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Effect of grain boundary character distribution (GBCD) on the cavitation behaviour during superplastic deformation of Al 7475
Materials Science and Engineering A338 (2002) 243 /252 www.elsevier.com/locate/msea
Effect of grain boundary character distribution (GBCD) on the cavitation behaviour during superplastic deformation of Al 7475 C.L. Chen, M.J. Tan * School of Mechanical and Production Engineering, Nanyang Technological University, Singapore 639798, Singapore Received 28 November 2001; received in revised form 25 January 2002
Abstract The formation of cavity stringers in superplastically deformed Al 7475 alloy was investigated in this work. After superplastic deformation, cavity stringers have a tendency to form intergranularly along grain boundaries that are parallel to the rolling direction of the material. The microstructure and microtexture showed that the material has anisotropic grain boundary character distribution (GBCD). In the as-received state, small precipitates were found to be distributed preferentially along the grain boundaries which are parallel to the rolling direction. After superplastic deformation, a low melting point eutectic phase was found along the grain boundaries, which are also mainly parallel to the rolling direction. At the test temperature of 516 8C, the eutectic phase is expected to be in the liquid state. The presence of the liquid phase reduces the grain boundary bonding strength. It is proposed that the formation of cavity stringers is closely related to the GBCD. Based on the anisotropic distribution of GBCD and extensive grain boundary sliding during superplastic deformation, a model is proposed to explain the formation of cavity stringers. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Al 7475; Superplasticity; Cavity stringers; Grain boundary character distribution (GBCD)
1. Introduction Intergranular cavitation has been frequently reported in a wide range of superplastic materials [1 /4], while some other alloy systems, such as titanium based alloys, suffer relatively less from this deleterious phenomenon. Cavities generally nucleate at triple junctions of grain boundaries, in the vicinity of particles and even on ledges of grain boundaries, and/or directly from preexisting defects. Cavity nuclei can then grow, coalesce and interlink as deformation continues, leading to the degradation of the postdeformation properties of the materials and may even result in premature fracture during processing. The effects of strain rate, strain, temperature, back pressure, etc. on cavitation behaviour have been extensively studied [1,5 /9], and these reported general common features for many alloy systems. The effects of the microstructure of the materials on cavitation
* Corresponding author. Tel.: /65-790-5582; fax: /65-791-1859. E-mail address: [email protected] (M.J. Tan).
behaviour, on the other hand, seems more complicated. Small chemical compositional differences or subtle differences of solidification and/or thermal processing conditions between batches of the same materials can result in a large difference in cavitational behaviour [1]. One of the interesting phenomena during superplastic deformation is the formation of cavity stringers where a group of cavities are aligned in a specific direction, and this has been reported both in microduplex (MD) and quasi-single phase (QSP) superplastic materials [4,10,11]. Yousefini et al. [11] reviewed this phenomenon recently and found that in most cases, the cavity stringers align themselves parallel to the tensile axis; and in some cases little or no arrangements can be discerned, and on a few occasions the cavities show an alignment perpendicular to the tensile axis. The morphology of cavity stringers varies from one material to another and is affected by the superplastic forming conditions such as strain, strain rate and temperature. Zelin et al. [12,13] proposed a model to depict the possibility of cavity stringer formation based on the concept of cooperative grain boundary sliding (CGBS). They suggested that during sheet forming under a
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biaxial state of stress, cavity nucleation occurs in the intersecting sites of cooperative sliding surfaces, which operates in a direction close to that of the principal shear stress and this thereby aligns the stringers along the same orientation. This was consistent with their observation on both biaxial and uniaxial tensile specimens. Chakraborty and Earthman [14] carried out the numerical analysis on the interaction between boundary sliding and cavitation. In the absence of accommodation by diffusion, deformation, and/or boundary migration, the local stress concentration is relieved by the opening of a cavity. According to their calculations, the opening of such a cavity results in the stressing of the boundaries inclined to the stress axis near this cavity more than the transverse boundaries, leading to the further nucleation of cavities at these boundaries. Short rows of cavities inclined to the tensile axis can then result. Yousefini et al. [11] furthered Chakraborty and Earthman’s work [14] and proposed, based on the results of Zn /22wt.%Al materials, that the formation of cavity stringers is related to: (1) the favourable nucleation sites on former a boundaries which are related to the impurity segregation and thermal history of the material; and (2) their change in orientation on approaching the tensile axis as deformation proceeds. This latter factor, as a result of superplastic flow, plays an important role in the development of cavity stringers, since the evolution of cavity stringers is controlled by the factors related to the nature and characteristics of superplasticity rather than the cavity nucleation process, once the cavities are nucleated. It is important to mention that, for the MD materials, the cavity stringers are usually aligned parallel to the tensile axis regardless of the rolling direction. For some MD materials and for most QSP materials, the formation of cavity stringers along the rolling direction has been reported. This phenomenon is consistent with the fact that, during prior thermomechanical treatment, most superplastic materials are rolled either to produce the fine grain size necessary for superplastic deformation or to be a part of the processing procedure. Large particles, if present, can therefore break down and align as small particles along the rolling direction. Since superplastic materials are commonly tested with tensile axis parallel to the rolling direction, it was suggested that cavity nucleation at these particles leads to the formation of cavity stringers [4,10]. This suggestion was verified by the observation that changing the orientation of the test samples led to a corresponding change in the arrangement of cavity stringers [4,10]. Jiang et al. [10] and Shin and Park [4] have reported the directionally aligned cavity stringer in the finegrained Al 7075 alloy. In addition to the cavity nucleation effects, Jiang et al. [10] also concluded that the tensile superplastic deformation contributes to the
formation of the cavity stringers aligned along tensile axis. Their explanation is similar to that of Yousefini et al. [11]. Shin and Park [4] studied both fine and coarse grain Al 7075 alloys and found tensile direction-aligned stringers in fine grain materials and rolling directionaligned stringers in large grain materials. They found that the cavity stringer formation parallel to the tensile axis could be associated with the existence of dispersoid free zone (DFZ) which was formed in the vicinity of grain boundaries perpendicular to the tensile axis during superplastic deformation. The formation of DFZs is frequently observed and it may lead to the decrease of new cavity nucleation in their vicinity. These works are helpful in the understanding of the stringer formation mechanism. However, there are still some aspects that have yet not been well clarified: (1) When directional precipitate stringers appear, what mechanisms cause these cavities to interlink with each other along precipitate stringers, regardless of the stress direction? (2) When cavities were uniformly nucleated, such as the cases of QSP materials with no precipitate stringer, and MD materials, what mechanisms lead cavities to develop along stress direction, instead of along the transverse or 458 direction which usually takes place in the room temperature tensile condition? (3) What factors dominate the cavity stringer direction? These ambiguities in the stringer formation are related to the cavity growth, coalescence and interlinkage mechanisms. Studies on the cavity stringer behaviour are important so as to achieve the optimum microstructure for superplasticity. It is widely recognised that the grain boundary character has great impact on the superplasticity of the material, as a number of grain boundary phenomena are involved during superplastic deformation. In a previous work [15] on a typical QSP material, Al 7475, preferential cavity interlinkage of the material during superplastic deformation was reported. The present work focuses on the effect of GBCD on the preferential cavity interlinkage.
2. Experimental procedures Al 7475 sheet with chemical composition listed in Table 1 was used for superplastic deformation testing. The as-received sheet was in T4 condition (solution heat treated, quenched, and naturally aged). The tensile specimens were of 15 mm length, 4 mm width and 2 mm thickness in gauge parts, with the specimen cutting directions either parallel (L-sample) or perpendicular (T-sample) to the rolling direction of the sheet. The surfaces of the tensile sample gauge part were polished down to 1 mm diamond paste. Superplastic uniaxial tensile test and cavitation ratio determination have been detailed elsewhere [15]. In the present work, tensile tests were performed at 516 8C
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Table 1 Chemical composition of the alloy Element
Zn
Mg
Cu
Cr
Fe
Si
Ti
Mn
Al
Wt.%
5.64
2.34
1.58
0.19
0.08
0.05
0.02
0.01
Balance
and at constant initial strain rate ranging from 10 4 to 3 /102 s1. It is to be noted that, following straining to a desired strain or to fracture, the samples were unloaded immediately and were: (1) quenched in water to preserve the microstructure for transmission election microscope (TEM) observation; or (2) cooled down to room temperature using force cooling, instead of quenching in water, to preserve clean surfaces for scanning electron microscope (SEM) observation. Optical microscope with image analyzer, SEM, and TEM were engaged in the observation of the microstructure evolution of the materials during superplastic deformation. The TEM specimens were taken from the center part of the tensile samples on the longitudinal / transverse (L /T) section, which were then thinned using dimple polishing to 10 /20 mm and perforated using ion milling with ion beam angle inclination of 5/58. For TEM (JEOL JEM-2010), ellipse-shaped samples with major axis parallel to the stress direction were used, enabling the relationship between the grain boundary direction and the stress direction to be determined. Electron BackScatter Diffraction (EBSD) test was performed using the JEOL JSM 5410 SEM with the attachment of an Oxford Link Opal EBSD test system to determine the grain boundary misorientation distribution of the material. The surfaces of the EBSD samples were prepared by polishing down to 0.06 mm colloidal silica slurry and etching with Keller’s reagent for 10 /50 s. EBSD was also done on the center part of the tensile sample on the L /T section. An average of 300 observations were collected for determining the misorientation distribution.
3. Results and discussions 3.1. Initial microstructure and elongation to failure The metallographic examination combined with image analysis showed that the as-received material had uniform grain size and shape (Fig. 1). The linear intercepts of grains measured along longitudinal, transverse, short transverse directions were 10, 10.1, and 7.2 mm, respectively. The elongations to failure, o f, of samples with cutting direction parallel and perpendicular to the rolling direction are shown in Fig. 2. The elongations of the L-samples were consistently higher than that of the T-samples in the whole test strain rate range. There is a relatively large difference between the
Fig. 1. Microstructure of as received Al 7475 sheet, on rolling and transverse section.
Fig. 2. Elongations of samples as a function of initial strain rate at temperature of 516 8C.
elongations of L- and T-samples at strain rate ranging from 10 4 to 3 /103 s 1. When the strain rate is above 102 s 1, the difference decreased. This is in good agreement with results reported by Shin and Park [4] on an Al 7075 alloy with average grain size of 12 mm. Fig. 3 shows the flow stress of the samples deformed at different initial strain rates along the L- and Tdirections to 100% strain. The strain rate sensitivity index, m , was about 0.6 when the strain rate is below 102 s 1 for both L- and T-samples, and it began to decrease when the strain rate is greater than 10 2 s 1. The lower elongation of the T-samples when the strain rate is lower than 102 s 1 may be partially attributed to the higher stress needed for deformation along transverse direction. The relatively large difference
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Fig. 3. Illustration of the variation of flow stresses at strain of 100% as a function of initial strain rate for the samples tested at 516 8C.
between the elongations of the L- and T-samples indicated that the microstructure of the materials was anisotropic in L- and T-directions. At high strain rates, the materials deformed non-superplastically and the influence of anisotropy microstructure may not be important. 3.2. Distribution of misorientation population The misorientation of grain boundaries (and interphase boundary property) of the polycrystalline materials can influence polycrystalline strength, ductility and corrosion behaviour, etc. [16 /20]. In the regime of superplasticity, high-angle grain boundaries are usually desired for achieving grain boundary sliding (GBS) as GBS is the dominant mechanism of superplasticity. In this work, the distributions of the contiguous grain boundary misorientation were characterized. Results of the as-received material and the L-sample after straining
Fig. 4. Misorientation distributions of the grain boundaries of the asreceived material: (a) parallel to transverse direction, (b) parallel to the longitudinal direction; and of the sample strained to 100% at an initial strain rate 10 3 s 1 at 516 8C; (c) parallel to the transverse direction; (d) parallel to the longitudinal direction.
to 100% at 516 8C, at an initial strain rate of 10 3 s 1, are shown in Fig. 4. The misorientation distribution of the as-received Al 7475 showed that the high-angle grain boundaries were in the majority to that of the total grain boundary population. It is also noted that the low-angle grain boundaries were more frequently present parallel to the transverse direction compared to the rolling direction, as showed in Fig. 4(a and b). This is an interesting point since anisotropic high-angle boundary distribution will result in different degrees and ease of GBS when strained along the T- or L-directions. The low-angle grain boundaries are not as mobile until they evolve into high-angle boundaries. This means that higher stress is needed to facilitate GBS on low-angle boundaries. During superplastic deformation, GBS takes place more on the grain boundaries along the stress axis than those along the transverse direction. The frequency of low-angle boundaries parallel to stress direction may thus determine the average stress needed to generate superplastic flow. For the current material, the high frequency of low-angle boundaries present parallel to the transverse direction of the material account for the larger flow stress needed when strained along the Tdirection (see Fig. 3). These low-angle boundaries gradually evolved into high-angle boundaries and became more mobile due to the accumulation of dislocations during the earlier stage of superplastic deformation [21]. This is evident from the comparison of Fig. 4(a /d). 3.3. Precipitates, thickness of the liquid phase and DFZ The degree of misorientation is most frequently used as a simple but important description of grain boundary character [17 /20]. However, for QSP materials deformed at the high temperature, other aspects like the distribution of precipitates and liquid phase (if it exists at the test temperature) along grain boundary are very important. The presence of precipitates and/or liquid phases along grain boundaries will greatly change the grain boundary character. 3.3.1. Precipitates and liquid phase TEM observations of the as-received material showed that precipitates are mainly present at or near grain boundaries parallel to the rolling direction, and are relatively sparse within the grains (Fig. 5(a)). This indicated that Zn, Mg, and Cu etc. solute atoms dissolved relatively uniformly in the matrix. The formation of such microstructures can be attributed to the thermal mechanical treatment (TMT). The large particles precipitated during TMT act as recrystallization nucleation sites as a consequence of the large misorientation of grain boundaries generated around them during cold deformation [22]. More high-angle grain boundaries thus might form along the rolling direction.
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Fig. 5. TEM micrographs of: (a) particles present on grain boundaries of as-received state. Relatively more precipitates are present on grain boundaries parallel (marked A) to the rolling direction than that of on the perpendicular direction ones (marked B). C is a pre-existed cavity. (b) Eutectic phase (marked L) present after 50% strain at 516 8C, 10 3 s 1.
These high-angle grain boundaries may subsequently be preferential sites for the small particles to precipitate in the latter stages of recrystallization. After the material has experienced strains 50% beyond, a semi-continuous eutectic phase along the grain boundaries together with copious precipitates inside the grains were present (Fig. 5(b)). This microstructure was similar to the microstructure reported by Dang et al. [23] on a 7000 series aluminum alloy. In that work, the eutectic phase was considered to be a quaternary eutectic phase of melting point of 477 8C, which is significantly lower than the melting point of the bulk material. EDS analysis on the eutectic phase present revealed that it contains high Zn, Mg and Al content and the atomic ratio of Zn /Mg /Al is close to 1:1:1. Considering the eutectic reactions of Liquid 0/ Al/MgZn2, Liquid0/Al/MgZn2/Mg3Zn2Al2 at 475 8C and Liquid 0/Al/Mg3Zn2Al2 at 489 8C [24], at 516 8C, this phase can be expected to be in the liquid state (hereafter, the term ‘liquid phase’ is used to refer the eutectic phase at deformation temperature). The atomic ratio revealed by EDS was in good agreement with these eutectic reactions. Inspections of the microstructure showed that the eutectic phase was not uniformly distributed. The eutectic phase was present in more copious amounts along the grain boundaries parallel to the rolling direction than in the other directions. This was similar to the distribution of the precipitates along the grain boundaries in the as-received material. This indicated that the formation of the eutectic phase is closely related to the original distribution of the precipitates. The precipitates may be fully or partially dissolved during the heat-up and early stages of deformation, and this results in solute segregation along grain boundaries. The segregation of solutes leads to the presence of low
melting phase along the grain boundaries. This liquid phase later serves as the source material for the filament. Discussion on the filament formation can be found elsewhere [15]. Once the liquid phase is present along the grain boundaries, the bonding strength of the grain boundaries will be reduced greatly [25,26]. Takayama et al. [24] have deduced a relationship between the flow stress and the thickness of liquid phase, s
3p h A0 vl¯ 8 h3 L 0
;
(1)
where h is the viscosity of the liquid; h, thickness of the liquid film; A0, cross sectional area of the specimen; L0, gauge length; l;¯ mean linear intercept of grains; and n, tensile velocity. For an approximate evaluation, replacing the viscosity of pure aluminum at melting point 1.30 nNs m2 as pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi the h of liquid phase, l¯9 m ( 3 d1 d2 d3 ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi p 3 1010:17:2); flow stress of 1.11 MPa at tensile velocity 15 /106 m s 1 (10 3 s 1) and other corresponding values into (Eq. (1)), the thickness of the liquid phase which can sustain the flow stress is about 46 nm. The thickness could be larger since the liquid phase is not present on all the grain boundaries and the viscosity of the liquid should be higher at the test temperature, as have been discussed elsewhere [15]. Two inferences therefore can be reached: (1) increasing the thickness of the liquid phase will significantly reduce the grain boundary bound strength as s 8/h3; and (2) when the ratio of the solid /solid contact along the grain boundaries decreases to a certain degree, grain boundary separation may take place, assuming that the liquid thickness is equal to the eutectic phase thickness, which is around 100/300 nm as shown in Fig. 5(b).
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3.3.2. Dispersoid free zone DFZs have been observed in 7xxx Al alloys after superplastic deformation. The formation of DFZs usually occurs in the grain boundaries approximately perpendicular to the stress direction [4,9]. In this work, large DFZs can be observed along the grain boundaries perpendicular to the stress direction, although there exist some relatively narrow DFZs at some grain boundaries parallel to the stress direction as shown in Fig. 6. It is also noted that DFZs generally formed on one side of the grain boundaries. This phenomenon can be explained by the relaxation of the stress, due to the softening of the matrix when DFZs are formed on one side of the grain boundaries. The driving force for DFZ formation on the other side of the grain boundaries is then reduced. Blandin et al. [9] suggested that the formation of DFZs could be the direct-result of the tensile stress acting on the grain boundaries perpendicular to the stress direction. DFZs present on the other grain boundaries can be due to the grain rotation and switching after large strains. However, in the present work, DFZs on the grain boundaries that are not perpendicular to the tensile axis were also observed in the early stages of deformation as shown in Fig. 6. The existence of DFZ on certain grain boundaries parallel to the stress direction indicated that, except for the effect of grain rotation, these grain boundaries could also be under tensile stress so as to accommodate GBS that took place on the neighbouring grains. This tensile stress is also the driving force for the debonding that took place on the grain boundaries parallel to the stress direction, if the bonding strength of these boundaries is decreased to a certain level when enough liquid phase is present.
The formation of DFZs may interact with the cavitation behaviour. The possibility of new cavity nucleation in DFZs may decrease as a result of lack of dispersoids in DFZs [4]. However, this mechanism is important only at the later stages of deformation. In the early stages of deformation, the stress concentration on the grain boundaries may cause either the nucleation of cavities or the formation of DFZs adjacent to the grain boundaries. Formation of either cavity nuclei on the grain boundaries or DFZs in the vicinity of the grain boundaries will lower the potential for further formation of cavities or DFZs. This effect is evidenced by the observation that there were few DFZs formed next to cavities, and vice versa, since either process can relieve stress concentration. Whether stress concentration is relieved by the cavity nucleation or DFZ formation may depend on the different bonding strengths of the different grain boundaries. The DFZs behave similarly to the soft phase in MD superplastic materials like Ti/ 6Al /4V, which suffers relatively less from cavitation. 3.4. Modelling of preferential cavity stringer formation Based on these mechanisms of cavity stringer formation discussed previously [4,10,11,13,15], the following considerations are summarized for developing a model for cavity stringer formation: (1) Only those cavities nucleated at precipitates on the grain boundaries and triple junctions are important for the latter stage of cavity growth, coalescence and interlinkage. The likelihood of coalescence and interlinkage of cavities nucleated within the grain is low and is not considered here. (2) GBS takes place extensively along the grain boundaries but is inhomogeneous due to the anisotropic GBCD. Cavities are nucleated at sites where high stress concentrations are induced by GBS (i.e. when GBS can
Fig. 6. TEM micrographs show of the presence of high Zn /Mg /Al content eutectic phase along the grain boundaries (marked L) and DFZ (marked D and A) of the L-sample strained to 50% at 516 8C at 10 3 s 1. Small DFZ can be found at the grain boundary parallel to stress direction, marked A in (a). Stress and rolling directions are shown by arrow in pictures.
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not be well accommodated by diffusion flow or dislocation slip etc.) and where the local stress exceeds the particle /boundary or the boundary/boundary bonding strength. Thus, inhomogeneous GBS also results in inhomogeneous cavity nucleation. (3) Under superplastic flow conditions, when the ratio of the grain boundary area covered by liquid phase to the liquid phase-free area along the grain boundary exceeds a certain value, grain boundary debonding will take place. The presence of liquid phase on the boundaries not only enhances the GBS, but also increases its potential for debonding. For the material used in this work, liquid phase was present frequently along the grain boundaries parallel to the rolling direction of the material. On the other hand, the liquid phase can only be found at a few boundaries perpendicular to the rolling direction. This suggests that debonding happens mainly on the grain boundaries parallel to the rolling direction, whilst some debonding may also take place on the grain boundaries perpendicular to the rolling direction. Therefore, considering a hexagonal grain structure configuration as shown in Fig. 7, two ways of cavity nucleation and development
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can be expected when it is deformed along the L- or Tdirections. 3.4.1. Mode I For the T-sample, the particles and/or liquid phase are present preferentially at the grain boundaries perpendicular to the stress direction. At 516 8C, the boundaries become easily separated during deformation due to the presence of large quantity of liquid phase. Cavities are nucleated here and can then propagate intergranularly. With an increasing number of the separated grain boundaries, the cavities can finally propagate along the grain boundaries and interlink with each other along the traverse direction, i.e. on the plane of the maximum principle stress, or along 458 direction, i.e. on the plane of the maximum shear stress. This cavity propagation mode results in the cavity stringers along 90 or 458 direction to stress axis. A schematic illustration of this mechanism is shown as Mode I, in Fig. 7. 3.4.2. Mode II For the L-sample, most of the grain boundaries parallel to the stress direction are covered by liquid phase at 516 8C. This means that not only GBS, but also debonding can take place on those grain boundaries during deformation. On the other hand, most grain boundaries perpendicular to the stress direction have little or no liquid phase, while only a few boundaries are covered by the liquid phase. This means that although GBS can take place along the grain boundaries perpendicular to the stress direction, debonding can take place only on a few of those grain boundaries where a large amount of liquid phase is present. Supposing the grain boundaries perpendicular to the stress direction, as shown in Mode II in Fig. 7, A /A? is the only boundary where debonding can take place. During deformation, this cavity can then develop the along stress direction by GBS on the other boundaries perpendicular to the stress direction, and both GBS and debonding occur on the grain boundaries inclined to the stress direction, as illustrated sequentially in Mode II (a), (b) and (c) in Fig. 7. In this process, the grain boundaries that were originally perpendicular to the stress direction are kept connected throughout except boundary A /A?. The grain boundaries that are inclined to the stress direction can be separated to facilitate the continuous deformation. The result of this procedure is the formation of a long and narrow cavity marked in black (A /A?) of Mode II (c) in Fig. 7. Long cavity stringers can then form as this process is repeated. 3.5. Experimental evidences
Fig. 7. Illustration of the procedure of anisotropic grain boundary debonding and GBS related cavity stringer formation (stressed vertically).
To investigate the cavity stringer behaviour in superplastic deformation, the observations in this work were
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limited to the samples tested at initial strain rates of 3/ 104 and 10 3 s 1 at 516 8C, which are the optimum or near optimum superplastic test conditions for both the L- and T-samples. Inspections of the cavity characteristics in the center portion of the specimens strained to fracture showed that there exist extensive cavity stringers in both the L- and T-samples. For the Lsamples, the cavity stringers are formed parallel to the stress direction as shown in Fig. 8. For the T-samples, they are generally formed in a direction perpendicular, or near 458, to the stress direction (Fig. 8). Two types of preferential cavity stringer were observed on the surfaces of the T- and L-samples, as shown in Fig. 9(a and b), respectively. Numerous short filaments were found on the separated grain boundaries which are perpendicular to the stress direction on the Tsamples. This clearly shows that, since most of the grain boundaries perpendicular to the stress direction can be separated easily when there is enough liquid phase, the cavity stringers will form preferentially along the transverse direction in the T-samples, as described in Mode I.
The stringer formation in Mode II involves a sequence of grain position configurations at different stages of deformation. By observing the L-samples (Fig. 9(b)), it is evident that some grain boundaries perpendicular to stress direction could have been separated earlier whilst others remain connected. As shown in Fig. 9(b), some long and continuous filaments link distant grains A and D, and grains C and E in a narrow and long surface cavity, whilst some short filaments link grains B and C, and grains A and C. The continuous filaments show that these grains were originally contiguous. The original grain position configuration can be deduced to be similar to the one in Fig. 9(c-i). It is clearly shown that this cavity stringer does not result simply from the boundary debonding and cavity plastic deformation or the cavity interlinking along stress direction. This is a cooperative procedure. The important factor in this procedure is the GBCD, especially its grain boundary bonding strength at the test temperature. This leads to the development of cavity stringer as proposed in Mode II. It is necessary to emphasize that the above model does not preclude the contribution of superplastic flow, cavity
Fig. 8. Cavity stringers of samples strained to fracture at 516 8C, at an initial strain rate of 10 3 s 1: (a and b) L-sample, 850%; (c and d) T-sample, 490%; (a and c) on sample L-T surface; (b and d) in center part of the sample on ST-L section. Stress direction is horizontal.
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Fig. 9. Surface morphologies of the samples after superplastic deformation at 516 8C, 10 3 s 1: (a) the T-sample strained to fracture; (b) the Lsample after fracture; (c) schematic illustration of the configuration variation of the grains shown in (b). The long continuous filament AD and CE showed that the originally contiguous grains A and D, C and E were far separated, while the originally contiguous grains B and C were connected via a relatively short filament BC. Stress direction is horizontal.
coalescence and interlinkage, etc. However, they are contributors and not the dominant mechanism. These factors will become important only when boundary or interface are bound well, as in the case of some MD materials [11,27], or they become insignificant for the low boundary strength materials, e.g. some of superplastic ceramics and intermetallics [25].
4. Conclusions The investigation was carried out on the directional cavity stringer formation in superplastically deformed Al 7475. The results obtained are as follows: (1) The anisotropy of GBCD of the Al 7475 resulted in a larger flow stress needed for deformation along the transverse direction (T-sample) compared to that of the longitudinal direction (L-sample). The larger flow stress is attributed to the lower elongation of the T-samples. The lower elongation of the T-samples is also due to the preferential cavity stringers formed perpendicular to the stress direction. (2) The sliding capability of grain boundaries is inhomogeneous due to the anisotropic GBCD. Cavities are nucleated at sites where high stress concentrations have been induced by GBS and where the local stress exceeds the particle /boundary or the boundary/boundary bonding strength.
(3) There are interactions between the DFZ formation and the cavity nucleation. The DFZ formation or the cavity nucleation may depend on the local stress condition and the grain boundary bonding strength. The formation of cavity nuclei at the grain boundaries will lower the potential of DFZ formation in their vicinity, and vice versa. (4) The GBCD, especially the grain boundary bonding strength at test temperature, is the key factor controlling the formation of cavity stringers. The presence of liquid phase on the grain boundaries significantly reduces the bonding strength. The formation of preferential cavity stringer is closely related to the GBCD, GBS and grain boundary debonding. Based on this, a model of cavity stringer formation is proposed.
References [1] C.C. Bamption, J.W. Edington, Metall. Trans. 13A (1982) 1721. [2] H. Iwasaki, T. Mori, T. Tagata, M. Matsuo, K. Higashi, Mater. Sci. Forum 233 /234 (1997) 81. [3] S. Sagat, P.A. Blenkinsop, D.M.R. Taplin, J. Inst. Metals 100 (1972) 268. [4] H. Shin Dong, Park Kyung-Tae, Mater. Sci. Eng. A268 (1999) 55. [5] J. Pilling, N. Ridley, Acta Metall. 34 (1986) 669. [6] A. Varloteaux, J.J Blandin, M. Suery, Mater. Sci. Technol. 5 (1989) 1109.
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[7] Xun Yuwei, Zhun Yiyuan, Mao Wenfeng, Cui Jianzhong, J. Mater. Proc. Technol. 72 (1997) 183. [8] W.D. Cao, X.P. Lu, H. Conrad, Acta Mater. 44 (1996) 697. [9] J.J. Blandin, B. Hong, A. Varloteaux, M. Suery, G. L’esperanve, Acta Mater. 44 (1996) 2317. [10] X. Jiang, J. Chui, L. Ma, Acta Metall. Mater. 41 (1993) 2721. [11] A. Yousefini, J.C. Earthman, F.A. Mohamed, Acta Mater. 46 (1998) 3557. [12] M.G. Zelin, H.S. Yang, R.Z. Valiev, A.K. Mukherjee, Metall. Trans. 24A (1993) 417. [13] M.G. Zelin, A.K. Mukherkee, Acta Metall. Mater. 43 (1995) 2359. [14] A. Chakraborty, J.C. Earthman, Acta Mater. 45 (1997) 4615. [15] C.L. Chen, M.J. Tan, Mater. Sci. Eng. A298 (2001) 235. [16] R. Raj, Acta Metall. 26 (1978) 995. [17] V. Randle, Mater. Sci. Forum 294 /296 (1999) 51.
[18] V. Randle, Acta Mater. 45 (1999) 4187. [19] Y. Pan, B.L. Adams, T. Olson, N. Panayotou, Acta Mater. 44 (1996) 4685. [20] Liu Hongwei, Gao Ming, G. Harlow David, P. Wei Robert, Scripta Metall. Mater. 32 (1995) 1807. [21] J. Liu, D.J. Chakrabarti, Acta Mater. 44 (1996) 4647. [22] G.J. Mahon, D. Warrington, R.G. Butler, R. Grimes, Mater. Sci. Forum 170 /172 (1994) 187. [23] M.C. Dang, J.J. Blandin, B. Baudelet, Acta Mater. 44 (1996) 3991. [24] Y. Takayama, T. Tozawa, H. Kato, Acta Mater. 47 (1999) 1263. [25] W.J. Kim, J. Wolfenstine, O.D. Sherby, Acta Metall. Mater. 39 (1991) 199. [26] M.C. Dang, B. Baudelet, J. Mater. Sci. 29 (1994) 2315. [27] K.T. Park, S.T. Yang, J.C. Earthman, F.A. Mohamed, Mater. Sci. Eng. A188 (1994) 59.
Effect of grain boundary character distribution (GBCD) on the cavitation behaviour during superplastic deformation of Al 7475 C.L. Chen, M.J. Tan * School of Mechanical and Production Engineering, Nanyang Technological University, Singapore 639798, Singapore Received 28 November 2001; received in revised form 25 January 2002
Abstract The formation of cavity stringers in superplastically deformed Al 7475 alloy was investigated in this work. After superplastic deformation, cavity stringers have a tendency to form intergranularly along grain boundaries that are parallel to the rolling direction of the material. The microstructure and microtexture showed that the material has anisotropic grain boundary character distribution (GBCD). In the as-received state, small precipitates were found to be distributed preferentially along the grain boundaries which are parallel to the rolling direction. After superplastic deformation, a low melting point eutectic phase was found along the grain boundaries, which are also mainly parallel to the rolling direction. At the test temperature of 516 8C, the eutectic phase is expected to be in the liquid state. The presence of the liquid phase reduces the grain boundary bonding strength. It is proposed that the formation of cavity stringers is closely related to the GBCD. Based on the anisotropic distribution of GBCD and extensive grain boundary sliding during superplastic deformation, a model is proposed to explain the formation of cavity stringers. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Al 7475; Superplasticity; Cavity stringers; Grain boundary character distribution (GBCD)
1. Introduction Intergranular cavitation has been frequently reported in a wide range of superplastic materials [1 /4], while some other alloy systems, such as titanium based alloys, suffer relatively less from this deleterious phenomenon. Cavities generally nucleate at triple junctions of grain boundaries, in the vicinity of particles and even on ledges of grain boundaries, and/or directly from preexisting defects. Cavity nuclei can then grow, coalesce and interlink as deformation continues, leading to the degradation of the postdeformation properties of the materials and may even result in premature fracture during processing. The effects of strain rate, strain, temperature, back pressure, etc. on cavitation behaviour have been extensively studied [1,5 /9], and these reported general common features for many alloy systems. The effects of the microstructure of the materials on cavitation
* Corresponding author. Tel.: /65-790-5582; fax: /65-791-1859. E-mail address: [email protected] (M.J. Tan).
behaviour, on the other hand, seems more complicated. Small chemical compositional differences or subtle differences of solidification and/or thermal processing conditions between batches of the same materials can result in a large difference in cavitational behaviour [1]. One of the interesting phenomena during superplastic deformation is the formation of cavity stringers where a group of cavities are aligned in a specific direction, and this has been reported both in microduplex (MD) and quasi-single phase (QSP) superplastic materials [4,10,11]. Yousefini et al. [11] reviewed this phenomenon recently and found that in most cases, the cavity stringers align themselves parallel to the tensile axis; and in some cases little or no arrangements can be discerned, and on a few occasions the cavities show an alignment perpendicular to the tensile axis. The morphology of cavity stringers varies from one material to another and is affected by the superplastic forming conditions such as strain, strain rate and temperature. Zelin et al. [12,13] proposed a model to depict the possibility of cavity stringer formation based on the concept of cooperative grain boundary sliding (CGBS). They suggested that during sheet forming under a
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biaxial state of stress, cavity nucleation occurs in the intersecting sites of cooperative sliding surfaces, which operates in a direction close to that of the principal shear stress and this thereby aligns the stringers along the same orientation. This was consistent with their observation on both biaxial and uniaxial tensile specimens. Chakraborty and Earthman [14] carried out the numerical analysis on the interaction between boundary sliding and cavitation. In the absence of accommodation by diffusion, deformation, and/or boundary migration, the local stress concentration is relieved by the opening of a cavity. According to their calculations, the opening of such a cavity results in the stressing of the boundaries inclined to the stress axis near this cavity more than the transverse boundaries, leading to the further nucleation of cavities at these boundaries. Short rows of cavities inclined to the tensile axis can then result. Yousefini et al. [11] furthered Chakraborty and Earthman’s work [14] and proposed, based on the results of Zn /22wt.%Al materials, that the formation of cavity stringers is related to: (1) the favourable nucleation sites on former a boundaries which are related to the impurity segregation and thermal history of the material; and (2) their change in orientation on approaching the tensile axis as deformation proceeds. This latter factor, as a result of superplastic flow, plays an important role in the development of cavity stringers, since the evolution of cavity stringers is controlled by the factors related to the nature and characteristics of superplasticity rather than the cavity nucleation process, once the cavities are nucleated. It is important to mention that, for the MD materials, the cavity stringers are usually aligned parallel to the tensile axis regardless of the rolling direction. For some MD materials and for most QSP materials, the formation of cavity stringers along the rolling direction has been reported. This phenomenon is consistent with the fact that, during prior thermomechanical treatment, most superplastic materials are rolled either to produce the fine grain size necessary for superplastic deformation or to be a part of the processing procedure. Large particles, if present, can therefore break down and align as small particles along the rolling direction. Since superplastic materials are commonly tested with tensile axis parallel to the rolling direction, it was suggested that cavity nucleation at these particles leads to the formation of cavity stringers [4,10]. This suggestion was verified by the observation that changing the orientation of the test samples led to a corresponding change in the arrangement of cavity stringers [4,10]. Jiang et al. [10] and Shin and Park [4] have reported the directionally aligned cavity stringer in the finegrained Al 7075 alloy. In addition to the cavity nucleation effects, Jiang et al. [10] also concluded that the tensile superplastic deformation contributes to the
formation of the cavity stringers aligned along tensile axis. Their explanation is similar to that of Yousefini et al. [11]. Shin and Park [4] studied both fine and coarse grain Al 7075 alloys and found tensile direction-aligned stringers in fine grain materials and rolling directionaligned stringers in large grain materials. They found that the cavity stringer formation parallel to the tensile axis could be associated with the existence of dispersoid free zone (DFZ) which was formed in the vicinity of grain boundaries perpendicular to the tensile axis during superplastic deformation. The formation of DFZs is frequently observed and it may lead to the decrease of new cavity nucleation in their vicinity. These works are helpful in the understanding of the stringer formation mechanism. However, there are still some aspects that have yet not been well clarified: (1) When directional precipitate stringers appear, what mechanisms cause these cavities to interlink with each other along precipitate stringers, regardless of the stress direction? (2) When cavities were uniformly nucleated, such as the cases of QSP materials with no precipitate stringer, and MD materials, what mechanisms lead cavities to develop along stress direction, instead of along the transverse or 458 direction which usually takes place in the room temperature tensile condition? (3) What factors dominate the cavity stringer direction? These ambiguities in the stringer formation are related to the cavity growth, coalescence and interlinkage mechanisms. Studies on the cavity stringer behaviour are important so as to achieve the optimum microstructure for superplasticity. It is widely recognised that the grain boundary character has great impact on the superplasticity of the material, as a number of grain boundary phenomena are involved during superplastic deformation. In a previous work [15] on a typical QSP material, Al 7475, preferential cavity interlinkage of the material during superplastic deformation was reported. The present work focuses on the effect of GBCD on the preferential cavity interlinkage.
2. Experimental procedures Al 7475 sheet with chemical composition listed in Table 1 was used for superplastic deformation testing. The as-received sheet was in T4 condition (solution heat treated, quenched, and naturally aged). The tensile specimens were of 15 mm length, 4 mm width and 2 mm thickness in gauge parts, with the specimen cutting directions either parallel (L-sample) or perpendicular (T-sample) to the rolling direction of the sheet. The surfaces of the tensile sample gauge part were polished down to 1 mm diamond paste. Superplastic uniaxial tensile test and cavitation ratio determination have been detailed elsewhere [15]. In the present work, tensile tests were performed at 516 8C
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Table 1 Chemical composition of the alloy Element
Zn
Mg
Cu
Cr
Fe
Si
Ti
Mn
Al
Wt.%
5.64
2.34
1.58
0.19
0.08
0.05
0.02
0.01
Balance
and at constant initial strain rate ranging from 10 4 to 3 /102 s1. It is to be noted that, following straining to a desired strain or to fracture, the samples were unloaded immediately and were: (1) quenched in water to preserve the microstructure for transmission election microscope (TEM) observation; or (2) cooled down to room temperature using force cooling, instead of quenching in water, to preserve clean surfaces for scanning electron microscope (SEM) observation. Optical microscope with image analyzer, SEM, and TEM were engaged in the observation of the microstructure evolution of the materials during superplastic deformation. The TEM specimens were taken from the center part of the tensile samples on the longitudinal / transverse (L /T) section, which were then thinned using dimple polishing to 10 /20 mm and perforated using ion milling with ion beam angle inclination of 5/58. For TEM (JEOL JEM-2010), ellipse-shaped samples with major axis parallel to the stress direction were used, enabling the relationship between the grain boundary direction and the stress direction to be determined. Electron BackScatter Diffraction (EBSD) test was performed using the JEOL JSM 5410 SEM with the attachment of an Oxford Link Opal EBSD test system to determine the grain boundary misorientation distribution of the material. The surfaces of the EBSD samples were prepared by polishing down to 0.06 mm colloidal silica slurry and etching with Keller’s reagent for 10 /50 s. EBSD was also done on the center part of the tensile sample on the L /T section. An average of 300 observations were collected for determining the misorientation distribution.
3. Results and discussions 3.1. Initial microstructure and elongation to failure The metallographic examination combined with image analysis showed that the as-received material had uniform grain size and shape (Fig. 1). The linear intercepts of grains measured along longitudinal, transverse, short transverse directions were 10, 10.1, and 7.2 mm, respectively. The elongations to failure, o f, of samples with cutting direction parallel and perpendicular to the rolling direction are shown in Fig. 2. The elongations of the L-samples were consistently higher than that of the T-samples in the whole test strain rate range. There is a relatively large difference between the
Fig. 1. Microstructure of as received Al 7475 sheet, on rolling and transverse section.
Fig. 2. Elongations of samples as a function of initial strain rate at temperature of 516 8C.
elongations of L- and T-samples at strain rate ranging from 10 4 to 3 /103 s 1. When the strain rate is above 102 s 1, the difference decreased. This is in good agreement with results reported by Shin and Park [4] on an Al 7075 alloy with average grain size of 12 mm. Fig. 3 shows the flow stress of the samples deformed at different initial strain rates along the L- and Tdirections to 100% strain. The strain rate sensitivity index, m , was about 0.6 when the strain rate is below 102 s 1 for both L- and T-samples, and it began to decrease when the strain rate is greater than 10 2 s 1. The lower elongation of the T-samples when the strain rate is lower than 102 s 1 may be partially attributed to the higher stress needed for deformation along transverse direction. The relatively large difference
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Fig. 3. Illustration of the variation of flow stresses at strain of 100% as a function of initial strain rate for the samples tested at 516 8C.
between the elongations of the L- and T-samples indicated that the microstructure of the materials was anisotropic in L- and T-directions. At high strain rates, the materials deformed non-superplastically and the influence of anisotropy microstructure may not be important. 3.2. Distribution of misorientation population The misorientation of grain boundaries (and interphase boundary property) of the polycrystalline materials can influence polycrystalline strength, ductility and corrosion behaviour, etc. [16 /20]. In the regime of superplasticity, high-angle grain boundaries are usually desired for achieving grain boundary sliding (GBS) as GBS is the dominant mechanism of superplasticity. In this work, the distributions of the contiguous grain boundary misorientation were characterized. Results of the as-received material and the L-sample after straining
Fig. 4. Misorientation distributions of the grain boundaries of the asreceived material: (a) parallel to transverse direction, (b) parallel to the longitudinal direction; and of the sample strained to 100% at an initial strain rate 10 3 s 1 at 516 8C; (c) parallel to the transverse direction; (d) parallel to the longitudinal direction.
to 100% at 516 8C, at an initial strain rate of 10 3 s 1, are shown in Fig. 4. The misorientation distribution of the as-received Al 7475 showed that the high-angle grain boundaries were in the majority to that of the total grain boundary population. It is also noted that the low-angle grain boundaries were more frequently present parallel to the transverse direction compared to the rolling direction, as showed in Fig. 4(a and b). This is an interesting point since anisotropic high-angle boundary distribution will result in different degrees and ease of GBS when strained along the T- or L-directions. The low-angle grain boundaries are not as mobile until they evolve into high-angle boundaries. This means that higher stress is needed to facilitate GBS on low-angle boundaries. During superplastic deformation, GBS takes place more on the grain boundaries along the stress axis than those along the transverse direction. The frequency of low-angle boundaries parallel to stress direction may thus determine the average stress needed to generate superplastic flow. For the current material, the high frequency of low-angle boundaries present parallel to the transverse direction of the material account for the larger flow stress needed when strained along the Tdirection (see Fig. 3). These low-angle boundaries gradually evolved into high-angle boundaries and became more mobile due to the accumulation of dislocations during the earlier stage of superplastic deformation [21]. This is evident from the comparison of Fig. 4(a /d). 3.3. Precipitates, thickness of the liquid phase and DFZ The degree of misorientation is most frequently used as a simple but important description of grain boundary character [17 /20]. However, for QSP materials deformed at the high temperature, other aspects like the distribution of precipitates and liquid phase (if it exists at the test temperature) along grain boundary are very important. The presence of precipitates and/or liquid phases along grain boundaries will greatly change the grain boundary character. 3.3.1. Precipitates and liquid phase TEM observations of the as-received material showed that precipitates are mainly present at or near grain boundaries parallel to the rolling direction, and are relatively sparse within the grains (Fig. 5(a)). This indicated that Zn, Mg, and Cu etc. solute atoms dissolved relatively uniformly in the matrix. The formation of such microstructures can be attributed to the thermal mechanical treatment (TMT). The large particles precipitated during TMT act as recrystallization nucleation sites as a consequence of the large misorientation of grain boundaries generated around them during cold deformation [22]. More high-angle grain boundaries thus might form along the rolling direction.
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Fig. 5. TEM micrographs of: (a) particles present on grain boundaries of as-received state. Relatively more precipitates are present on grain boundaries parallel (marked A) to the rolling direction than that of on the perpendicular direction ones (marked B). C is a pre-existed cavity. (b) Eutectic phase (marked L) present after 50% strain at 516 8C, 10 3 s 1.
These high-angle grain boundaries may subsequently be preferential sites for the small particles to precipitate in the latter stages of recrystallization. After the material has experienced strains 50% beyond, a semi-continuous eutectic phase along the grain boundaries together with copious precipitates inside the grains were present (Fig. 5(b)). This microstructure was similar to the microstructure reported by Dang et al. [23] on a 7000 series aluminum alloy. In that work, the eutectic phase was considered to be a quaternary eutectic phase of melting point of 477 8C, which is significantly lower than the melting point of the bulk material. EDS analysis on the eutectic phase present revealed that it contains high Zn, Mg and Al content and the atomic ratio of Zn /Mg /Al is close to 1:1:1. Considering the eutectic reactions of Liquid 0/ Al/MgZn2, Liquid0/Al/MgZn2/Mg3Zn2Al2 at 475 8C and Liquid 0/Al/Mg3Zn2Al2 at 489 8C [24], at 516 8C, this phase can be expected to be in the liquid state (hereafter, the term ‘liquid phase’ is used to refer the eutectic phase at deformation temperature). The atomic ratio revealed by EDS was in good agreement with these eutectic reactions. Inspections of the microstructure showed that the eutectic phase was not uniformly distributed. The eutectic phase was present in more copious amounts along the grain boundaries parallel to the rolling direction than in the other directions. This was similar to the distribution of the precipitates along the grain boundaries in the as-received material. This indicated that the formation of the eutectic phase is closely related to the original distribution of the precipitates. The precipitates may be fully or partially dissolved during the heat-up and early stages of deformation, and this results in solute segregation along grain boundaries. The segregation of solutes leads to the presence of low
melting phase along the grain boundaries. This liquid phase later serves as the source material for the filament. Discussion on the filament formation can be found elsewhere [15]. Once the liquid phase is present along the grain boundaries, the bonding strength of the grain boundaries will be reduced greatly [25,26]. Takayama et al. [24] have deduced a relationship between the flow stress and the thickness of liquid phase, s
3p h A0 vl¯ 8 h3 L 0
;
(1)
where h is the viscosity of the liquid; h, thickness of the liquid film; A0, cross sectional area of the specimen; L0, gauge length; l;¯ mean linear intercept of grains; and n, tensile velocity. For an approximate evaluation, replacing the viscosity of pure aluminum at melting point 1.30 nNs m2 as pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi the h of liquid phase, l¯9 m ( 3 d1 d2 d3 ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi p 3 1010:17:2); flow stress of 1.11 MPa at tensile velocity 15 /106 m s 1 (10 3 s 1) and other corresponding values into (Eq. (1)), the thickness of the liquid phase which can sustain the flow stress is about 46 nm. The thickness could be larger since the liquid phase is not present on all the grain boundaries and the viscosity of the liquid should be higher at the test temperature, as have been discussed elsewhere [15]. Two inferences therefore can be reached: (1) increasing the thickness of the liquid phase will significantly reduce the grain boundary bound strength as s 8/h3; and (2) when the ratio of the solid /solid contact along the grain boundaries decreases to a certain degree, grain boundary separation may take place, assuming that the liquid thickness is equal to the eutectic phase thickness, which is around 100/300 nm as shown in Fig. 5(b).
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3.3.2. Dispersoid free zone DFZs have been observed in 7xxx Al alloys after superplastic deformation. The formation of DFZs usually occurs in the grain boundaries approximately perpendicular to the stress direction [4,9]. In this work, large DFZs can be observed along the grain boundaries perpendicular to the stress direction, although there exist some relatively narrow DFZs at some grain boundaries parallel to the stress direction as shown in Fig. 6. It is also noted that DFZs generally formed on one side of the grain boundaries. This phenomenon can be explained by the relaxation of the stress, due to the softening of the matrix when DFZs are formed on one side of the grain boundaries. The driving force for DFZ formation on the other side of the grain boundaries is then reduced. Blandin et al. [9] suggested that the formation of DFZs could be the direct-result of the tensile stress acting on the grain boundaries perpendicular to the stress direction. DFZs present on the other grain boundaries can be due to the grain rotation and switching after large strains. However, in the present work, DFZs on the grain boundaries that are not perpendicular to the tensile axis were also observed in the early stages of deformation as shown in Fig. 6. The existence of DFZ on certain grain boundaries parallel to the stress direction indicated that, except for the effect of grain rotation, these grain boundaries could also be under tensile stress so as to accommodate GBS that took place on the neighbouring grains. This tensile stress is also the driving force for the debonding that took place on the grain boundaries parallel to the stress direction, if the bonding strength of these boundaries is decreased to a certain level when enough liquid phase is present.
The formation of DFZs may interact with the cavitation behaviour. The possibility of new cavity nucleation in DFZs may decrease as a result of lack of dispersoids in DFZs [4]. However, this mechanism is important only at the later stages of deformation. In the early stages of deformation, the stress concentration on the grain boundaries may cause either the nucleation of cavities or the formation of DFZs adjacent to the grain boundaries. Formation of either cavity nuclei on the grain boundaries or DFZs in the vicinity of the grain boundaries will lower the potential for further formation of cavities or DFZs. This effect is evidenced by the observation that there were few DFZs formed next to cavities, and vice versa, since either process can relieve stress concentration. Whether stress concentration is relieved by the cavity nucleation or DFZ formation may depend on the different bonding strengths of the different grain boundaries. The DFZs behave similarly to the soft phase in MD superplastic materials like Ti/ 6Al /4V, which suffers relatively less from cavitation. 3.4. Modelling of preferential cavity stringer formation Based on these mechanisms of cavity stringer formation discussed previously [4,10,11,13,15], the following considerations are summarized for developing a model for cavity stringer formation: (1) Only those cavities nucleated at precipitates on the grain boundaries and triple junctions are important for the latter stage of cavity growth, coalescence and interlinkage. The likelihood of coalescence and interlinkage of cavities nucleated within the grain is low and is not considered here. (2) GBS takes place extensively along the grain boundaries but is inhomogeneous due to the anisotropic GBCD. Cavities are nucleated at sites where high stress concentrations are induced by GBS (i.e. when GBS can
Fig. 6. TEM micrographs show of the presence of high Zn /Mg /Al content eutectic phase along the grain boundaries (marked L) and DFZ (marked D and A) of the L-sample strained to 50% at 516 8C at 10 3 s 1. Small DFZ can be found at the grain boundary parallel to stress direction, marked A in (a). Stress and rolling directions are shown by arrow in pictures.
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not be well accommodated by diffusion flow or dislocation slip etc.) and where the local stress exceeds the particle /boundary or the boundary/boundary bonding strength. Thus, inhomogeneous GBS also results in inhomogeneous cavity nucleation. (3) Under superplastic flow conditions, when the ratio of the grain boundary area covered by liquid phase to the liquid phase-free area along the grain boundary exceeds a certain value, grain boundary debonding will take place. The presence of liquid phase on the boundaries not only enhances the GBS, but also increases its potential for debonding. For the material used in this work, liquid phase was present frequently along the grain boundaries parallel to the rolling direction of the material. On the other hand, the liquid phase can only be found at a few boundaries perpendicular to the rolling direction. This suggests that debonding happens mainly on the grain boundaries parallel to the rolling direction, whilst some debonding may also take place on the grain boundaries perpendicular to the rolling direction. Therefore, considering a hexagonal grain structure configuration as shown in Fig. 7, two ways of cavity nucleation and development
249
can be expected when it is deformed along the L- or Tdirections. 3.4.1. Mode I For the T-sample, the particles and/or liquid phase are present preferentially at the grain boundaries perpendicular to the stress direction. At 516 8C, the boundaries become easily separated during deformation due to the presence of large quantity of liquid phase. Cavities are nucleated here and can then propagate intergranularly. With an increasing number of the separated grain boundaries, the cavities can finally propagate along the grain boundaries and interlink with each other along the traverse direction, i.e. on the plane of the maximum principle stress, or along 458 direction, i.e. on the plane of the maximum shear stress. This cavity propagation mode results in the cavity stringers along 90 or 458 direction to stress axis. A schematic illustration of this mechanism is shown as Mode I, in Fig. 7. 3.4.2. Mode II For the L-sample, most of the grain boundaries parallel to the stress direction are covered by liquid phase at 516 8C. This means that not only GBS, but also debonding can take place on those grain boundaries during deformation. On the other hand, most grain boundaries perpendicular to the stress direction have little or no liquid phase, while only a few boundaries are covered by the liquid phase. This means that although GBS can take place along the grain boundaries perpendicular to the stress direction, debonding can take place only on a few of those grain boundaries where a large amount of liquid phase is present. Supposing the grain boundaries perpendicular to the stress direction, as shown in Mode II in Fig. 7, A /A? is the only boundary where debonding can take place. During deformation, this cavity can then develop the along stress direction by GBS on the other boundaries perpendicular to the stress direction, and both GBS and debonding occur on the grain boundaries inclined to the stress direction, as illustrated sequentially in Mode II (a), (b) and (c) in Fig. 7. In this process, the grain boundaries that were originally perpendicular to the stress direction are kept connected throughout except boundary A /A?. The grain boundaries that are inclined to the stress direction can be separated to facilitate the continuous deformation. The result of this procedure is the formation of a long and narrow cavity marked in black (A /A?) of Mode II (c) in Fig. 7. Long cavity stringers can then form as this process is repeated. 3.5. Experimental evidences
Fig. 7. Illustration of the procedure of anisotropic grain boundary debonding and GBS related cavity stringer formation (stressed vertically).
To investigate the cavity stringer behaviour in superplastic deformation, the observations in this work were
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limited to the samples tested at initial strain rates of 3/ 104 and 10 3 s 1 at 516 8C, which are the optimum or near optimum superplastic test conditions for both the L- and T-samples. Inspections of the cavity characteristics in the center portion of the specimens strained to fracture showed that there exist extensive cavity stringers in both the L- and T-samples. For the Lsamples, the cavity stringers are formed parallel to the stress direction as shown in Fig. 8. For the T-samples, they are generally formed in a direction perpendicular, or near 458, to the stress direction (Fig. 8). Two types of preferential cavity stringer were observed on the surfaces of the T- and L-samples, as shown in Fig. 9(a and b), respectively. Numerous short filaments were found on the separated grain boundaries which are perpendicular to the stress direction on the Tsamples. This clearly shows that, since most of the grain boundaries perpendicular to the stress direction can be separated easily when there is enough liquid phase, the cavity stringers will form preferentially along the transverse direction in the T-samples, as described in Mode I.
The stringer formation in Mode II involves a sequence of grain position configurations at different stages of deformation. By observing the L-samples (Fig. 9(b)), it is evident that some grain boundaries perpendicular to stress direction could have been separated earlier whilst others remain connected. As shown in Fig. 9(b), some long and continuous filaments link distant grains A and D, and grains C and E in a narrow and long surface cavity, whilst some short filaments link grains B and C, and grains A and C. The continuous filaments show that these grains were originally contiguous. The original grain position configuration can be deduced to be similar to the one in Fig. 9(c-i). It is clearly shown that this cavity stringer does not result simply from the boundary debonding and cavity plastic deformation or the cavity interlinking along stress direction. This is a cooperative procedure. The important factor in this procedure is the GBCD, especially its grain boundary bonding strength at the test temperature. This leads to the development of cavity stringer as proposed in Mode II. It is necessary to emphasize that the above model does not preclude the contribution of superplastic flow, cavity
Fig. 8. Cavity stringers of samples strained to fracture at 516 8C, at an initial strain rate of 10 3 s 1: (a and b) L-sample, 850%; (c and d) T-sample, 490%; (a and c) on sample L-T surface; (b and d) in center part of the sample on ST-L section. Stress direction is horizontal.
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Fig. 9. Surface morphologies of the samples after superplastic deformation at 516 8C, 10 3 s 1: (a) the T-sample strained to fracture; (b) the Lsample after fracture; (c) schematic illustration of the configuration variation of the grains shown in (b). The long continuous filament AD and CE showed that the originally contiguous grains A and D, C and E were far separated, while the originally contiguous grains B and C were connected via a relatively short filament BC. Stress direction is horizontal.
coalescence and interlinkage, etc. However, they are contributors and not the dominant mechanism. These factors will become important only when boundary or interface are bound well, as in the case of some MD materials [11,27], or they become insignificant for the low boundary strength materials, e.g. some of superplastic ceramics and intermetallics [25].
4. Conclusions The investigation was carried out on the directional cavity stringer formation in superplastically deformed Al 7475. The results obtained are as follows: (1) The anisotropy of GBCD of the Al 7475 resulted in a larger flow stress needed for deformation along the transverse direction (T-sample) compared to that of the longitudinal direction (L-sample). The larger flow stress is attributed to the lower elongation of the T-samples. The lower elongation of the T-samples is also due to the preferential cavity stringers formed perpendicular to the stress direction. (2) The sliding capability of grain boundaries is inhomogeneous due to the anisotropic GBCD. Cavities are nucleated at sites where high stress concentrations have been induced by GBS and where the local stress exceeds the particle /boundary or the boundary/boundary bonding strength.
(3) There are interactions between the DFZ formation and the cavity nucleation. The DFZ formation or the cavity nucleation may depend on the local stress condition and the grain boundary bonding strength. The formation of cavity nuclei at the grain boundaries will lower the potential of DFZ formation in their vicinity, and vice versa. (4) The GBCD, especially the grain boundary bonding strength at test temperature, is the key factor controlling the formation of cavity stringers. The presence of liquid phase on the grain boundaries significantly reduces the bonding strength. The formation of preferential cavity stringer is closely related to the GBCD, GBS and grain boundary debonding. Based on this, a model of cavity stringer formation is proposed.
References [1] C.C. Bamption, J.W. Edington, Metall. Trans. 13A (1982) 1721. [2] H. Iwasaki, T. Mori, T. Tagata, M. Matsuo, K. Higashi, Mater. Sci. Forum 233 /234 (1997) 81. [3] S. Sagat, P.A. Blenkinsop, D.M.R. Taplin, J. Inst. Metals 100 (1972) 268. [4] H. Shin Dong, Park Kyung-Tae, Mater. Sci. Eng. A268 (1999) 55. [5] J. Pilling, N. Ridley, Acta Metall. 34 (1986) 669. [6] A. Varloteaux, J.J Blandin, M. Suery, Mater. Sci. Technol. 5 (1989) 1109.
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C.L. Chen, M.J. Tan / Materials Science and Engineering A338 (2002) 243 /252
[7] Xun Yuwei, Zhun Yiyuan, Mao Wenfeng, Cui Jianzhong, J. Mater. Proc. Technol. 72 (1997) 183. [8] W.D. Cao, X.P. Lu, H. Conrad, Acta Mater. 44 (1996) 697. [9] J.J. Blandin, B. Hong, A. Varloteaux, M. Suery, G. L’esperanve, Acta Mater. 44 (1996) 2317. [10] X. Jiang, J. Chui, L. Ma, Acta Metall. Mater. 41 (1993) 2721. [11] A. Yousefini, J.C. Earthman, F.A. Mohamed, Acta Mater. 46 (1998) 3557. [12] M.G. Zelin, H.S. Yang, R.Z. Valiev, A.K. Mukherjee, Metall. Trans. 24A (1993) 417. [13] M.G. Zelin, A.K. Mukherkee, Acta Metall. Mater. 43 (1995) 2359. [14] A. Chakraborty, J.C. Earthman, Acta Mater. 45 (1997) 4615. [15] C.L. Chen, M.J. Tan, Mater. Sci. Eng. A298 (2001) 235. [16] R. Raj, Acta Metall. 26 (1978) 995. [17] V. Randle, Mater. Sci. Forum 294 /296 (1999) 51.
[18] V. Randle, Acta Mater. 45 (1999) 4187. [19] Y. Pan, B.L. Adams, T. Olson, N. Panayotou, Acta Mater. 44 (1996) 4685. [20] Liu Hongwei, Gao Ming, G. Harlow David, P. Wei Robert, Scripta Metall. Mater. 32 (1995) 1807. [21] J. Liu, D.J. Chakrabarti, Acta Mater. 44 (1996) 4647. [22] G.J. Mahon, D. Warrington, R.G. Butler, R. Grimes, Mater. Sci. Forum 170 /172 (1994) 187. [23] M.C. Dang, J.J. Blandin, B. Baudelet, Acta Mater. 44 (1996) 3991. [24] Y. Takayama, T. Tozawa, H. Kato, Acta Mater. 47 (1999) 1263. [25] W.J. Kim, J. Wolfenstine, O.D. Sherby, Acta Metall. Mater. 39 (1991) 199. [26] M.C. Dang, B. Baudelet, J. Mater. Sci. 29 (1994) 2315. [27] K.T. Park, S.T. Yang, J.C. Earthman, F.A. Mohamed, Mater. Sci. Eng. A188 (1994) 59.