim of the present study was to assess the effect of fusion zone microstructure on the fracture strength and fracture behaviour of the weldment.
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Experimental The nominal composition of the alloy 8090 used in this study is given in Table 1. Autogenous bead-on-plate CO2 laser welds were made using a Ferranti 5kW laser which had a continuous wave output at 10.61am. The laser beam was focused onto the sheet surface and welding was performed on a 4mm thick sheet with a beam power of -2kW and a welding speed of -70mrrds. These conditions caused full penetration of the sheet by the bead-on-plate weld. A coaxial helium gas flow was employed to provide a protective atmosphere. Post-weld heat treatments of tensile samples were carried out at a solution treatment temperature of 803 K for -40 rains and heating rates to 803K of 1 and 100 K/min were employed. Samples were air cooled after solution treatment. Aging was performed at 463K for 16 hours. Air cooling was utilised rather than water quenching in order to simulate as closely as possible a proposed manufacturing route. TABLE 1 Nominal Composition of the Alloy 8090 (wt%).
Tensile tests were performed on five notched samples in each heat treatment condition (9). The notch geometry and orientation in the weld is illustrated in Fig. 1. The notch was added in the first instance to ensure that the test piece failed in the fusion zone. A MAYES DM4 tensile testing machine was used operating at a cross-head speed of 1.5mm/min and the load vs. cross-head displacement was continuously recorded on an X-Y plotter. Fractography was carried out on a JEOL 6400 SEM with WINSEM computer software and an accelerating voltage of 15kV.
6
191 _
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1 2 . 5
A
.
"~
_-
.
-
-
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.
' o.~""
IB i~'
47
l
. A
r=o 1
/ 6 5 -
.
.
.
.
-
•
•
FIG. 1. Illustration of a notch tensile test specimen cut from 4mm thick sheet, dimensions in mm. Notch tensile tests were performed on base and bead-on-plate CO2 weld metal in the following heat treatment conditions where 'S.T.' defines solution treatment at 803K for -40 minutes, with a heating rate control of 1 or 100 K/min for the weld samples followed by air cooling and 'A' defines aging for 16 hours at 463K: Heat Treatment #1 Heat Treatment #2 Heat Treatment #3 Heat Treatment #4 Heat Treatment #5
base metal, S.T. + A, weld metal, as welded, weld metal, heated at 1 K/min to the S.T. temperature, weld metal, heated at 100 K/rain to the S.T. temperature, weld metal, heated at 1 K/min to the S.T. temperature + A,
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Results Table 2 lists the fracture strengths of samples heat treated as described above. Considering the weld metal, the results show that the notched weld metal strength is lowest in the as-welded condition, with a value of 235_+18 MPa. However, solution treatment and air cooling can increase this to 312_+38 MPa and 323-+28 MPa for heating rates of 1 and 100 K/min to the solution treatment temperature, respectively. Maximum strength was achieved after solution treating and aging which resulted in a tensile strength of 415-+19 MPa. This value is comparable to the notched base metal strength in the same heat treatment condition, 426-+5 MPa. An SEM micrograph from the fracture surface of a notched base metal tensile specimen which was tested after solution treating and aging is shown in Fig. 2. The crack has propagated intergranularly along high angle grain boundaries to produce relatively large grain faces, e.g. marked T , due to the pancake structure of the base metal grains in the rolling direction of the sheet. Smaller grains can be seen where the crack has followed lower angle 2grain boundaries. Furthermore, shallow dimples are also visible on grain faces. These were probably caused by void formation around particles such as S' (A12CuMg) or T t (AI2CuLi) which can form at the grain boundaries during aging. TABLE 2 8090 Base and CO2 Weld Metal Tensile Strengths From Notched Samples Heat Treatment Fracture Strength (MPa)
#1
#2
#3
#4
#5
426+_5 235_+18 312_+38 323_+28 415_+19
The error indicated is two standard deviations about the mean. The failure mode of the notched as-welded sample is shown in Fig. 3. At a low magnifcation in Fig. 3(a) two very different regions can be seen on the fracture surface. On the left, large rupture steps are present with cracks which run the entire length of the step, e.g. marked 'S'. This region represents the fracture of the coarse columnar grain structure in the region of the weld near the top surface of the sheet and failure has occurred along solidification cell boundaries.
FIG. 2. Fracture surface of the 8090 base metal. after heat treatment #1.
FIG. 3(a). Fracture surface of the 8090 weld in the as-welded condition (heat treatment #2).
Fig. 3(b) reveals that the steps were caused by a mixed mode of failure. The crack has followed the cell boundaries but some transgranular failure has occurred as shown by the large ductile dimples separating the rupture steps, marked 'DD'. What appear to be slip bands, marked 'SB', can also be seen on the surfaces of the elongated grains. The fracture surface of the lower region of the weld near the bottom surface of the sheet is shown in Fig. 3(c) and fine equiaxed grains are clearly visible. Individual grains e.g. marked 'C' can be seen
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which suggests that failure occurred intergranularly and the dendritic morphology is visible on the grain surfaces, marked T. However, some regions of ductility can be seen. At positions marked 'T' there are obvious ductile dimples due to transgranular failure. Again, secondary cracks e.g. marked 'S' in Fig. 3(c) are present.
FIG. 3(b). Higher magnifieation of the left of Fig. 3(a).
FIG. 3(e). Higher magnification of the right of Fig. 3(a).
The fracture surfaces of the welds which were solution treated with heating rates of 1 and 100 K/min are very similar in appearance. Figs. 4(a) and (b) show regions of equiaxed grains that are comparable in that failure has occurred predominantly along grain boundaries and individual grains are visible in both micrographs. However, Fig. 4(b) shows generally larger grains and the region ringed 'A' appears to consist of a larger than average, single grain with an internal dendritic sub-structure. It is likely that the crack has propagated along the spatial distribution of undissolved intermetallic particles within this grain.
FIG. 4(a). Weld fracture surface after heat treatment #3.
HG. 4(b). Weld fracture surface after heat treatment #4.
After solution treating and aging, failure is intergranular predominandy at high angle grain boundaries, though
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transgranular regions are visible marked 'T' in Fig. 5 which contain large ductile dimples. On the intergranular region marked 'I' in Fig. 5 the dendritic structure is present together with shallow dimples. Discussion The fracture behaviour of A1-Li based alloys has been addressed generally by Quist et al.(1). More specifically, several workers have studied the fracture of 8090 type alloys (10)(11)(12). Indeed, Weft et al.(12) have shown that in an A1-Li based alloy secondary cracks formed in the short-transverse direction of the rolled sheet. However, as well as the intergranular fracture, final failure occurred transgranularly across the delaminated grains. After aging, parallel ridges were observed on the transgranular fracture surface which were possibly caused by slip bands. In an ingot-cast A1-Li based alloy Butler at al.(13) determined that fracture occurred by a mixture of intergranular and transgranular modes. However, in the solution treated condition the intergranular surface was faceted in appearance while in the over-aged condition ductile microvoiding had occurred. Thus void nucleation was made easier in aged material and Fe-Cu rich particles were associated with the microvoids. In the present study the failure of solution treated and aged base metal sheet is entirely consistent with the behaviour discussed above and limited dimpling was indeed observed, but no attempt was made at determining the mechanism of its formation. There has been much more limited work on the fractography of A1-Li alloy welds. However, Edwards et al.(3) showed that tungsten-inert-gas and electron beam welds in 8090 failed intergranularly but with enhanced grain boundary ductility after solution treatment and aging. Microvoids formed at grain boundary precipitates and the low ductility observed was associated with the grain boundary structure rather than the grain size. These findings are very similar to those in the present work on bead-on-plate laser welds in 8090. As indicated in Fig. 6, it seems likely that it is the distribution of intermetallic particles in the solution treated bead-on-plate welds, which exist at prior solidification cell boundaries, that is responsible for controlling the fracture behaviour. Such a fracture path is indicated as A-A', distinct from the high angle grain boundaries indicated as B-B'. This satisfactorily explains why the fracture strength did not vary significantly with heating rate to the solution treatment temperature (#3 and #4) whereas markedly different grain sizes were formed (8)(9). Though no attempt was made to identify the particles within the shallow microvoids, the intermetallic particles which exist at prior solidification cell boundaries were shown by transmission electron microscopy to be enriched in Fe and Cu as reported previously (9). Therefore, following the model suggested by Butler et al.(13), the sequence of events leading to fracture could be as follows. Planar slip in the grains containing 5' precipitates causes dislocation pile-ups at these intermetallic particles be they inter or intragranular. As a result of the local stress concentration, particle-matrix decohesion could occur and where several particles are grouped in stringers the voids can coalesce during decohesion resulting in a dimpled, intergranular fracture surface. That microvoid formation is enhanced after solution treating and aging is probably due to the ease of coplanar slip across larger 5' precipitates. The fracture strengths clearly show that similar properties can be obtained between base and weld metal after an appropriate post-weld heat treatment, despite the large difference in grain size which can occur between the heat treated weld fusion zone and the base metal. Moreover the fracture strength of the solution treated and aged weld is approximately equal to that of the base metal, thus being higher than that achieved in several other welding studies. Conclusions A post-weld solution heat treatment followed by aging was able to produce a good match in fracture strengths between 8090 base metal and CO 2 bead-on-plate weld metal with the weld strength being -97% of the base metal strength. In this respect the alloy can be classed as weldable. The 8090 CO2 weld failed in a brittle, predominantly intergranular manner and the fracture strength did not depend on the weld metal grain size. This was due to the fracture path being determined by the residual solidification structure i.e. the secondary dendrite arm spacing resulting from weld solidification and not the grain size resulting from post weld heat treatment.
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;,-
FIG. 5. Weld fracture surface after heat treatment #5.
" :-\ ___----- B'
FIG. 6. Schematic diagram of the fracture path in the weld metal.
A..cknowledgments The authors would like to thank N. Calder and Professor W.M. Steen for their part in providing the welded sheets and also Dr. B. Noble for elucidating aspects of the fracture surfaces. This work was carried out at Nottingham University and was supported by the Science and Engineering Research Council and British Aerospace through a Co-operative Research Grant and CASE award research studentship. References
1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13.
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