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Cast Al-7 Si composites: Effect of particle type and size on mechanical properties
Scripta METALLURGICA
Vol. 2 2 , pp. 1307-1312, 1988 Printed in the U.S.A.
Pergamon Press plc
CAST Al - 7 Si COMPOSITES: EFFECT OF PARTIt"I.E TYPE AND SIZE ON MECHANICAL PROPERTIES J.J. Stephens*, J.P. Lucas+ and F.M. Hosking++ *Lockheed Missiles and Space Co., Inc. Pale Alto, CA 94304 +Sandia National Laboratories Livermore, CA 94550 ++Sandia National Laboratories Albuquerque, NM 87185
(Received April 2 7 , 1988) (Revised June 7, 1988)
1. Inm~ducfion Al-minum-based metal matrix composites (MMCs) offer significant increases in modulus and strength over their unreinforeed counterparts, and in recent years have been the focus of intensive efforts aimed at understanding the relation between processing and properties of these materials. Discontinuous reinforcement aluminum MMC's are especially auxactive for applications where low thermal expansion coefficient along with low density are desired. While the majority of discontinuous a]umimun M]VIC'sconsidered for aerospace applications have been processed with powder metallurgical methods, cast aluminum MMCs offer significant cost advantages over their powder processed counterparts [1]. With the recent commercial availability of cast al,~mlnum alloy composites, micrnsuuctoral and mechanical characterization has become an important issue in composite development. Issues such as the araount of ductility possible for reinforced alloys in the solution treated and peak aged condition become quite important when one is considering using matrix alloys which tend to have low ductility in the heat treated condition. The present study was undertaken in order to characterize the effect of particulate type and size on room temperature mechanical prope~ies for cast Al-7 Si (alloys A356 and A357) MMCs. In particular, we have examined whether a specific reinforcement type (B4C or SIC), and size can provide a useful combination of tensile, fracture toughness and impact properties. 2. Exverimental The. .present investigation focuses exclusively on cast ~uminum MMCs with matrices of either A356 (nominal composluon, wt%: Al - 7 St - 0.35 Mg) or A357 (AI - 7 Sl - 0.55 Mg). For purposes of comparison with unreinfomed a
mold
of,,356
was also
ind L The
were
I
al Aluminom
pomtes Corporation, La Jolla, Ca, m the form of 19.5 x 3.4 x 1.25" (L x T x ST directions) bars. The various combinations of m,trix, reinforcement and particle size studied are shown in Table I. All of the composites had reinfmcement volume fractions between 25 and 30 vol%. Exact details are considered proprietary by the manufacturer, but basic process~g steps include surface preparation of the reinfon:ements and casting of the composites using a process such as stir..casting [2,3] which permits the r e i n f ~ c n t to remai, suspended in the molten metal. In all cases, the cast MMC's were hot isostafically pressed (HIPped) in order to reduce casting porosity. Sheet specimens cut from the bar with a longitodin,l orientation were machined for tensile testing with a gage length of 2" and a cross section of 0.5 x 0.5" [4]. Specimens were machined for impact testing with a cross section of 0.395 x 0.395", 2.165" long and a 45 ° notch with a depth of 0.080" [5], in both the L-T and T-L orientations [6]. Short bar fracture toughness specimens were machined with a width of 1", length of 1.5" and 0.87" thickness, in both the T-L and ST-L orientations. Test specimens for materials #1-4 were solution Ireated and aged in salt pots using standard heat treating procedures for the unreinfor~d A356 alloy (see Table I); materials #5-7 were heat treated by the manufacturer. Since the presence of a reinforeement can lead to er,celeratiun of aging kinetics [7-9], test coupons of materials #2-4 were aged for various times at 163°C in order to establish the time to peak age for these composites. Results for these MMC's indicated that the times for peak hardness fell within the range of tunes requited (4-5 hr) at 163°C for peak aging of the un~inforeedA356 alloy. This lack of aecelenUion in aging kinetics is apparently due to the relatively large size of the reinforcement. Nevertheless, other
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investigators have observed acceleration in aging kinetics - about a factor of 2 in time to peak hardness - in 6061 Al alloy with fine (-5 ixm size) reinforcements of either B4C or SiC [7,8]. Separate work in our laboratory on 6061/SIC and 7091/B4C composites with 30 vol%, 5 ~ size particulate has also shown acceleration in aging kinetics [9] relative to the unreinforced alloys. Tensile specimens were tested to fracture at room temperature using a MTS 810 servohydranlic testing machine equipped with a 50 kip capacity actuator and 30 kip capacity wedge type gips. A cfip-on extensometer with a gage length of 1" was used to measure strain. All tests were ran at an engineming strain rate of 10-3 s-l. The impact tests were performed using a ManI abs Model CIM-24 (24 fl-lb, capacity) pendulmn type tester. The fracture toughness tests were performed usinG a Fractometer lI system. Detailed procedures of chevron notched short bar fracture toughness testing have been prevxously presented[10]. In this investigation, fracture toughness testing was conducted at a load-line displacement rate of 5.8x10-3 mm/s. For all samples, crack initiation and growth were extremely stable. All of the short rod specimens in this study failed with e x i , ~ e l y planar crack fronts, and in vimmlly all cases crack tip plasticity was such that linear elastic fracture mechanics (LEFM) conditions prevailed [11]. These observations are consistent with the results of Lewandewski, et. al.[12], who observed good correlation of short rod fracture toughness results with tests which used ASTM E-399 compact tension specimens for the case of a particulate reinforced 7000 series aluminum alloy. 3. Results and Discussion Both particulate type and size wine found to have a strong effect on both tensile suength and fracture toughness values. The results for fracture toughness are shown in Figure 1. The results for the two sample orientations differed less than 5%, and are simply averaged in subsequent figures. For a specific particle size, composites with sificon carbide reinforcement were found to possess higher levels of fracture toughness compared to those with boron carbide reinforcement. In general, the composites with larger particle sizes lead to higher values of fracture toughness, but smaller particle sizes tend to have superior tensile prvpe~es (see Table ID. Plotting the fi'acture toughness results as a function of ultimate tensile strength, Figure 2a, indicates that fracture toughness generally decreases with increasing tensile strength. The composites with silicon carbide reinforcement appear to exhibit a better combination of strength and fracture toughness than those with boron carbide reinforcement. However, it should be emphasized that an absolute comparison is complicated by the small variations in volume fraction listed in Table I. Note also that the fracture toughness of uureinfon:ed A356 alloy is siiznificantly higher than that of any of the composites, while it has much lower tensile strength compared to the strongest composites studied. The composites with 30 tun size particulate (02 and 4) tended to have higher elongation and lower fracture toughness, see Figure 2b. All of the composites had ductility values which were much less than that of uureinforced A356 alloy. The low ductility at particulate volume fractions of 25-30 vol% are consistent with the results of McDanels [13], who observed relatively brittle fracture behavior and ductility o f - l % for composites of 6061, 7075 and 2024 alloys reinforced with 30 vol% silicon carbide particulate fabricated using powder metallurgy techniques, including hot pressing, followed by extrusion and cross-roliing. As shown in Table IT, all of the composites studied had uniformly low values of impact strength ((~arpy V-Notch (CVN) Energy ): less than 20% of the level for the unreinforced A356 alloy. As in the case of the fracture toughness specimens, no significant difference between the two sample orientations was observed. No trend in impact strength is observed as a function of tensile suength or panicle type. Similarly, no correlation could be observed between impact strength and fracture toughness. These values of CVN Energy are comparable to the results of Crowe, et al [14] for 6061T6 reinforced with 20 vol% silicon carbide whiskers. Major differences in microstrueture were observed for composites with silicon carbide as opposed to boron carbide reinforcement. Typical microstructures for materials #2 (A357/30% SIC/30 ttm) and #4 (A356/25% B4C/30 txm) are shown in Figure 3. The composites with silicon carbide reinforcement tended to exhibit uniform precipitation of the Mg2Si phase throughout the matrix, while the composites with boron carbide reinforcement had little or no Mg2Si phase precipitated within the matrix. Significant reaction zones (~1 ~ thick) were observed adjacent to the particle-matrix interface in the boron carbide reinforced composites. Scanningtranmnission electron microscopy has identified the phases in the outer part of the reaction zone as AlxMg(l-x)B2,A]4C3, as well as Ti(O,B) and complex compounds of the type AIx(B,C,O)y at the particle/matrix interface. This is in distinct contrast with the A357/SIC interface, which was sharp and straight, without any reaction zone. The presence of large amounts of silicon in the A356 and A357 matrix apparently tends to decrease the driving force for reaction between the matrix and silicon carbide during casting of the composite. This was demonstrated by Iseki, et al [15], who observed a decrease in the extent of reaction between SiC and liquid aluminum when silicon additions were made to the liquid aluminum. The extent of reaction at the particle/matrix interface appears to influence the load wansfer across the interface, and causes significant differences in fracture surface features. Figure 4 shows two different magnification views of the fracture surface of composite #2 (A357/30% SIC/30 tun) from a short rod fracture toughness specimen. Multiple-faceted fracture of the SiC particulates was observed on the fracture surface using SEM. Moreover, multiple-cracked particulates were observed below the main crack fracture surface extending to depth 1 to 2 times the intexparticulate distance. A significant fraction of particles showed evidence of ductile decohesion fracture of the matrix at the panicle/matrix interface, as in Figure 4b. Both of these features suggest good transfer of load across the panicle/matrix interface for the case of silicon carbide reinforced A356/7. This behavior is in distinct contrast with the fracture surface features observed in the boron carbide reinforced material. Figure 5 shows views of the fracture surface of a short rod specimen from material #4 (A356/25%
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B4C/30 ~m): few particles exhibit secondary cracking to a .~milar degree as found in composite #2, and there is evidence of void formation and growth around malrix-precipitated particles in the reaction zone at the B4C/mauix interface (see Figure Yo). In contrast with the silicon carbidc reinforced material, those with B4C exhibited mostly single-facewxl U~svcrse fracture of particles on the fracture surface and cracked B4C panicles were not readily found beneath the main crack fracture surface. For the case of tensile specimens, fracture initiation occurred at larger than average particles, as indicated by the convergence of classic "fiver patterns" on the fracture s/u'face. In one tensile specimen from m~tedal #3, SEM/Fmergy dispersive analysis identified ,he initiation particles as A1203 inclusions. Higher magnification SEM of these fracture surfaces revealed features similar to those discussed above for the fracture toughness specimens. The results of this investigation suggest that for the same volume fraction of reinforcement, the highest fracture toughness is obtained in composites with relatively large particulates, while the best tensile properties are realized with smaller particulates. Recent work [16] has suggested that an optimum combination of yield slrength and fracture toughness in MMC's can be obtained if these two quantifies are plotted as a function of interparticle spacing, 7,, with the MMC yield stress calculated using either shear-lag [17] or thermal s~'ain mL~ma.tch/dislocafionloop punching [18] models, and the fracture toughness calculated using a modified dimple height correlation incorporating ~. The calculated values of 7.for the composites in this study are shown in Table I, and range from 55 to 229 ~m. While there is insufficient yield strength data in this investigation (due to the high reinforcement volume fractions and low ductilities) to comment on the yield strength models, the predicted values of fracture toughness in [16] as 7.increases >10 lan are much grcate~ than the experimental values in this study. The ex~emely high predicted values of fracture toughness from the calculations in [16] indicate that the fracture toughness model based primarily on a simple linear relationship between dimple height of void-initiating particles and interparficle space may not be entirely appropriate. It has been suggested [19] that ductile fracture tongimess models with a primary void link-up premise, neglecting secondary particles, void/void interactions and s~'ess state contributions, tend to overpredict fracture toughness.
This experimental study has indicated that for cast and heat Zreated aluminum alloy MMC's with a high silicon ~ontent matrix, silicon carbide reinforcement appears to provide a better combination of fractore toughness and tensile l~roperties than boron carbide. For both types of reinforcement, composites with the largest particulate sizes tended to exhibit the highest fracture toughness. Conversely, the best tensile properties were obtained in composites with the smaller O0 lun) particulate size. Microslructural analysis of fractured specimens has indicated that a significant amount of reaction has occurred between the boron carbide and the A356/7 maU'ix,while the silicon carbide/A356/7 interfaces tended to be free of any reaction zone. This suggests that load transfer mechanisa~s may differ across the particle/matrix interface of silicon carbide and boron carbide reinforced composites. Indeed, the fracture surface morphology of the p.ardcle/m.a.~.. inte.ff.ace. differed considerably for the two types of reinforcement. Finally, the values of fracture toughness increase vnth paructe raze and interparficulate spacing. Acknowledmnents The assistance of Mr. Frank Vigil on tensile testing and impact testing, and T. Sage and A. Gardes for fracture touglmess testing and metallography is appreciated. We also thank F. Greulich for his assistance in electron microscopy. The authors would like to acknowledge the support of both Sandia National Laboratories, under comract number DE-AC04-76DP00789, and the Lockheed Independent Research and Development Program. References 1. S.G. Fishman, ASTM Standardization News, Vol. 14,(#10), October, 1986, p. 46. 2. F.M. Hosking, F.F. Portillo, R. Wunderlin and R. Mehrabian, Journal of Materials Science, VUl. 17 (1982), pp. 477-498. 3. A.K. Gupta, T.K. Dan and P.IC Rohatgi, Journal of Materials Science, Vol. 21 (1986), pp. 3413-3419. 4. "Standard Methods for Tension Testing of Metallic Materials (E8-86)", 1987 Annual Book of ASTM Standards. Sec. 3, %oi. 03.01, p. 188 (1987). 5. "Standard Methods for Notched Bar Impact Testing of Metallic Materials (E23-86)", 1987 Annual Book of ASTM Standards. Sec. 3, Vol. 03.01, p. 294 (1987). 6. D. Breek, l~.!ementarvEnt,ineenne Fracture Mechanics. Sijthoffand Noordhoff, 1978. p. 284. 7. T.G. Nieh and R.F. Karlak, Scfipta Metallurgica, Vol. 18 (1984), pp. 25-28. 8. W.C. Harrigan, Jr., presented at the Annual Meeting of TMS-AIME, Los Angeles, February, 1984. 9. J.J. Stephens. Unpublished research, Sandia National Laboratories, January, 1987. 10. TerraTek Fractometer H Owner's Manual, Salt Lake City, Utah 84108, 1984 11. L. M. Barker, International Journal of Fracture, %'ol 15 (1979) pp. 515-536. 12. J.J. Lowandowski,C. Liu and W.H. Hunt, Jr., "Microstructural Effects on the Fracture Micromechanisms in 7XXX AI P/M-SiC Particulate Metal Matrix Composites" in Powder MetallurL,v Comr)c~ites, eds. P. Kumar, A. Ritter and K. Vedula, 1988, (in press). 13. D.L. McDanels, Metallurgical Transactions, Vol. 16A (1985), pp. 1105-1115.
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14. C.R. Cmwe, R.A. Gray and D.F. ~ "Microslructure Comrolled Fracture TOuehne-~Sof SiC/AI Metal Matrix ~ t e s , " in Prec. Fifth International Conference on Comvosite Materials. ICCM-V, eds. W.C. Harrigan, Jr., J. Strife, and A.K. Dhingra, 1985,p. 843. 15.T. Iseki, T. Kameda and T. Maruysma, Journal of Materials Science, Vol. 19 (1984), pp. 1692-1698. 16.R.I-L Jones, C.A. Lavender and M.T. Smith, Scripta Metallurgica, Vol. 21 (1987), pp. 1565-1570. 17. V.C. Nardone and K.M. Prewo, Scripta Metallurgica, Vol. 20 (1986), pp. 43-48. 18.R.J. Arsenanh and N. Shi, Materials Science and Engineering, Vol. 81 (1986), p. 175. 19. D.M. Tracey, Engineering Fracture Mechanics, Vol. 3, (1971), p. 301.
Table I - Combinations of matrix, reinfotr,ement type. volume fraction, and particulate size studied. Ma,eri~ 1 2 3 4 5 6 7
Mttrix
Avg. Particle ftim.flmfl
Seinfo,t z m ~ t V o l ~
A356" A357" A356" A356" A357"* A357"* A357"*
NONE SiC SiC B4C B4C B4C B4C
. 30 30 25 25 30 28
.
.
.
Calculated Particle Spacin~fum~+ .
30 100 30 100 100 122
55 183 60 200 183 229
* Cast, and HIPped by supplier, then heat treated to T6 condition: 530°C/17 h sulutionize + hot water quench, aged 163°C/51x ** Cast, HIPped sud heat treated by supplier (DACC). + Intetparticle spacing, X, calculated from 7.= ((Average Particle Size)2/Vol. Fraction)t/2
Table II- Average room temperatme mechanical properties of the composites studied: T6 condition. Duplicate tests unless otherwise noted. ~TImlAL
Uy (0.1%) (MPa)
ity (0.2%~ (MPa)
Strrs (MPa)
~ (%)
K~cSR (Mpa,~m)
(Joules)
#I: A356 unreinforced
220.3
237.1
253.9
1.14
38.4
3.13
#2 : A357/SIC 30%/30 ~m
349.9**
377.4*
374.3**
0.50**
18.6
0.56
204.7
0.18
22.0
0.49
344.4
0.67
13.6
0.56
303.2***
0.35***
19.3
0.56
#6: A357/B4C 30%/100 ILm
309.3
0.29
19.7
0.41
#7: A357/B4C 28%/122 tim
240.1
0.26
18.2
0.42
#3: A356/SIC 30%/100 pro3 #4: A356/B4C 25%/30 lain
297.7
#5: A357/B4C 25%/100 [Lm
306.6*
* only 01~ test attained this Oy. ** average of two tests. *** average of three tests.
326.0
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3O
E -'F
A356/7 Matrix 25
m 4) ciFigure 1. Effect of particle size on fracture toughne25. A356/7 alnmlnum MMC's with either SiC or B4C particulate, T6 condition. Volume fractions not equal to 30% are
indicated.
20
:3 jo-
.o :3 o
• ~
•
28
15
.m
25
ii
10 20
40
I
I
I
I
60
80
100
120
140
Particle Size, gm 40 E Unreinforced A356
m
a.
c6 30 4)
ttFigure 2. Correllafion betweentensile properties andfracture toughnessfor A356/7 aluminum MMC's with either SiC or B4C particulate, T6 condition. Data for unreinforced A356 alloy is also shown. Volume fractions not equal to 30% are indicated. (A). Ultimate tensile strength vs. fracture toughness. (B). Tensile elongation vs. fracture toughness.
:3 O 4)
[] ~
SiC
20
L_ e~
B4C
o u.
10 150
•
25
I
I
I
I
200
250
300
350
400
Ultimate Tensile Strength, MPa
E
g
4O
Unreinforced A356 - •
t~
IX
~,,~C i"
e-
1
20 28
e
~
O
I-
L_ O s_ LI.
15
10 0.00
B4C
~
I
I
0.25
0.50
% Elongation
25
g, 1.00
1.25
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Figure 3. Optical micrographs inuswafing diffeaencesin matrix precipitation. Both composites heat treated to I"6 condition. (A) Material #4, A356/'25% B4C./30)am. (B) Material #2, A357/30% SIC/30 tun.
A357 + 30% SiC (30 pm)
A356 + 25% B4C (30 Jam)
Figure 4. ScanningelecU'onmicrographs of the acture surface of a short rod fracture toughness sample of Material #2 ( A357/30% SIC/30 gin). Secondary cracking of the SiC particulate is apparent.
•
Figure 5. Scanningele~lron mimeographsof the surface of a short rod fracture toughness &~mpl¢o f ~ @4 (A356/25~ B,~Cd]0/.tin). Eiriden~eof void formation is apparent at the maa'ix/re~or.emem int~ace.
i:" i
. ~ "3,, ~
::,Ira a
8
Vol. 2 2 , pp. 1307-1312, 1988 Printed in the U.S.A.
Pergamon Press plc
CAST Al - 7 Si COMPOSITES: EFFECT OF PARTIt"I.E TYPE AND SIZE ON MECHANICAL PROPERTIES J.J. Stephens*, J.P. Lucas+ and F.M. Hosking++ *Lockheed Missiles and Space Co., Inc. Pale Alto, CA 94304 +Sandia National Laboratories Livermore, CA 94550 ++Sandia National Laboratories Albuquerque, NM 87185
(Received April 2 7 , 1988) (Revised June 7, 1988)
1. Inm~ducfion Al-minum-based metal matrix composites (MMCs) offer significant increases in modulus and strength over their unreinforeed counterparts, and in recent years have been the focus of intensive efforts aimed at understanding the relation between processing and properties of these materials. Discontinuous reinforcement aluminum MMC's are especially auxactive for applications where low thermal expansion coefficient along with low density are desired. While the majority of discontinuous a]umimun M]VIC'sconsidered for aerospace applications have been processed with powder metallurgical methods, cast aluminum MMCs offer significant cost advantages over their powder processed counterparts [1]. With the recent commercial availability of cast al,~mlnum alloy composites, micrnsuuctoral and mechanical characterization has become an important issue in composite development. Issues such as the araount of ductility possible for reinforced alloys in the solution treated and peak aged condition become quite important when one is considering using matrix alloys which tend to have low ductility in the heat treated condition. The present study was undertaken in order to characterize the effect of particulate type and size on room temperature mechanical prope~ies for cast Al-7 Si (alloys A356 and A357) MMCs. In particular, we have examined whether a specific reinforcement type (B4C or SIC), and size can provide a useful combination of tensile, fracture toughness and impact properties. 2. Exverimental The. .present investigation focuses exclusively on cast ~uminum MMCs with matrices of either A356 (nominal composluon, wt%: Al - 7 St - 0.35 Mg) or A357 (AI - 7 Sl - 0.55 Mg). For purposes of comparison with unreinfomed a
mold
of,,356
was also
ind L The
were
I
al Aluminom
pomtes Corporation, La Jolla, Ca, m the form of 19.5 x 3.4 x 1.25" (L x T x ST directions) bars. The various combinations of m,trix, reinforcement and particle size studied are shown in Table I. All of the composites had reinfmcement volume fractions between 25 and 30 vol%. Exact details are considered proprietary by the manufacturer, but basic process~g steps include surface preparation of the reinfon:ements and casting of the composites using a process such as stir..casting [2,3] which permits the r e i n f ~ c n t to remai, suspended in the molten metal. In all cases, the cast MMC's were hot isostafically pressed (HIPped) in order to reduce casting porosity. Sheet specimens cut from the bar with a longitodin,l orientation were machined for tensile testing with a gage length of 2" and a cross section of 0.5 x 0.5" [4]. Specimens were machined for impact testing with a cross section of 0.395 x 0.395", 2.165" long and a 45 ° notch with a depth of 0.080" [5], in both the L-T and T-L orientations [6]. Short bar fracture toughness specimens were machined with a width of 1", length of 1.5" and 0.87" thickness, in both the T-L and ST-L orientations. Test specimens for materials #1-4 were solution Ireated and aged in salt pots using standard heat treating procedures for the unreinfor~d A356 alloy (see Table I); materials #5-7 were heat treated by the manufacturer. Since the presence of a reinforeement can lead to er,celeratiun of aging kinetics [7-9], test coupons of materials #2-4 were aged for various times at 163°C in order to establish the time to peak age for these composites. Results for these MMC's indicated that the times for peak hardness fell within the range of tunes requited (4-5 hr) at 163°C for peak aging of the un~inforeedA356 alloy. This lack of aecelenUion in aging kinetics is apparently due to the relatively large size of the reinforcement. Nevertheless, other
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investigators have observed acceleration in aging kinetics - about a factor of 2 in time to peak hardness - in 6061 Al alloy with fine (-5 ixm size) reinforcements of either B4C or SiC [7,8]. Separate work in our laboratory on 6061/SIC and 7091/B4C composites with 30 vol%, 5 ~ size particulate has also shown acceleration in aging kinetics [9] relative to the unreinforced alloys. Tensile specimens were tested to fracture at room temperature using a MTS 810 servohydranlic testing machine equipped with a 50 kip capacity actuator and 30 kip capacity wedge type gips. A cfip-on extensometer with a gage length of 1" was used to measure strain. All tests were ran at an engineming strain rate of 10-3 s-l. The impact tests were performed using a ManI abs Model CIM-24 (24 fl-lb, capacity) pendulmn type tester. The fracture toughness tests were performed usinG a Fractometer lI system. Detailed procedures of chevron notched short bar fracture toughness testing have been prevxously presented[10]. In this investigation, fracture toughness testing was conducted at a load-line displacement rate of 5.8x10-3 mm/s. For all samples, crack initiation and growth were extremely stable. All of the short rod specimens in this study failed with e x i , ~ e l y planar crack fronts, and in vimmlly all cases crack tip plasticity was such that linear elastic fracture mechanics (LEFM) conditions prevailed [11]. These observations are consistent with the results of Lewandewski, et. al.[12], who observed good correlation of short rod fracture toughness results with tests which used ASTM E-399 compact tension specimens for the case of a particulate reinforced 7000 series aluminum alloy. 3. Results and Discussion Both particulate type and size wine found to have a strong effect on both tensile suength and fracture toughness values. The results for fracture toughness are shown in Figure 1. The results for the two sample orientations differed less than 5%, and are simply averaged in subsequent figures. For a specific particle size, composites with sificon carbide reinforcement were found to possess higher levels of fracture toughness compared to those with boron carbide reinforcement. In general, the composites with larger particle sizes lead to higher values of fracture toughness, but smaller particle sizes tend to have superior tensile prvpe~es (see Table ID. Plotting the fi'acture toughness results as a function of ultimate tensile strength, Figure 2a, indicates that fracture toughness generally decreases with increasing tensile strength. The composites with silicon carbide reinforcement appear to exhibit a better combination of strength and fracture toughness than those with boron carbide reinforcement. However, it should be emphasized that an absolute comparison is complicated by the small variations in volume fraction listed in Table I. Note also that the fracture toughness of uureinfon:ed A356 alloy is siiznificantly higher than that of any of the composites, while it has much lower tensile strength compared to the strongest composites studied. The composites with 30 tun size particulate (02 and 4) tended to have higher elongation and lower fracture toughness, see Figure 2b. All of the composites had ductility values which were much less than that of uureinforced A356 alloy. The low ductility at particulate volume fractions of 25-30 vol% are consistent with the results of McDanels [13], who observed relatively brittle fracture behavior and ductility o f - l % for composites of 6061, 7075 and 2024 alloys reinforced with 30 vol% silicon carbide particulate fabricated using powder metallurgy techniques, including hot pressing, followed by extrusion and cross-roliing. As shown in Table IT, all of the composites studied had uniformly low values of impact strength ((~arpy V-Notch (CVN) Energy ): less than 20% of the level for the unreinforced A356 alloy. As in the case of the fracture toughness specimens, no significant difference between the two sample orientations was observed. No trend in impact strength is observed as a function of tensile suength or panicle type. Similarly, no correlation could be observed between impact strength and fracture toughness. These values of CVN Energy are comparable to the results of Crowe, et al [14] for 6061T6 reinforced with 20 vol% silicon carbide whiskers. Major differences in microstrueture were observed for composites with silicon carbide as opposed to boron carbide reinforcement. Typical microstructures for materials #2 (A357/30% SIC/30 ttm) and #4 (A356/25% B4C/30 txm) are shown in Figure 3. The composites with silicon carbide reinforcement tended to exhibit uniform precipitation of the Mg2Si phase throughout the matrix, while the composites with boron carbide reinforcement had little or no Mg2Si phase precipitated within the matrix. Significant reaction zones (~1 ~ thick) were observed adjacent to the particle-matrix interface in the boron carbide reinforced composites. Scanningtranmnission electron microscopy has identified the phases in the outer part of the reaction zone as AlxMg(l-x)B2,A]4C3, as well as Ti(O,B) and complex compounds of the type AIx(B,C,O)y at the particle/matrix interface. This is in distinct contrast with the A357/SIC interface, which was sharp and straight, without any reaction zone. The presence of large amounts of silicon in the A356 and A357 matrix apparently tends to decrease the driving force for reaction between the matrix and silicon carbide during casting of the composite. This was demonstrated by Iseki, et al [15], who observed a decrease in the extent of reaction between SiC and liquid aluminum when silicon additions were made to the liquid aluminum. The extent of reaction at the particle/matrix interface appears to influence the load wansfer across the interface, and causes significant differences in fracture surface features. Figure 4 shows two different magnification views of the fracture surface of composite #2 (A357/30% SIC/30 tun) from a short rod fracture toughness specimen. Multiple-faceted fracture of the SiC particulates was observed on the fracture surface using SEM. Moreover, multiple-cracked particulates were observed below the main crack fracture surface extending to depth 1 to 2 times the intexparticulate distance. A significant fraction of particles showed evidence of ductile decohesion fracture of the matrix at the panicle/matrix interface, as in Figure 4b. Both of these features suggest good transfer of load across the panicle/matrix interface for the case of silicon carbide reinforced A356/7. This behavior is in distinct contrast with the fracture surface features observed in the boron carbide reinforced material. Figure 5 shows views of the fracture surface of a short rod specimen from material #4 (A356/25%
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B4C/30 ~m): few particles exhibit secondary cracking to a .~milar degree as found in composite #2, and there is evidence of void formation and growth around malrix-precipitated particles in the reaction zone at the B4C/mauix interface (see Figure Yo). In contrast with the silicon carbidc reinforced material, those with B4C exhibited mostly single-facewxl U~svcrse fracture of particles on the fracture surface and cracked B4C panicles were not readily found beneath the main crack fracture surface. For the case of tensile specimens, fracture initiation occurred at larger than average particles, as indicated by the convergence of classic "fiver patterns" on the fracture s/u'face. In one tensile specimen from m~tedal #3, SEM/Fmergy dispersive analysis identified ,he initiation particles as A1203 inclusions. Higher magnification SEM of these fracture surfaces revealed features similar to those discussed above for the fracture toughness specimens. The results of this investigation suggest that for the same volume fraction of reinforcement, the highest fracture toughness is obtained in composites with relatively large particulates, while the best tensile properties are realized with smaller particulates. Recent work [16] has suggested that an optimum combination of yield slrength and fracture toughness in MMC's can be obtained if these two quantifies are plotted as a function of interparticle spacing, 7,, with the MMC yield stress calculated using either shear-lag [17] or thermal s~'ain mL~ma.tch/dislocafionloop punching [18] models, and the fracture toughness calculated using a modified dimple height correlation incorporating ~. The calculated values of 7.for the composites in this study are shown in Table I, and range from 55 to 229 ~m. While there is insufficient yield strength data in this investigation (due to the high reinforcement volume fractions and low ductilities) to comment on the yield strength models, the predicted values of fracture toughness in [16] as 7.increases >10 lan are much grcate~ than the experimental values in this study. The ex~emely high predicted values of fracture toughness from the calculations in [16] indicate that the fracture toughness model based primarily on a simple linear relationship between dimple height of void-initiating particles and interparficle space may not be entirely appropriate. It has been suggested [19] that ductile fracture tongimess models with a primary void link-up premise, neglecting secondary particles, void/void interactions and s~'ess state contributions, tend to overpredict fracture toughness.
This experimental study has indicated that for cast and heat Zreated aluminum alloy MMC's with a high silicon ~ontent matrix, silicon carbide reinforcement appears to provide a better combination of fractore toughness and tensile l~roperties than boron carbide. For both types of reinforcement, composites with the largest particulate sizes tended to exhibit the highest fracture toughness. Conversely, the best tensile properties were obtained in composites with the smaller O0 lun) particulate size. Microslructural analysis of fractured specimens has indicated that a significant amount of reaction has occurred between the boron carbide and the A356/7 maU'ix,while the silicon carbide/A356/7 interfaces tended to be free of any reaction zone. This suggests that load transfer mechanisa~s may differ across the particle/matrix interface of silicon carbide and boron carbide reinforced composites. Indeed, the fracture surface morphology of the p.ardcle/m.a.~.. inte.ff.ace. differed considerably for the two types of reinforcement. Finally, the values of fracture toughness increase vnth paructe raze and interparficulate spacing. Acknowledmnents The assistance of Mr. Frank Vigil on tensile testing and impact testing, and T. Sage and A. Gardes for fracture touglmess testing and metallography is appreciated. We also thank F. Greulich for his assistance in electron microscopy. The authors would like to acknowledge the support of both Sandia National Laboratories, under comract number DE-AC04-76DP00789, and the Lockheed Independent Research and Development Program. References 1. S.G. Fishman, ASTM Standardization News, Vol. 14,(#10), October, 1986, p. 46. 2. F.M. Hosking, F.F. Portillo, R. Wunderlin and R. Mehrabian, Journal of Materials Science, VUl. 17 (1982), pp. 477-498. 3. A.K. Gupta, T.K. Dan and P.IC Rohatgi, Journal of Materials Science, Vol. 21 (1986), pp. 3413-3419. 4. "Standard Methods for Tension Testing of Metallic Materials (E8-86)", 1987 Annual Book of ASTM Standards. Sec. 3, %oi. 03.01, p. 188 (1987). 5. "Standard Methods for Notched Bar Impact Testing of Metallic Materials (E23-86)", 1987 Annual Book of ASTM Standards. Sec. 3, Vol. 03.01, p. 294 (1987). 6. D. Breek, l~.!ementarvEnt,ineenne Fracture Mechanics. Sijthoffand Noordhoff, 1978. p. 284. 7. T.G. Nieh and R.F. Karlak, Scfipta Metallurgica, Vol. 18 (1984), pp. 25-28. 8. W.C. Harrigan, Jr., presented at the Annual Meeting of TMS-AIME, Los Angeles, February, 1984. 9. J.J. Stephens. Unpublished research, Sandia National Laboratories, January, 1987. 10. TerraTek Fractometer H Owner's Manual, Salt Lake City, Utah 84108, 1984 11. L. M. Barker, International Journal of Fracture, %'ol 15 (1979) pp. 515-536. 12. J.J. Lowandowski,C. Liu and W.H. Hunt, Jr., "Microstructural Effects on the Fracture Micromechanisms in 7XXX AI P/M-SiC Particulate Metal Matrix Composites" in Powder MetallurL,v Comr)c~ites, eds. P. Kumar, A. Ritter and K. Vedula, 1988, (in press). 13. D.L. McDanels, Metallurgical Transactions, Vol. 16A (1985), pp. 1105-1115.
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14. C.R. Cmwe, R.A. Gray and D.F. ~ "Microslructure Comrolled Fracture TOuehne-~Sof SiC/AI Metal Matrix ~ t e s , " in Prec. Fifth International Conference on Comvosite Materials. ICCM-V, eds. W.C. Harrigan, Jr., J. Strife, and A.K. Dhingra, 1985,p. 843. 15.T. Iseki, T. Kameda and T. Maruysma, Journal of Materials Science, Vol. 19 (1984), pp. 1692-1698. 16.R.I-L Jones, C.A. Lavender and M.T. Smith, Scripta Metallurgica, Vol. 21 (1987), pp. 1565-1570. 17. V.C. Nardone and K.M. Prewo, Scripta Metallurgica, Vol. 20 (1986), pp. 43-48. 18.R.J. Arsenanh and N. Shi, Materials Science and Engineering, Vol. 81 (1986), p. 175. 19. D.M. Tracey, Engineering Fracture Mechanics, Vol. 3, (1971), p. 301.
Table I - Combinations of matrix, reinfotr,ement type. volume fraction, and particulate size studied. Ma,eri~ 1 2 3 4 5 6 7
Mttrix
Avg. Particle ftim.flmfl
Seinfo,t z m ~ t V o l ~
A356" A357" A356" A356" A357"* A357"* A357"*
NONE SiC SiC B4C B4C B4C B4C
. 30 30 25 25 30 28
.
.
.
Calculated Particle Spacin~fum~+ .
30 100 30 100 100 122
55 183 60 200 183 229
* Cast, and HIPped by supplier, then heat treated to T6 condition: 530°C/17 h sulutionize + hot water quench, aged 163°C/51x ** Cast, HIPped sud heat treated by supplier (DACC). + Intetparticle spacing, X, calculated from 7.= ((Average Particle Size)2/Vol. Fraction)t/2
Table II- Average room temperatme mechanical properties of the composites studied: T6 condition. Duplicate tests unless otherwise noted. ~TImlAL
Uy (0.1%) (MPa)
ity (0.2%~ (MPa)
Strrs (MPa)
~ (%)
K~cSR (Mpa,~m)
(Joules)
#I: A356 unreinforced
220.3
237.1
253.9
1.14
38.4
3.13
#2 : A357/SIC 30%/30 ~m
349.9**
377.4*
374.3**
0.50**
18.6
0.56
204.7
0.18
22.0
0.49
344.4
0.67
13.6
0.56
303.2***
0.35***
19.3
0.56
#6: A357/B4C 30%/100 ILm
309.3
0.29
19.7
0.41
#7: A357/B4C 28%/122 tim
240.1
0.26
18.2
0.42
#3: A356/SIC 30%/100 pro3 #4: A356/B4C 25%/30 lain
297.7
#5: A357/B4C 25%/100 [Lm
306.6*
* only 01~ test attained this Oy. ** average of two tests. *** average of three tests.
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E -'F
A356/7 Matrix 25
m 4) ciFigure 1. Effect of particle size on fracture toughne25. A356/7 alnmlnum MMC's with either SiC or B4C particulate, T6 condition. Volume fractions not equal to 30% are
indicated.
20
:3 jo-
.o :3 o
• ~
•
28
15
.m
25
ii
10 20
40
I
I
I
I
60
80
100
120
140
Particle Size, gm 40 E Unreinforced A356
m
a.
c6 30 4)
ttFigure 2. Correllafion betweentensile properties andfracture toughnessfor A356/7 aluminum MMC's with either SiC or B4C particulate, T6 condition. Data for unreinforced A356 alloy is also shown. Volume fractions not equal to 30% are indicated. (A). Ultimate tensile strength vs. fracture toughness. (B). Tensile elongation vs. fracture toughness.
:3 O 4)
[] ~
SiC
20
L_ e~
B4C
o u.
10 150
•
25
I
I
I
I
200
250
300
350
400
Ultimate Tensile Strength, MPa
E
g
4O
Unreinforced A356 - •
t~
IX
~,,~C i"
e-
1
20 28
e
~
O
I-
L_ O s_ LI.
15
10 0.00
B4C
~
I
I
0.25
0.50
% Elongation
25
g, 1.00
1.25
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Figure 3. Optical micrographs inuswafing diffeaencesin matrix precipitation. Both composites heat treated to I"6 condition. (A) Material #4, A356/'25% B4C./30)am. (B) Material #2, A357/30% SIC/30 tun.
A357 + 30% SiC (30 pm)
A356 + 25% B4C (30 Jam)
Figure 4. ScanningelecU'onmicrographs of the acture surface of a short rod fracture toughness sample of Material #2 ( A357/30% SIC/30 gin). Secondary cracking of the SiC particulate is apparent.
•
Figure 5. Scanningele~lron mimeographsof the surface of a short rod fracture toughness &~mpl¢o f ~ @4 (A356/25~ B,~Cd]0/.tin). Eiriden~eof void formation is apparent at the maa'ix/re~or.emem int~ace.
i:" i
. ~ "3,, ~
::,Ira a
8