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Casting defects and the tensile properties of an AlSiMg alloy
MATERIALS SCIENCE & ENGINEERING ELSEVIER
Materials Science and Engineering A220 (1996) I09-116
A
Casting defects and the tensile properties of an A 1 - S i - M g alloy C.H. C/tceres a,*, B.I. Selling 1,b aCRC for Alloy and Solidification Technology (CAST), Department of Mining, Minerals and Materials Engineering, The University of Queensland, Brisbane QLD 4072, Australia bCSIRO Division of Manufactur#~g Technology, PO Box 883, Kenmore QLD 4069, Australia
Received 13 October 1995; revised t0 June 1996
Abstract
Samples containing either entrapped dross and oxide films, gas porosity or small drilled holes have been used to study the effect of different types of defects on the tensile behaviour of an A1-7Si-0.4Mg casting alloy. The tensile properties show little or no correlation with the bulk porosity content, especially in the case of samples containing dross and oxide films. In contrast, the decrease in tensile ductility and strength correlates with the area fraction of defects in the fracture surface of the samples. The experimental results are in agreement with the predictions of a simple analysis based on models for the growth of a plastic instability in a tensile sample. Keywords: Oxide films; Porosity; Tensile properties
1. Introduction
The presence of defects, in the form of gas or shrinkage porosity formed during solidification as well as dross or oxide films entrapped during the filling of the mould, can make the tensile behaviour of casting alloys unpredictable. Castings with thin sections are especially vulnerable to the effects of porosity since a single macropore (defined for the purposes of the following analysis as a pore or equivalent casting defect larger than the dendrite cell size, i.e., bigger than about 100 btm), may cover a significant fraction of the cross-sectional area. Even high integrity castings are expected to contain defects of one kind or another and thus it is important to be able to predict their effect on the mechanical performance of the material. The effect of porosity on the tensile properties of castings has been the matter of several studies [1-4]. Eady and Smith [1] considered the effect of large average volume fractions o f porosity, up to 7%, on the mechanical behaviour of A1-7%Si with Mg levels of 0.1-0.47%. They found that only the low yield stress varieties of the alloy (Mg content of less than 0.26%)
* Corresponding author. Present address: Volvo Aero Corp., Trollhattan, S-46181 Sweden. 0921-5093/96/$15.00 © 1996 -- Elsevier Science S.A. All rights reserved PII S0921-5093(96)10433-0
tolerate porosity levels in excess of 1%, in which case the tensile ductility decreased with the level of porosity. McLellan [4] reported that increasing soundness results in higher elongation to fracture in alloy A357. Herrera and Kondic [2] and Surappa et al. [3] carried out studies on A1-Si and A I - S i - M g alloys with low average levels of porosity ( < 0 . 4 % ) . Herrera and Kondic found a very large scatter when trying to correlate the tensile strength to the average volume fraction of porosity, whilst Surappa et al. showed that the elongation to fracture may actually increase with an increase in the average volume fraction of porosity, pointing out that this could be caused by a non uniform distribution of pores in the test bars. Instead, both studies showed that a better correlation and reduced scatter is obtained if the decrease in tensile strength is plotted against either the length [2] or the projected area [2,3] of the pores in the fracture surface. Similarly, Surappa et al. showed that the decrease in the elongation to fracture could be related to the projected area of the pores in the fracture surface. They also observed intense deformation and microcracking near the pores, but concluded that the concentration of stress and plastic strain is of the order of what can be expected from the reduction in cross sectional area. Herrera and Kondik's results prompted them to conclude that density is the least promising parameter for
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studying quantitative effects of cavities on the strength properties of a cast alloy. Similarly, Surappa et al. pointed out that the ductility and strength of the A I 7Si-0.3Mg-T6 alloy depend mainly on the size of the macropores on the fracture surface, rather than on volume percent porosity obtained by density measurement. More recently, a model [5] was presented which attempted to rationalise Surappa et al.'s results by assuming that defect-containing regions in a tensile sample reduce the load bearing area and, therefore, yield first, concentrating the strain. As observed by Surappa et al., the concentration of the strain results in the cracking of Si particles near the voids and thus the model assumed that this increase in the local damage caused premature fracture in the defect containing section. The predictions of this model showed good agreement with Surappa et al.'s results. In the present paper the previous model [5] is rediscussed and some of its predictions contrasted with new experiments aimed at quantifying the effect of different casting defects on the tensile behaviour of the Al-7%Si-0.4%Mg alloy in the T6 condition.
2. Micromechanics of fracture in the A1-7Si-0.4Mg casting alloy Plastic deformation results in the cracking of a significant fraction of eutectic Si particles [6-11] in this alloy. The cracking of Si particles occurs gradually with strain, constituting the dominant form of damage. As the applied strain increases, cracks in the particles become voids that grow and link forming larger cracks in the A1 matrix. These microcracks eventually become unstable, causing fracture. Recent detailed studies using the Bauschinger effect and quantitative metallography [10,11] have shown that the rate of cracking of Si particles is determined by the size of the dendritic cells and the size and shape of the Si particles. Tensile fracture has been observed to occur at a roughly constant level of damage in the form of cracked Si particles, about 15% of the total population [8,1t]. The differences in the fracture strain of sound material with different dendrite cell size are due to the different rates at which damage develops within each microstructure [9,11]. Fine microstructures, i.e. microstructures containing small dendrite cells and Si particles, develop damage at a low rate and thus require a large strain to reach the critical level of damage for fracture, whilst in coarser microstructures with large cell sizes, large and elongated Si particles tend to crack at low strains, lowering the ductility. Thus, in principle, it may be assumed that any given microstructure, free from porosity or other casting de-
fects, is characterised by a particular fracture strain determined by the scale of the dendritic structure and the size and shape of the Si particles [9,11].
3. The effect of geometric defects on the tensile ductility When porosity or an equivalent defect is present in a tensile sample, the load bearing area is reduced. It can thus be assumed that the defective region will yield first, concentrating the strain. The rate of strain concentration can be calculated considering the strain hardening ability of the material as described below. A geometric defect that locally reduces the load bearing area of a tensile sample results in the formation of an incipient neck. The growth of this sort of neck can be described using current models for the development of plastic instabilities [5,12-17]. If the neck is not sharp or the strains involved are not large, it may be assumed that only one significant stress exists in either the uniform section or the local inhomogeneity. Under these conditions, Jonas and Baudelet [14] pointed out that a geometrical defect (e.g. a machining defect) and an internal void in a tensile sample are equivalent and, therefore, their effect on the strain distribution may be analysed using the same formalism. The elongation to fracture in this alloy is fairly low in the T6 condition. Depending on the microstructure, values between 2 and 14% are normally observed [7,9,18] and fracture involves little or no necking [4]. Thus, strain gradients in the neck are unlikely to be very steep even at fracture and it seems reasonable to assume an uniaxial stress-strain state when the material is deformed.
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0.00 0,02 0.04 0,06 0.08 0.10 strain outside the defect region Fig. 2. The strain in the defect-containing region as a function of the strain outside the defect region, for area fractions of defects f in the fracture surface between 0 and 12%. Values computed using Eq. (3) for n = 0.1. The dashed line represents a fracture strain s* = 0.11.
On the other hand, voids in a cross section create a multiaxial stress state and cause strain concentrations near them. This perturbation to the assumed uniaxial stress state is clearly more important when the voids cover a significant fraction of the cross section. However, Surappa et aL pointed out that the concentration of stress and plastic strain near the voids is of the order of what can be expected from the reduction in cross sectional area. It will be thus assumed that the only effect of the voids is to reduce the cross section. Other possible effects on the neck's stress-strain state will be ignored. The approach o f Ghosh [16] is followed in the present analysis, with the exception that strain rate effects are neglected. The geometry for the model is depicted in Fig. 1. F o r simplicity, only a single spherical void is assumed in the gauge length of an otherwise perfect specimen. The geometry in Fig. 1 is probably realistic in the case of a macropore in a thin walled casting and is close to the physical situation o f Surappa et al.'s experiments, since their tensile samples had only one macropore in the gauge length. Other related situations will be considered at the end of this section. With reference to Fig. 1, the cross section in the smooth region, Ao, and the cross section in the region containing the defect, &, are such that Ai = Ao(1 --J'), where f i s the area fraction covered by the defect. Thus, load equilibrium along the tensile axis is maintained if [5,16] o-i(1 - J ) A o e - ~ -- o-hAo e -~'
(1)
where o-:, e4 and o'h, Sh are the true stresses and strains in and outside the defect region, respectively. A numerical solution of Eq. (1) requires the use of a constitutive equation. Assuming that the material follows the relationship ~r = K s n
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where o- and e are the true stress and true plastic strain, respectively, K is a constant and n is the strain-hardening exponent [19], substitution of Eq. (2) into (1) leads to (1 - f ) e - ~ ef = e - ~ e~
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which relates the strain inside the defect region, e~, to the strain outside, eh. The relation between the two strains is shown in Fig. 2 for a material with n = 0.1 and f between 0 and 12%. The line for f = 0 represents
Fig. 4. The dark, featureless region in the middIe of this SEM photomicrograph illustrates the size and shape of oxide films observed on the fracture surface of the samples.
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a sound sample for which the strains inside and outside the defect area are equal. As concluded in the previous section, it can be assumed that the material is characterised by a critical strain, e*, at which fracture occurs. With reference to the curves in Fig. 2, this assumption implies that for any given value of f, fracture occurs whenever the strain in the void region ~ reaches the critical strain e*. Thus, by recording the maximum homogeneous strain, eft, at which this happens, the maximum uniform strain (in Marciniak and Kuczinsky's terms, the limit strain [12,16]) can be predicted. If the contribution from necking to the total elongation is small, as in the present alloy, and can be neglected, the limit strain represents the tensile ductility of the material. The loss in tensile strength resulting from the reduced ductility is estimated as follows. The maximum in the load deflection curve of a sound sample occurs when [17,19] eh--=s=n in Eq. (2). The true tensile strength is then o-*= Kn n. If the elongation to fracture for a level of porosity f is e*, the true stress at fracture, o-~, is such that °'--2 or* = [ ~ 1 "
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which allows a straightforward calculation of the tensile strength from the elongation to fracture data and the strain-hardening exponent. The predictions of this model depend explicitly only on the values of n and f That is, for a given n-value, only the fractional cross section area covered by the defect is important. Whether the defect in Fig. 1 is spherical, irregularly shaped or even 'flat,' as in the case of dross or oxides films entrapped in the metal, should be unimportant. Likewise, whether the defect is single or multiple, or whether there are different types of defects in the cross section should be unimportant as well. Again, these two assertions imply that any stress field assod a t e d with the defects can be ignored. The experiments described in what follows were aimed at testing these hypotheses in an A I - 7 % S i 0.4%Mg casting alloy in the T6 condition. Three types of defects were introduced in the tensile samples: entrapped dross and oxide films to represent two-dimensional defects, gas/shrinkage porosity to represent near-spherical voids and drilled holes to simulate very targe casting defects. For each sample the bulk porosity and the area fraction of defects on the fracture surface were measured and the correlation between these parameters and the mechanical properties was assessed. The results were then contrasted with the predictions of the analytical model.
4. Experimental details A commercial A1-7%Si-0.4%Mg casting alloy (Australian denomination AA601, which is equivalent to the USA A356.0 alloy) was used for these experiments. Casting was done in a permanent mould with a cavity of dimensions 190 x 80 x 15 mm 3. The mould was preheated to about 400°C prior to casting. A few plates were cast without degassing, and, in order to ensure the entrapment of dross and oxide films, several others were cast after vigorously stirring the melt. The cast plates were then cut into small (15 x 15 x 60 m m 3) bars which were solution treated for 20 h at 540°C, quenched in water and subsequently aged for 6 h at 170°C. The average porosity content of the tests samples was determined by the density method. A small block of alloy was HIP-ed for 2 h at 540°C and an applied pressure of 100 MPa. Its density value, 6.677 g cm -3, was used as a reference for the density measurements. Tensile samples were machined from the bars, with a gauge length of 15 mm and rectangular cross section of 4 x 5 mm 2. Tensile testing was done in a screw driven machine, at a cross head speed of 0.5 mm s - 1, with an extensometer attached. Photographs were taken of both fracture surfaces after testing. The fractured surfaces were then observed under the stereomicroscope at a magnification of about x 50, and the defects delineated on the photographs. The defects were then transferred onto a single photograph and their area measured using an image analyser. The total measured area value was then divided by the initial cross section of the sample to find the defect area fraction. A number of samples obtained from the plates cast from the unstirred melt were examined with X-ray radiography. Twelve samples that showed no defects had a hole drilled through the thickness using drill bits of diameters 0.25, 0.35 and 0.50 ram. The fracture surfaces of these samples were subsequently checked under the stereomicroscope to ensure that no other significant defects were present.
5. Experimental results 5. I. Metallography o f fi'acture surfaces
Two main types of defects were identified on the fracture surfaces: spherical voids with shiny surfaces and clearly visible dendrites, identified as gas or shrinkage porosity (hereafter referred to as pores), and rough, dull or coloured regions, identified as dross or oxide films. Figs. 3 and 4 illustrate the shape and size of the observed defects. The samples obtained from the stirred melt showed a combination of oxides and pores in the
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Fig. 7. The true eIongation to fracture as a function o f the volumetric porosity content, for samples with different types o f defects.
fracture surface, although in a few cases dross and oxides accounted for most of the defect area, especially at large defect area fractions. Samples cast from the unstirred melt contained only gas type pores. The samples were thus sorted according to whether they had only gas pores or a combination of dross and pores. The number o f defects per frac-
ture surface varied between 3 and 30. Fracture commonly occurred on a flat plane perpendicular to the tensile axis, although the fracture surface showed some degree of tortuosity. A group of defects always fell into focus at the same time when observed on the microscope, indicating that they were intersected by a single cross sectional plane, while others lay above or below the apparent main fracture plane. However, trying to determine exactly whether a particular defect was intersected by the fracture plane or the locus of the actual intersection was a rather arbitrary exercise. Thus, it was assumed that all visible defects on the fracture surface were sitting on a single cross sectional plane and their projected area considered for the calculations.
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Fig. 5 compares the average porosity, fv, o f the test bars as determined by density measurement with the area fraction of defects on the fracture surface, f. F o r samples containing pores, (i.e., samples from the unstirred casting, circles) the area fraction of defects increases more or less proportionally to the average porosity content. I n contrast, there is no obvious correlation between the average density and the area fraction o f defects on the fracture surface of samples containing both pores and dross (i.e., samples from the stirred casting, squares). The local f value w a s up to some 20 times higher than the volumetric average, fv.
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C.H. Cdceres, B,L Selling / Materials Science and EngineerhTgA220 (1996) I09-It6
5.3. Tensile testing, smooth samples
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Elongations to fracture varied f r o m 2 to 11%. The samples f r o m the stirred melt were on average less ductile than those obtained f r o m the unstirred melt. Fig. 6 shows an experimental nominal stress-nominal strain curve, together with the true stress-true plastic strain curve, the latter superimposed on a curve drawn according to Eq. (2), with K = 459 M P a and n = 0.1. Twenty samples were selected for defect area fraction analysis. The true elongation to fracture o f these samples, calculated as e~ = ln(1 + s), where s is the engineering elongation to fracture, is plotted as a function of the volumetric porosity and the area fraction o f defects in the fracture surface in Figs. 7 and 8, respectively. The tensile strength, TS, of each sample was converted to true tensile strength as [19] G ~ = T S ( 1 + In(1 + s)), normalised to the tensile strength o f the m o s t ductile sample and plotted against the volumetric porosity and the area fraction o f defects in the fracture surface in Figs. 9 and (10), respectively. Figs. 7 and 9 show that although in general the tensile ductility and strength decrease with the average porosity content, the scatter in the data is very large. The correlation is especially p o o r in the case o f the samples containing film oxides and dross. In contrast, the data in Figs. 8 and 10 show a more clear trend o f decreasing ductility for increasing area fraction o f defects, with less scatter. It can also be seen that there is I
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Fig. 8. The true elongation to fracture as a function of the area fraction of defects in the fracture surface, for samples with different types of defects. The solid line represents the strain outside the defect region in Fig. 2 for a fracture strain ~* = 0.11 in the defect containing section.
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C,H. Cdceres, B.L Selling / Materials Science and Engineering A220 (i996) I 0 9 - l l d
5.4. Artificially damaged samples A total of 12 samples with artificial holes were tested. The elongation to fracture and tensile strength results for these samples are also plotted in Figs. 8 and 10. It can be seen that they blend very well with the results for the smooth samples.
5.5. Comparison with calculated results The elongation to fracture of the most ductile sample in these experiments was 11% (see results for smooth samples). This value was used in Fig. 2 (s* = 0.11, indicated by the dashed line) to find the maximum homogeneous strain, s*, for the different f-values. The e* values (joined by the solid line) are compared in Fig. 8 with the experimental tensile ductility as a function of the defect density on the fracture surface. It is seen that the calculated curve correctly predicts the trend of the experimental points. The effect on the tensile strength was calculated using Eq. (4), with n = 0 . 1 and the st values from Fig. 8 (solid line). The computed values are compared in Fig. 10 (joined by the solid line) with the experimental tensile strength as a function of the defect density on the fracture surface. Again, it can be seen that the calculated curve indicates the trend o f the experimental results.
6. Discussion
The results shown in Figs. 7 and 9 indicate, in agreement with Herrera and Kondic and Surappa et al., that the average porosity content is not a reliable parameter to predict the tensile behaviour of the material. This is especially true in the case of material contaminated with dross or oxide films. In contrast, and also in agreement with Surappa et al., the area fraction of defects in the fracture surface results in a much better correlation, with a clear trend despite some scatter in the data (Figs. 8 and 10). The results in Figs. 8 and l0 are also in good agreement with the analytical model. A general conclusion can thus be made that the dominant parameter is the area fraction of defects in the fracture surface, f. A few experimental and analytical aspects of this correlation are discussed next. The scatter observed in the experimental points in Figs. 8 and 10 can be considered, at least in part, to be inherent to the fracture behaviour o f the material. Even samples from degassed, sound castings, tend to show some scatter in the tensile ductility [7,9,18] and at least variations of the same order were expected to affect the present measurements of ductility. In the case of the tensile strength, part of the vertical scatter can be ascribed to the fact that the values were
115
normalised to a single reference value, while in fact variations of up to + 10 MPa were observed in the tensile strength of samples with comparable ductility, which represents a variation in excess of _+ 3% in the normalised values of Fig. 10. Another possible source of scatter, this time along the x-axis, stems from the measurements of the defect sizes since the void boundaries on the fracture surface were often ill-defined, making some of the measurements somewhat arbitrary. A related and probably more significant effect stems from the fact that not all defects lay on the same cross sectional plane and therefore the projected area used to calculate f may be an overestimate of the actual defect area in the cross section. It should be noted that the defect area fraction was calculated by referring the measured area to the initial cross section of the sample. This procedure ignores any possible strain effects on the size of the defects and thus it may also tend to overestimate the actual f value. However, Bourcier et at. [20] observed that even after strains considerably larger than those involved in these experiments the size of the pores found in the fracture surface of porous alloys was very similar to those in the undeformed alloy. This suggests that the error introduced by the procedure is likely to be very small. The results in Fig. 5 show that the local area fraction of defects may be some 20 times higher than the average bulk porosity content. Bourcier at al. [20] also observed that the defect content in the fracture surface was much higher (some 10 times) than the bulk porosity in their experiments, and concluded that the non regularity of defect distribution along the gauge length determines where fracture occurs. They pointed out that paths or regions of high porosity in an otherwise homogeneous continuum can be viewed as inhomogeneities which, in terms of Marciniak and Kuczinski's analysis [12], can always be reduced to geometrical imperfections and thus expressed by the parameter f, as done in the present analysis. Surappa et al. [3] suggested that the envelope of the defect, defined as a circle with diameter equal to the largest dimension of the defect, correlates better with the ductility than the projected area, suggesting that the shape of the void can be important. The degree of clustering for a given defect area density may also influence the results as shown by experiments carried out with small drilled holes in wrought aluminium alloys [21]. Thus, there is the possibility that some of the scatter in Figs. 8 and 10 can also be ascribed to these causes. However, Figs. 8 and 10 show no systematic differences between Surappa et al.'s results obtained with samples with only one macropore in the gauge section and the present results involving either smooth samples with multiple defects and both flat (oxides) and three dimensional defects (pores) or sam-
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pies with artificial (single) defects. This, once again, supports the conclusion that the dominant parameter is the area fraction of defects in the cross section, regardless of the defect shape, number or distribution, in accord with the predictions of the analytical model. It must be kept in mind that the total elongation to fracture is actually the limit strain plus any contribution from the post-uniform deformation to the extension of the sample. In that sense, the limit strain (the solid line in Fig. 8) represents only a lower bound for the tensile ductility. The tensile strength of the material, on the other hand, is a sole function of the limit strain and therefore the calculated curve in Fig. 10 is a more rigorous prediction of the effect of the defects. It is apparent that although the present results show that the detrimental effect of casting defects on the mechanical properties can be quantified in a fairly simple way, the procedure requires the knowledge of the maximum defect density f in a given section of the casting. In the present experiments f was obtained by a post-mortem examination of the tensile samples, which is hardly a practical procedure for real castings. The practical value of this analysis would be greatly enhanced if the correlation of Figs. 8 and 10 could be demonstrated for defects in the gauge length detected prior to testing, e.g., by standard N D T methods (ultrasound or X-ray radiography).
7. Conclusions A series of experiments has been carried out to quantify the effect of casting defects on the tensile properties of an A 1 - 7 % S i - 0 . 4 % M g - T 6 casting alloy. Samples containing different types of defects, namely, gas porosity, a combination of gas pores and entrapped dross and oxides and artificial holes, have been studied. Both the ductility and the tensile strength show little or no correlation with the bulk porosity content, especially in the case of samples containing dross and oxide films. In contrast, the mechanical performance decreases monotonically with an increase in the area fraction of defects in the fracture surface of the samples. The projected area fraction of defects in a given cross
section of the sample seems to determine the effect of casting defects on the tensile behaviour. Whether the defect is single or multiple, or whether it is cylindrical, near-spherical (gas porosity), or two-dimensional (dross or oxide film), does not seem to be significant. The results are in very good agreement with predictions based on models for the growth of a plastic instability in a tensile sample.
Acknowledgements The authors are indebted to Dr J.R. Griffiths for valuable suggestions during the course of the work.
References [1] J.A. Eady and D.M. Smith, Mater. Forum, 9 (1986) 217. [2] A. Herrera and V. Kondic, Proc. Int. Conf. Solidification and Cast Metals, The Metals Society, Siwffield, UK, 1977, p. 460. [3] M.K. Snrappa, E.W. Blank and J.C. Jaquet, Seripta Metatl., 20 (1986) 1281. [4] D,L. McLellan, J. Test. Eval., 8 (1980) 170, [5] C.H. C~ceres, Scripta Metall. Mater., 32 (1995) 1851. [6] A. Gangulee and J. Gurland, Trans. Metatl. Soc. AIME, 239 (1967) 269. [7] S.F. Frederick and W.A. Bailey, Trans. Metall. Soc. AIME, 242 (1968) 2063. [8] M.K. Surappa, E.W. Blank and J.C. Jaquet, 3rd Int. Conf. Solidification Processing, Sheffield, UK, 1987, pp. 424-427. [9] C.H. C~iceres, C.J. Davidson and J.R. Griffiths, Mater. Sci. Eng., A197 (1995) 171. [10] C.H. Cfieeres, J.R. Griffiths and P. Reiner, Acta Mater., 44 (1996) 15. [1I] C.H. Chceres and J.R. Griffiths,Acta Mater., 44 (1996) 25. [12] Z.Marciniak and K. Kuczinsky,Int. J. Mech. Sci., 9 (1967) 609. [13] E.W. Hart, Acta Metatt., 15 (1967) 351. [14] J.J. Jonas and B. Baudelet, Acta Metatl., 25 (1977) 43. [15] E. Duncombe, Int. J. Mech. Sci., I4 (1972) 325. [16] A.K. Ghosh, Acta Metall., 25 (1977) 1413. [17] F.A. Nichols, Acta Metalt., 28 (1980) 663. [18] R.E. Spear and G.R. Gardner, Trans. Am. Foundrymen's Soc., 7I (1963) 209. [19] G.E. Dieter, Mechanical Metalh~rgy, 3rd edn., McGraw-Hill, New York, 1986, p. 287. [20] R.J. Bourcier,D.A. Koss, R.E. Smelser and O. Richmond, Acta Metatt., 34 (1986) 2443. [21] E.M. Dubenskyand D.A. Koss, Metall. Trans., I8A (1987) 1887.
Materials Science and Engineering A220 (1996) I09-116
A
Casting defects and the tensile properties of an A 1 - S i - M g alloy C.H. C/tceres a,*, B.I. Selling 1,b aCRC for Alloy and Solidification Technology (CAST), Department of Mining, Minerals and Materials Engineering, The University of Queensland, Brisbane QLD 4072, Australia bCSIRO Division of Manufactur#~g Technology, PO Box 883, Kenmore QLD 4069, Australia
Received 13 October 1995; revised t0 June 1996
Abstract
Samples containing either entrapped dross and oxide films, gas porosity or small drilled holes have been used to study the effect of different types of defects on the tensile behaviour of an A1-7Si-0.4Mg casting alloy. The tensile properties show little or no correlation with the bulk porosity content, especially in the case of samples containing dross and oxide films. In contrast, the decrease in tensile ductility and strength correlates with the area fraction of defects in the fracture surface of the samples. The experimental results are in agreement with the predictions of a simple analysis based on models for the growth of a plastic instability in a tensile sample. Keywords: Oxide films; Porosity; Tensile properties
1. Introduction
The presence of defects, in the form of gas or shrinkage porosity formed during solidification as well as dross or oxide films entrapped during the filling of the mould, can make the tensile behaviour of casting alloys unpredictable. Castings with thin sections are especially vulnerable to the effects of porosity since a single macropore (defined for the purposes of the following analysis as a pore or equivalent casting defect larger than the dendrite cell size, i.e., bigger than about 100 btm), may cover a significant fraction of the cross-sectional area. Even high integrity castings are expected to contain defects of one kind or another and thus it is important to be able to predict their effect on the mechanical performance of the material. The effect of porosity on the tensile properties of castings has been the matter of several studies [1-4]. Eady and Smith [1] considered the effect of large average volume fractions o f porosity, up to 7%, on the mechanical behaviour of A1-7%Si with Mg levels of 0.1-0.47%. They found that only the low yield stress varieties of the alloy (Mg content of less than 0.26%)
* Corresponding author. Present address: Volvo Aero Corp., Trollhattan, S-46181 Sweden. 0921-5093/96/$15.00 © 1996 -- Elsevier Science S.A. All rights reserved PII S0921-5093(96)10433-0
tolerate porosity levels in excess of 1%, in which case the tensile ductility decreased with the level of porosity. McLellan [4] reported that increasing soundness results in higher elongation to fracture in alloy A357. Herrera and Kondic [2] and Surappa et al. [3] carried out studies on A1-Si and A I - S i - M g alloys with low average levels of porosity ( < 0 . 4 % ) . Herrera and Kondic found a very large scatter when trying to correlate the tensile strength to the average volume fraction of porosity, whilst Surappa et al. showed that the elongation to fracture may actually increase with an increase in the average volume fraction of porosity, pointing out that this could be caused by a non uniform distribution of pores in the test bars. Instead, both studies showed that a better correlation and reduced scatter is obtained if the decrease in tensile strength is plotted against either the length [2] or the projected area [2,3] of the pores in the fracture surface. Similarly, Surappa et al. showed that the decrease in the elongation to fracture could be related to the projected area of the pores in the fracture surface. They also observed intense deformation and microcracking near the pores, but concluded that the concentration of stress and plastic strain is of the order of what can be expected from the reduction in cross sectional area. Herrera and Kondik's results prompted them to conclude that density is the least promising parameter for
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C.H. Cdceres, B.L Selling / Materials Science and Engineering A220 (i996) I09-116
studying quantitative effects of cavities on the strength properties of a cast alloy. Similarly, Surappa et al. pointed out that the ductility and strength of the A I 7Si-0.3Mg-T6 alloy depend mainly on the size of the macropores on the fracture surface, rather than on volume percent porosity obtained by density measurement. More recently, a model [5] was presented which attempted to rationalise Surappa et al.'s results by assuming that defect-containing regions in a tensile sample reduce the load bearing area and, therefore, yield first, concentrating the strain. As observed by Surappa et al., the concentration of the strain results in the cracking of Si particles near the voids and thus the model assumed that this increase in the local damage caused premature fracture in the defect containing section. The predictions of this model showed good agreement with Surappa et al.'s results. In the present paper the previous model [5] is rediscussed and some of its predictions contrasted with new experiments aimed at quantifying the effect of different casting defects on the tensile behaviour of the Al-7%Si-0.4%Mg alloy in the T6 condition.
2. Micromechanics of fracture in the A1-7Si-0.4Mg casting alloy Plastic deformation results in the cracking of a significant fraction of eutectic Si particles [6-11] in this alloy. The cracking of Si particles occurs gradually with strain, constituting the dominant form of damage. As the applied strain increases, cracks in the particles become voids that grow and link forming larger cracks in the A1 matrix. These microcracks eventually become unstable, causing fracture. Recent detailed studies using the Bauschinger effect and quantitative metallography [10,11] have shown that the rate of cracking of Si particles is determined by the size of the dendritic cells and the size and shape of the Si particles. Tensile fracture has been observed to occur at a roughly constant level of damage in the form of cracked Si particles, about 15% of the total population [8,1t]. The differences in the fracture strain of sound material with different dendrite cell size are due to the different rates at which damage develops within each microstructure [9,11]. Fine microstructures, i.e. microstructures containing small dendrite cells and Si particles, develop damage at a low rate and thus require a large strain to reach the critical level of damage for fracture, whilst in coarser microstructures with large cell sizes, large and elongated Si particles tend to crack at low strains, lowering the ductility. Thus, in principle, it may be assumed that any given microstructure, free from porosity or other casting de-
fects, is characterised by a particular fracture strain determined by the scale of the dendritic structure and the size and shape of the Si particles [9,11].
3. The effect of geometric defects on the tensile ductility When porosity or an equivalent defect is present in a tensile sample, the load bearing area is reduced. It can thus be assumed that the defective region will yield first, concentrating the strain. The rate of strain concentration can be calculated considering the strain hardening ability of the material as described below. A geometric defect that locally reduces the load bearing area of a tensile sample results in the formation of an incipient neck. The growth of this sort of neck can be described using current models for the development of plastic instabilities [5,12-17]. If the neck is not sharp or the strains involved are not large, it may be assumed that only one significant stress exists in either the uniform section or the local inhomogeneity. Under these conditions, Jonas and Baudelet [14] pointed out that a geometrical defect (e.g. a machining defect) and an internal void in a tensile sample are equivalent and, therefore, their effect on the strain distribution may be analysed using the same formalism. The elongation to fracture in this alloy is fairly low in the T6 condition. Depending on the microstructure, values between 2 and 14% are normally observed [7,9,18] and fracture involves little or no necking [4]. Thus, strain gradients in the neck are unlikely to be very steep even at fracture and it seems reasonable to assume an uniaxial stress-strain state when the material is deformed.
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0.00 0,02 0.04 0,06 0.08 0.10 strain outside the defect region Fig. 2. The strain in the defect-containing region as a function of the strain outside the defect region, for area fractions of defects f in the fracture surface between 0 and 12%. Values computed using Eq. (3) for n = 0.1. The dashed line represents a fracture strain s* = 0.11.
On the other hand, voids in a cross section create a multiaxial stress state and cause strain concentrations near them. This perturbation to the assumed uniaxial stress state is clearly more important when the voids cover a significant fraction of the cross section. However, Surappa et aL pointed out that the concentration of stress and plastic strain near the voids is of the order of what can be expected from the reduction in cross sectional area. It will be thus assumed that the only effect of the voids is to reduce the cross section. Other possible effects on the neck's stress-strain state will be ignored. The approach o f Ghosh [16] is followed in the present analysis, with the exception that strain rate effects are neglected. The geometry for the model is depicted in Fig. 1. F o r simplicity, only a single spherical void is assumed in the gauge length of an otherwise perfect specimen. The geometry in Fig. 1 is probably realistic in the case of a macropore in a thin walled casting and is close to the physical situation o f Surappa et al.'s experiments, since their tensile samples had only one macropore in the gauge length. Other related situations will be considered at the end of this section. With reference to Fig. 1, the cross section in the smooth region, Ao, and the cross section in the region containing the defect, &, are such that Ai = Ao(1 --J'), where f i s the area fraction covered by the defect. Thus, load equilibrium along the tensile axis is maintained if [5,16] o-i(1 - J ) A o e - ~ -- o-hAo e -~'
(1)
where o-:, e4 and o'h, Sh are the true stresses and strains in and outside the defect region, respectively. A numerical solution of Eq. (1) requires the use of a constitutive equation. Assuming that the material follows the relationship ~r = K s n
(2)
where o- and e are the true stress and true plastic strain, respectively, K is a constant and n is the strain-hardening exponent [19], substitution of Eq. (2) into (1) leads to (1 - f ) e - ~ ef = e - ~ e~
(3)
which relates the strain inside the defect region, e~, to the strain outside, eh. The relation between the two strains is shown in Fig. 2 for a material with n = 0.1 and f between 0 and 12%. The line for f = 0 represents
Fig. 4. The dark, featureless region in the middIe of this SEM photomicrograph illustrates the size and shape of oxide films observed on the fracture surface of the samples.
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a sound sample for which the strains inside and outside the defect area are equal. As concluded in the previous section, it can be assumed that the material is characterised by a critical strain, e*, at which fracture occurs. With reference to the curves in Fig. 2, this assumption implies that for any given value of f, fracture occurs whenever the strain in the void region ~ reaches the critical strain e*. Thus, by recording the maximum homogeneous strain, eft, at which this happens, the maximum uniform strain (in Marciniak and Kuczinsky's terms, the limit strain [12,16]) can be predicted. If the contribution from necking to the total elongation is small, as in the present alloy, and can be neglected, the limit strain represents the tensile ductility of the material. The loss in tensile strength resulting from the reduced ductility is estimated as follows. The maximum in the load deflection curve of a sound sample occurs when [17,19] eh--=s=n in Eq. (2). The true tensile strength is then o-*= Kn n. If the elongation to fracture for a level of porosity f is e*, the true stress at fracture, o-~, is such that °'--2 or* = [ ~ 1 "
(4)
which allows a straightforward calculation of the tensile strength from the elongation to fracture data and the strain-hardening exponent. The predictions of this model depend explicitly only on the values of n and f That is, for a given n-value, only the fractional cross section area covered by the defect is important. Whether the defect in Fig. 1 is spherical, irregularly shaped or even 'flat,' as in the case of dross or oxides films entrapped in the metal, should be unimportant. Likewise, whether the defect is single or multiple, or whether there are different types of defects in the cross section should be unimportant as well. Again, these two assertions imply that any stress field assod a t e d with the defects can be ignored. The experiments described in what follows were aimed at testing these hypotheses in an A I - 7 % S i 0.4%Mg casting alloy in the T6 condition. Three types of defects were introduced in the tensile samples: entrapped dross and oxide films to represent two-dimensional defects, gas/shrinkage porosity to represent near-spherical voids and drilled holes to simulate very targe casting defects. For each sample the bulk porosity and the area fraction of defects on the fracture surface were measured and the correlation between these parameters and the mechanical properties was assessed. The results were then contrasted with the predictions of the analytical model.
4. Experimental details A commercial A1-7%Si-0.4%Mg casting alloy (Australian denomination AA601, which is equivalent to the USA A356.0 alloy) was used for these experiments. Casting was done in a permanent mould with a cavity of dimensions 190 x 80 x 15 mm 3. The mould was preheated to about 400°C prior to casting. A few plates were cast without degassing, and, in order to ensure the entrapment of dross and oxide films, several others were cast after vigorously stirring the melt. The cast plates were then cut into small (15 x 15 x 60 m m 3) bars which were solution treated for 20 h at 540°C, quenched in water and subsequently aged for 6 h at 170°C. The average porosity content of the tests samples was determined by the density method. A small block of alloy was HIP-ed for 2 h at 540°C and an applied pressure of 100 MPa. Its density value, 6.677 g cm -3, was used as a reference for the density measurements. Tensile samples were machined from the bars, with a gauge length of 15 mm and rectangular cross section of 4 x 5 mm 2. Tensile testing was done in a screw driven machine, at a cross head speed of 0.5 mm s - 1, with an extensometer attached. Photographs were taken of both fracture surfaces after testing. The fractured surfaces were then observed under the stereomicroscope at a magnification of about x 50, and the defects delineated on the photographs. The defects were then transferred onto a single photograph and their area measured using an image analyser. The total measured area value was then divided by the initial cross section of the sample to find the defect area fraction. A number of samples obtained from the plates cast from the unstirred melt were examined with X-ray radiography. Twelve samples that showed no defects had a hole drilled through the thickness using drill bits of diameters 0.25, 0.35 and 0.50 ram. The fracture surfaces of these samples were subsequently checked under the stereomicroscope to ensure that no other significant defects were present.
5. Experimental results 5. I. Metallography o f fi'acture surfaces
Two main types of defects were identified on the fracture surfaces: spherical voids with shiny surfaces and clearly visible dendrites, identified as gas or shrinkage porosity (hereafter referred to as pores), and rough, dull or coloured regions, identified as dross or oxide films. Figs. 3 and 4 illustrate the shape and size of the observed defects. The samples obtained from the stirred melt showed a combination of oxides and pores in the
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Fig. 7. The true eIongation to fracture as a function o f the volumetric porosity content, for samples with different types o f defects.
fracture surface, although in a few cases dross and oxides accounted for most of the defect area, especially at large defect area fractions. Samples cast from the unstirred melt contained only gas type pores. The samples were thus sorted according to whether they had only gas pores or a combination of dross and pores. The number o f defects per frac-
ture surface varied between 3 and 30. Fracture commonly occurred on a flat plane perpendicular to the tensile axis, although the fracture surface showed some degree of tortuosity. A group of defects always fell into focus at the same time when observed on the microscope, indicating that they were intersected by a single cross sectional plane, while others lay above or below the apparent main fracture plane. However, trying to determine exactly whether a particular defect was intersected by the fracture plane or the locus of the actual intersection was a rather arbitrary exercise. Thus, it was assumed that all visible defects on the fracture surface were sitting on a single cross sectional plane and their projected area considered for the calculations.
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Fig. 5 compares the average porosity, fv, o f the test bars as determined by density measurement with the area fraction of defects on the fracture surface, f. F o r samples containing pores, (i.e., samples from the unstirred casting, circles) the area fraction of defects increases more or less proportionally to the average porosity content. I n contrast, there is no obvious correlation between the average density and the area fraction o f defects on the fracture surface of samples containing both pores and dross (i.e., samples from the stirred casting, squares). The local f value w a s up to some 20 times higher than the volumetric average, fv.
114
C.H. Cdceres, B,L Selling / Materials Science and EngineerhTgA220 (1996) I09-It6
5.3. Tensile testing, smooth samples
1.1
Elongations to fracture varied f r o m 2 to 11%. The samples f r o m the stirred melt were on average less ductile than those obtained f r o m the unstirred melt. Fig. 6 shows an experimental nominal stress-nominal strain curve, together with the true stress-true plastic strain curve, the latter superimposed on a curve drawn according to Eq. (2), with K = 459 M P a and n = 0.1. Twenty samples were selected for defect area fraction analysis. The true elongation to fracture o f these samples, calculated as e~ = ln(1 + s), where s is the engineering elongation to fracture, is plotted as a function of the volumetric porosity and the area fraction o f defects in the fracture surface in Figs. 7 and 8, respectively. The tensile strength, TS, of each sample was converted to true tensile strength as [19] G ~ = T S ( 1 + In(1 + s)), normalised to the tensile strength o f the m o s t ductile sample and plotted against the volumetric porosity and the area fraction o f defects in the fracture surface in Figs. 9 and (10), respectively. Figs. 7 and 9 show that although in general the tensile ductility and strength decrease with the average porosity content, the scatter in the data is very large. The correlation is especially p o o r in the case o f the samples containing film oxides and dross. In contrast, the data in Figs. 8 and 10 show a more clear trend o f decreasing ductility for increasing area fraction o f defects, with less scatter. It can also be seen that there is I
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Fig. 8. The true elongation to fracture as a function of the area fraction of defects in the fracture surface, for samples with different types of defects. The solid line represents the strain outside the defect region in Fig. 2 for a fracture strain ~* = 0.11 in the defect containing section.
pores pores/dross drilled holes Surappa et al.[3] calculated
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Fig. 10. The true tensile strength normalised to the true tensile strength of the most ductile sample, as a function of the area fraction of defects in the fracture surface, for samples with different types of defects. The solid line represents values computed using Eq. (4).
C,H. Cdceres, B.L Selling / Materials Science and Engineering A220 (i996) I 0 9 - l l d
5.4. Artificially damaged samples A total of 12 samples with artificial holes were tested. The elongation to fracture and tensile strength results for these samples are also plotted in Figs. 8 and 10. It can be seen that they blend very well with the results for the smooth samples.
5.5. Comparison with calculated results The elongation to fracture of the most ductile sample in these experiments was 11% (see results for smooth samples). This value was used in Fig. 2 (s* = 0.11, indicated by the dashed line) to find the maximum homogeneous strain, s*, for the different f-values. The e* values (joined by the solid line) are compared in Fig. 8 with the experimental tensile ductility as a function of the defect density on the fracture surface. It is seen that the calculated curve correctly predicts the trend of the experimental points. The effect on the tensile strength was calculated using Eq. (4), with n = 0 . 1 and the st values from Fig. 8 (solid line). The computed values are compared in Fig. 10 (joined by the solid line) with the experimental tensile strength as a function of the defect density on the fracture surface. Again, it can be seen that the calculated curve indicates the trend o f the experimental results.
6. Discussion
The results shown in Figs. 7 and 9 indicate, in agreement with Herrera and Kondic and Surappa et al., that the average porosity content is not a reliable parameter to predict the tensile behaviour of the material. This is especially true in the case of material contaminated with dross or oxide films. In contrast, and also in agreement with Surappa et al., the area fraction of defects in the fracture surface results in a much better correlation, with a clear trend despite some scatter in the data (Figs. 8 and 10). The results in Figs. 8 and l0 are also in good agreement with the analytical model. A general conclusion can thus be made that the dominant parameter is the area fraction of defects in the fracture surface, f. A few experimental and analytical aspects of this correlation are discussed next. The scatter observed in the experimental points in Figs. 8 and 10 can be considered, at least in part, to be inherent to the fracture behaviour o f the material. Even samples from degassed, sound castings, tend to show some scatter in the tensile ductility [7,9,18] and at least variations of the same order were expected to affect the present measurements of ductility. In the case of the tensile strength, part of the vertical scatter can be ascribed to the fact that the values were
115
normalised to a single reference value, while in fact variations of up to + 10 MPa were observed in the tensile strength of samples with comparable ductility, which represents a variation in excess of _+ 3% in the normalised values of Fig. 10. Another possible source of scatter, this time along the x-axis, stems from the measurements of the defect sizes since the void boundaries on the fracture surface were often ill-defined, making some of the measurements somewhat arbitrary. A related and probably more significant effect stems from the fact that not all defects lay on the same cross sectional plane and therefore the projected area used to calculate f may be an overestimate of the actual defect area in the cross section. It should be noted that the defect area fraction was calculated by referring the measured area to the initial cross section of the sample. This procedure ignores any possible strain effects on the size of the defects and thus it may also tend to overestimate the actual f value. However, Bourcier et at. [20] observed that even after strains considerably larger than those involved in these experiments the size of the pores found in the fracture surface of porous alloys was very similar to those in the undeformed alloy. This suggests that the error introduced by the procedure is likely to be very small. The results in Fig. 5 show that the local area fraction of defects may be some 20 times higher than the average bulk porosity content. Bourcier at al. [20] also observed that the defect content in the fracture surface was much higher (some 10 times) than the bulk porosity in their experiments, and concluded that the non regularity of defect distribution along the gauge length determines where fracture occurs. They pointed out that paths or regions of high porosity in an otherwise homogeneous continuum can be viewed as inhomogeneities which, in terms of Marciniak and Kuczinski's analysis [12], can always be reduced to geometrical imperfections and thus expressed by the parameter f, as done in the present analysis. Surappa et al. [3] suggested that the envelope of the defect, defined as a circle with diameter equal to the largest dimension of the defect, correlates better with the ductility than the projected area, suggesting that the shape of the void can be important. The degree of clustering for a given defect area density may also influence the results as shown by experiments carried out with small drilled holes in wrought aluminium alloys [21]. Thus, there is the possibility that some of the scatter in Figs. 8 and 10 can also be ascribed to these causes. However, Figs. 8 and 10 show no systematic differences between Surappa et al.'s results obtained with samples with only one macropore in the gauge section and the present results involving either smooth samples with multiple defects and both flat (oxides) and three dimensional defects (pores) or sam-
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C.H. Cdceres, B.L Selt#~g / Materials Science and Engineering A220 (1996) 109-116
pies with artificial (single) defects. This, once again, supports the conclusion that the dominant parameter is the area fraction of defects in the cross section, regardless of the defect shape, number or distribution, in accord with the predictions of the analytical model. It must be kept in mind that the total elongation to fracture is actually the limit strain plus any contribution from the post-uniform deformation to the extension of the sample. In that sense, the limit strain (the solid line in Fig. 8) represents only a lower bound for the tensile ductility. The tensile strength of the material, on the other hand, is a sole function of the limit strain and therefore the calculated curve in Fig. 10 is a more rigorous prediction of the effect of the defects. It is apparent that although the present results show that the detrimental effect of casting defects on the mechanical properties can be quantified in a fairly simple way, the procedure requires the knowledge of the maximum defect density f in a given section of the casting. In the present experiments f was obtained by a post-mortem examination of the tensile samples, which is hardly a practical procedure for real castings. The practical value of this analysis would be greatly enhanced if the correlation of Figs. 8 and 10 could be demonstrated for defects in the gauge length detected prior to testing, e.g., by standard N D T methods (ultrasound or X-ray radiography).
7. Conclusions A series of experiments has been carried out to quantify the effect of casting defects on the tensile properties of an A 1 - 7 % S i - 0 . 4 % M g - T 6 casting alloy. Samples containing different types of defects, namely, gas porosity, a combination of gas pores and entrapped dross and oxides and artificial holes, have been studied. Both the ductility and the tensile strength show little or no correlation with the bulk porosity content, especially in the case of samples containing dross and oxide films. In contrast, the mechanical performance decreases monotonically with an increase in the area fraction of defects in the fracture surface of the samples. The projected area fraction of defects in a given cross
section of the sample seems to determine the effect of casting defects on the tensile behaviour. Whether the defect is single or multiple, or whether it is cylindrical, near-spherical (gas porosity), or two-dimensional (dross or oxide film), does not seem to be significant. The results are in very good agreement with predictions based on models for the growth of a plastic instability in a tensile sample.
Acknowledgements The authors are indebted to Dr J.R. Griffiths for valuable suggestions during the course of the work.
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