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Fracture toughness and fracture mechanisms of Al-Al2O3 composites at cryogenic and elevated temperatures
lW ELSEVIER
Materials Science and Engineering A206 (19961 183 193
A
Fracture toughness and fracture mechanisms of A1-A1203 composites at cryogenic and elevated temperatures P. Poza, J. Llorca Department q[' Materials Science, Polytechnic Universitj' of Madrid, E.T.S. de lngenieros de Caminos, 28040 Madrid, Spain Received 31 May 1995: in revised form 13 July 1995
Abstract
The fracture toughness of two Al alloys (2014 and 6061) reinforced with 15 vol.% AI20 3 particulates was measured in the temperature range from - 150 °C to 300 °C. Both composites maintained acceptable levels of toughness (over 20 MPa m ~';2)up to 200 °C, but the fracture toughness dropped very quickly above this temperature. Fracture always took place by a ductile mechanism involving the nucleation of voids from the ceramic particulates, the progressive growth of these voids, and the final coalescence through the matrix. Quantitative microscopy analyses demonstrated that voids were initiated from reinforcement fracture at low temperatures and by failure at the interface or in the matrix near the interface at elevated temperatures. Regardless of these differences, the changes in toughness could be explained by the deterioration in the mechanical properties with temperature rather than by the microstructural changes induced by the exposure to elevated temperatures. Keywords: Fracture; Composites; High temperature; Cryogenic temperatures; Toughness
1. Introduction
The low fracture toughness of discontinuously reinforced metal-matrix composites (MMCs) has been one of the main drawbacks to their industrial application. The plane strain fracture toughness of Al-based composites is usually in the range 10-25 M P a m ~/2 [1] and this factor, together with their reduced tensile ductility, has precluded their use in structural elements in the aerospace industry. As a result, the fracture behaviour of particulate- and whisker-reinforced M M C s has received considerable attention from the research community in quest of the mechanisms responsible of this brittle behaviour [2-11]. The investigations of the fracture toughness and the fracture mechanisms in M M C have been focused on the ambient temperature behaviour, and in fact data in the fracture toughness above 100°C are scarce [12-14]. Somerday et al. [12] measured the fracture resistance of a 2009 AI alloy reinforced with 20 vol.% SiC particulates in the temperature range 25-316 °C. They found that the fracture toughness increased slightly from 20 MPa m 12 at ambient temperature to a m a x i m u m of 29 MPa m ~'2 at 220°C. Beyond this temperature, the 0921-5093/96/$15.00 ,~2 1996 S S D I 0921-5093(95)09999-9
Elsevier Science S.A. All rights reserved
composite toughness decreased to a minimum of 14 M P a m L:2 at 316 °C. In addition, a transition from unstable to stable crack propagation was found at 175 °C. Fractographic observations indicated that crack intiation and propagation took place by microvoid nucleation, growth and coalescence in the matrix at all temperatures. The preferred void initiation site changed from the S i C - m a t r i x interface at low temperatures to within the matrix above 200 °C. A micromechanical model based on a critical strain concept was used to predict the changes in the initiation toughness and tearing modulus with the temperature, with the local fracture strain being inferred from the tensile data. In another investigation, we also measured the fracture toughness of another A I - S i C composite in the temperature range from - 1 3 6 °C to 190 °C [13]. The fracture toughness of this material was fairly constant (between 17 and 2 0 M P a m ~:2) in this temperature range and crack propagation was unstable in all cases. The mechanism of void initiation in this composite was reinforcement fracture, regardless of the temperature and aging condition, and the volume fraction of broken reinforcements ahead of the crack tip was between 10% and 15%.
P, Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
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These preliminary results seem to indicate that the fracture properties of discontinuously reinforced MMCs are maintained up to 200 °C, but it is evident that more research is needed to ascertain the influence of the temperature on the fracture toughness and fracture mechanisms in these materials. This need is more critical when it is considered that MMCs are likely to be used in high temperature applications owing to their excellent creep [15] and wear [16] resistance. In addition, precipitation- and solution-hardened A1 alloys are currently used for the storage and handling of cryogenic liquids (such as nitrogen, oxygen, natural gas, and carbon dioxide) and in other low temperature applications owing to their non-magnetic behaviour, stable microstructure, and retention of strength, ductility and toughness at very low temperatures [17]. Potential savings are also envisaged in the use of A1-SiC an AIA1203 composites for cryogenic applications but it is necessary to determine the mechanical properties of these materials under such conditions. This investigation was aimed at determining the effect of the temperature (in the range from - 150 °C to 300 °C) on the fracture resistance and the fracture mechanisms in two precipitation-hardened A1 alloys reinforced with A1203 particulates. The results of the fracture tests, together with detailed analyses of the failure mechanisms by means of quantitative microscopy, were used to elucidate the relationships between the macroscopic behaviour and the microstructural features.
2. M a t e r i a l s
Two composite materials were used in this investigation: a 2014 A1 alloy reinforced with 14.4 vol.% A1203 particulates, and a 6061 A1 alloy reinforced with 14.9 vol.% A1203 particulates. The cast materials w e r e supplied by the manufacturer (Duralcan, San Diego, CA, USA) in the form of extruded circular bars 50 mm in diameter in the peak-aged condition (T6). The heat treatments to achieve this temper in the 2014 A1 based composite included a solution heat treatment at 502 °C for 2 h, water quenching, and aging at 160 °C for 16 h. The heat treatments in the 6061 A1 based material w e r e the same, although the solution heat treatment temperature was 560 °C. The chemical compositions of the
Table 2 Geometric characteristics of the AI203 particulates in the longitudinal direction Composite
AI 2014 + AI203 AI 6061 + A1203
Areaa (I.tm2)
Aspect ratio a
145 + 81 144 + 70
0.44 + 0.16 0.51 + 0.19
Dmin/Dmax
a+ standard deviation.
matrices in both composites were provided by the manufacturer and are given in Table 1. The microstructural characteristics of both composites w e r e detailed in a previous publication [18]. The sizes and morphologies of the ceramic particulates were very similar in the two materials, and the average particulate area and aspect ratio are given in Table 2. The particulates were oriented with their longer axis parallel to the extrusion direction (Fig. 1), and the aspect ratio is approximately the quotient between the particulate width (perpendicular to the extrusion axis) and length (parallel to the extrusion axis). The area and aspect ratio distributions were wide, as indicated by the standard deviations in Table 2. The area of 80% of the particulates was in the range 60-225 gm 2, although reinforcements of up to 500 ~tm2 were occasionally found. It is worth noting that broken A1203 particulates w e r e found in both composites prior to testing. Most of the broken reinforcements were observed in regions w h e r e the ceramic particulates were very close together in a cluster and they were shattered (Fig. 2) rather than broken with a single crack. Porosity (black regions in Fig. 2) was often associated with clusters of shattered particulates. This fragmentation of the ceramic particulates probably took place during extrusion in regions w h e r e the local volume fraction of reinforcement was high. In addition to shattered reinforcements, some isolated particulates which were not associated with a cluster were also broken. The fracture mode was differ-
Table 1 Chemical composition (weight per cent) of the matrices Si
Cu
Mg
Mn
Fe
Cr
Zn
Ti
AI
2014 A1 0.76 4.70 0.47 0.78 0.06 0,01 0,06 0.03 Remainder 6061 AI 0.62 0.24 0.94 0.04 0.07 0.10 0.02 0.01 Remainder
Fig. 1. Low magnification micrograph of the 2014 A1 based composite. Notice that the AI203 particulates are oriented with their longer axis in the extrusion direction (vertical).
P. Poza, J, LIorca / Materials' Science and Engineering A206 (1996) 18.7 193
Fig. 2. Backscattered electron micrograph of a cluster of shattered particulates in the 2014 A1 alloy reinforced with 15v01.% AI203 particulates. AI203 particulates are seen as dark grey zones, whereas AI C u - M n inclusions are white and the black areas are pores. Extrusion axis is vertical.
ent in these cases, the reinforcement being fractured by a single crack perpendicular to the extrusion direction. The fraction of isolated particulates was, however, much less important than damage associated with clusters, as indicated in Table 3, where the percentage of isolated and clustered broken AI203 particulates is given for each composite. Finally, it should be indicated that the ceramic reinforcements were well bonded to the A1 matrix, and interfacial decohesion was rarely observed. In addition to the A1203 particulates, inclusions of irregular shape and about 1 jam in diameter were distributed within the matrix of the 2014 A1 based composites (Fig. 2). They appeared to be intermetallic inclusions, whose main constituents (determined by energy dispersive X-ray microanalysis) were A1, Cu and Mn. The study by means of backscattered electrons of the 6061 A1 based material indicated that the volume fraction of inclusions in this material was very small, and areas with a chemical composition different from the matrix or the particulates were scarce. Small dispersoids (around 0.1 jam) were found in the matrix, and they were identified as Mn- and Cr-rich compounds, which inhibit recrystallization during cooling.
3. Experimental techniques Compact tension specimens were machined from the bars for the fracture tests. Specimen thickness was 18 ram, the crack being oriented in the L-T direction. Specimens were precracked in fatigue, in accordance with the recommendations of ASTM Standard E 399, in a servohydraulic mechanical testing machine. To reduce the number of fatigue cycles to initiate the fatigue crack, a through-thickness cut 1 mm in depth was given at the notch root with a very thin (0.3 ram)
185
diamond saw. Fracture tests were performed in stroke control, and the displacement rate was set to obtain a stress intensity factor rate of 0.6 M P a m 1 2 s ~. Load and crack mouth opening displacement (CMOD) were continuously monitored through a computer-controlled data acquisition system. For high and low temperature tests, the specimen was placed in a temperature chamber. Heating was by an electric resistance and liquid nitrogen was used for cooling. The specimen temperature was measured with a K-type thermocouple, which was introduced in a 2 mm deep hole machined in the specimen. Heating (or cooling) rate was always slower than 5 °C rain -~. The specimen was held for 20 min at the test temperature (_+ 2 °C) prior to testing, to ensure a homogeneous temperature distribution through the thickness. The C M O D was measured with a standard clip displacement gauge in the tests carried out below 200°C. As this gauge could not work beyond this temperature, a laser extensometer developed to measure the C M O D at high temperatures was used [19]. This extensometer is essentially similar to that reported by Carroll et al. [20]. An He Ne laser emitter sends a scanning laser beam through two silica windows across the furnace walls. Two alumina pins glued on the points which determine the C M O D are interposed in the laser beam and the resulting intensity distribution is measured by a detector placed on the opposite side of the furnace. A microcomputer processes the intensity distribution to determine the distance between the pins. Although the resolution of the system is very high, the accuracy is limited mainly by two factors: the air seepage around the loading rods and the thermal gradients along the optical paths. Preliminary calibrations of the extensometer indicated that air seepage was dominant as compared with the thermal gradients in our system, perhaps because the large size of the furnace led to smaller temperature gradients along the laser path. Thus special care was taken to reduce the chimney effect as much as possible by closing the gaps between the loading rods and the furnace. In addition, the specimen remained at the test temperature for 20 min before testing to achieve a stationary distribution of temperatures within the furnace. Moreover, the good thermal insulation of the furnace allowed the laser source and detector to be located very close, to the silica glass without danger of overheating, reducing the fluctuations produced outside the furnace. In addition to all these precautions, the measurements were averaged over 150 scans for each displacement to minimize the noise. Metallographic samples for both materials were prepared in the longitudinal (extrusion) direction from the specimens after testing to analyse the failure micromechanisms. The specimens were cut through the middle with a low speed diamond blade wheel to minimize
P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
186
damage, using oil as a lubricant, and initially polished on SiC abrasive paper to 600 grit finish. This was followed by polishing on diamond slurry (9, 3 and 1 ~tm) and finally on magnesia. The polished samples were first cleaned in deionized water and afterwards by ultrasound in acetone. Scanning electron microscope micrographs were taken randomly from the polished surfaces at less than 100 tam from the fracture surface. This distance was always less than the size of the plastic zone around the crack tip for all the materials and temperatures. These micrographs were processed by means of a computer-controlled image analysis system to determine the fraction of void nucleating particles in the composite and the origin of the voids (either particulae fracture or matrix-reinforcement decohesion). In addition, the fracture surfaces were examined by scanning electron microscopy to ascertain the main failure mechanisms for each material and temperature.
30
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-200
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0
100
200
300
Temperature (°C) Fig. 3. Fracture toughness of the composite materials as a function of the test temperature.
4. R e s u l t s
4.1. Fracture toughness The results of the fracture tests were examined according to the procedures indicated in the ASTM Standard E-399. All the tests fulfilled the conditions prescribed for fatigue precracking, crack propagation during fatigue, crack front curvature, and ratio between the load for crack initiation and the maximum load during the test. The specimen thickness also met the requirements indicated in the standard for all the tests except those carried out at 300 °C. Given the values of the yield strength reported for these composites at 300 °C by the manufacturer [21], the specimen thickness should have been 26 mm for the 2014 A1 based composite and 38mm for the 6061 A1 based material. (According to Ref. [21], the yield strength of the composites at different temperatures was measured after 30min soak time. The actual values of the yield strength as a function of the temperature are plotted in Fig. 10(a)). Thus the results presented in Fig. 3 represent the plane strain fracture toughness of the composites up to 190 °C. The data obtained at 300 °C cannot be regarded as Ktc values, properly speaking, because plane strain conditions did not prevail at the crack tip. In fact, while the fracture surfaces up to
190 °C were flat, shear lips oriented on +45 ° bands to the specimen thickness were found at 300 °C. These shear lips, indicative of plane stress conditions at the crack tip, covered around 35% of the fracture surface in the 2014 A1 based composite and around 55% in the 6061 A1 based material. Even so, it is evident that both composites maintained acceptable levels of fracture toughness in the temperature range from - 1 5 0 °C to 190 °C. On the contrary, the fracture toughness decreased by 30%40% from 190 °C to 300 °C. This result is more dramatic when it is considered that the actual values of the plane strain fracture toughness at 300 °C should be lower than those obtained from our fracture tests. It should be noted that the reduction in the resistance to crack initiation was accompanied by an increase in the resistance to crack propagation, as is shown by the load-CMOD curves presented in Fig. 4 for the 6061 A1 based composite. The crack began to propagate unstably near to the maximum load at low temperatures, while a substantial amount of stable crack growth was observed at 300°C. The tests carried out at 190°C revealed an intermediate behaviour. Similar results were observed in the other composite and were reported by other researchers [12].
Table 3 Fraction of broken A1203 particulates in the as-received materials
4.2. Fracture mechanL~ms
Composite
AI 2014 + A1203 A1 6061 + A1203
Isolated particulates
Clustered particulates
5.5 2.7
8.5 13.0
(%)
(%)
The analysis of the fracture mechanisms began with the study of the fracture surfaces in the scanning electron microscope. All the fracture surfaces presented the typical morphology of failure by a ductile mechanism involving the nucleation, growth and coalescence of
P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183 193
voids (Figs. 5 and 6). At low temperatures (between - 1 5 0 °C and 20 °C), the principal feature found in both materials was the presence of a large number of A1203 particulates on the fracture surfaces (Fig. 5). The ceramic particulates were clean and no traces of the A1 matrix were observed on them. The A1 matrix around the reinforcements was heavily deformed and another microvoid population was observed in the matrix. Microvoids were formed around the A 1 - C u - M n inclusions and smaller dispersoids and precipitates in the 2014 A1 based composite, while the Mn- and Cr-rich dispersoids were identified as microvoid nucleation sites in the 6061 A1 based material. It should also be noted that it was not possible to distinguish the fracture surfaces created at ambient temperature from those formed at - 150 °C. The fracture surfaces of the specimens tested at 300 °C also presented the dual population of voids formed around the ceramic reinforcements and the inclusions and dispersoids in the matrix. However, the majority of the A1203 particulates (whose size and morphology were clearly visible on the fracture surface) were covered with AI matrix (Fig. 6), while they were free from matrix in the fracture surfaces created at ambient and cryogenic temperatures. The fracture surfaces at 190 °C (not shown here) exhibited an intermediate behaviour where ceramic reinforcements both clean and covered with matrix could be seen. It should be noted that the average void sizes were very similar at all temperatures. The previous observations seemed to indicate a change in the void nucleation sites from the reinforce-
AI 6061 + 15 % AI203 10--,
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(°C)
Temperature -150
8 ---=---
Z
6
o .._1
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0
0.5
1
190
1.5
CMOD (mm) Fig. 4. Load C M O D curves for the 6061 AI composite reinforced with AI~O 3 particulates at different temperatures.
187
Fig. 5. Ambient temperature fracture surfaces. (a) 2014 AI alloy reinforced with AI203 particulates. (b) 6061 AI alloy reinforced with AI203 particulates. The ceramic reinforcements on the fracture surfaces were free of matrix.
ment at low temperatures to those near the interface at elevated temperatures. To check this hypothesis, the specimens tested at - 1 5 0 °C, 190 °C and 300 °C were sliced through the middle and polished, and the region around the crack path was observed in the scanning electron microscope. Three representative micrographs of the mechanisms seen at the three temperatures are shown in Fig. 7. The dominant void initiation mechanism at - 1 5 0 °C was reinforcement fracture by one crack oriented perpendicularly to the loading direction, as shown by the arrows in Fig. 7(a). Afterwards, the voids grew in the longitudinal but not in the radial direction. These observations are in agreement with the morphology of the fracture surfaces, where the most typical feature was a matrix void containing one clean A1203 particulate, the void diameter and shape being similar to those of the reinforcement. At 300 °C, on the contrary, interracial decohesion became the main void nucleation mechanism (Fig. 7(c)). Matrix-reinforcement decohesion usually took place at the end of elongated reinforcements or in regions where the ends of two reinforcements were very close in the longitudinal direction. The specimens tested at 190 °C presented both kinds of failure mechanisms, and rein-
188
P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
forcement fracture and matrix-reinforcement decohesion were observed around the fracture surfaces (Fig. 7(b)). To quantify the importance o f each mechanism, the micrographs were processed by means of a computer-controlled image analysis system to determine the fraction of void nucleating particulates in the composite and the origin of the voids (either particulate fracture or matrix-reinforcement decohesion). This study did not include the ceramic particulates shattered in clusters, because they were already broken prior to testing.
Fig. 7. Void nucleation mechanisms as a function of the test temperature. (a) 6061 AI based composite tested at -150 °C. (b) 2014 AI based material tested at 190 °C. (c) 6061 AI based composite tested at 300 °C.
Fig. 6. Fracture surfaces created at 300 °C. (a) 2014 AI alloy reinforced with AI203 particulates. (b), (c) 6061 A1 alloy reinforced with AI203 particulates. Matrix voids were observed on the ceramic reinforcements, which are easily recognized by their size and shape.
The results of this analysis for both composites are shown in Fig. 8. It shows that only a small fraction (between 15% and 20%) of the ceramic particulates acted as void nucleation sites, regardless of the matrix alloy and test temperature. In addition, the void nucleation mechanism changed from reinforcement fracture at - 1 9 0 °C to interfacial decohesion at 300 °C. Ap-
P. Poza, J, Llorca / Materials Science and Engineering A 206 (1996) 183-193
proximately the same numbers of voids were nucleated by reinforcement fracture and by interfacial decohesion at 190 °C. It is worth noting that the fracture of broken reinforcements at 300 °C was very similar to the fraction of isolated A120~ particulates fractured before testing (Table 3). Thus reinforcement fracture had practically disappeared at this temperature.
5. Discussion The experimental observations showed that void nucleation by reinforcement fracture was progressively substituted by failure at or near the reinforcement-matrix interface as the temperature increased. This effect was also reported by Zhao et al. [22] who analysed the influence of the temperature on the tensile behaviour of a 2014 AI alloy reinforced with AI203 particulates, In agreement with our results, reinforcement fracture was dominant at ambient temperature while interfacial decohesion and matrix voiding were preponderant at 300 °C. Around 200 °C, they found that the number of voids was evenly divided between both mechanisms, in perfect agreement with the results plotted in Fig. 8. The reduction in the number of broken reinforcements as a function of the temperature can be explained through the changes in the matrix mechanical properties. The stresses on the reinforcements during monotonic deformation were recently calculated by Justice et al. [23] who carried out a parametric study of the influence of the matrix properties on the reinforcement stresses by means of the finite element analysis of a unit cell, representative of the composite. The ceramic reinforcements were located at the centre of each cell. Unit cylinders (the cylinder length was equal to the diameter) were used to model reinforcements with irregular shape, sharp corners and low aspect ratios, such as those used in particulate-reinforced Al-based composites. For a given reinforcement volume fraction and shape, the numerical results were fitted to polynomial expressions of the matrix properties. The tensile properties of the 2014 unreinforced alloy in the peak-aged condition, processed in the same way, were obtained by Zhao et al. [22] up to 300 °C after 30 min of exposure and they are shown in Fig. 9(a). The matrix flow stress o-m can be expressed by a power law of the plastic strain ~ as am=
Ag "
(1)
where the coefficients A and n, which characterize the matrix stress strain behaviour, can be obtained from the results in Fig. 9(a) and they are shown in Table 4. The stresses acting on the ceramic particulates as a function of the plastic strain can then be calculated by introducing these values of A and n in the polynomial expressions reported in Ref. [23]. The results are shown in Fig. 9(b),
189
where the stresses on the reinforcements are plotted as a function of the composite plastic strain. They clearly show that the stresses on the reinforcements remain almost constant up to 100 °C, decrease slightly at 200 °C and drop to very low levels at 300 °C as a consequence of the reduction in the matrix flow stress. The evolution of the stresses on the reinforcements around the crack tip during the fracture test cannot be inferred directly from the results in Fig. 9(b). The differences in triaxiality between a tensile and a fracture
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P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
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Fig. 9. (a) Mechanical properties of the unreinforced 2014 AI alloy as a function of the temperature [22]. (b) Stresses acting on the A1203 particulates during monotonic tensile deformation of the 2014 AIbased composite at different temperatures. Table 4 Coefficients A and n for the 2014 A1 matrix as a function of the temperature Temperature (°C)
20
100
200
300
A (MPa) n
1825 0.1063
1848 0.1230
1572 0.1259
346 0.0498
specimen, and the high stress gradients around the crack tip, modify the amount of load transferred from the matrix to the reinforcements. However, the changes in the particulate stresses presented in Fig. 9(b) are in perfect qualitative agreement with the levels of reinforcement fracture observed just below the fracture surfaces (Fig. 8(a)). The fraction of broken reinforce-
ments decreased only slightly up to 200 °C, as did the stresses on the reinforcements. On the contrary, the stresses on the reinforcements were reduced by factor of 4-5 between 200 °C and 300 °C and the fraction of broken reinforcements at 300 °C was almost negligible. Thus void nucleation around the crack tip took place at low temperatures by the progressive fracture of the reinforcements, as indicated by in situ observations [4,7,9]. Once a reinforcement has cracked, any additional strain in that region can be accommodated by the opening of an intraparticle void, as shown in Fig. 7(a). At high temperatures, however, the stresses transferred from the matrix to the reinforcements were not high enough to fracture the ceramic particulates. When reinforcement fracture is inhibited, all the strain has to be accommodated by the ductile matrix, and damage is concentrated at the interface or in the matrix near the interface [2,4,5,7,12]. Two regions of the microstructure are particularly susceptible to void nucleation: individual particulate ends and the region between closely spaced particulates aligned in the loading direction (Fig. 7(c)). Finite element analyses have demonstrated that the hydrostatic stresses at maxima in these areas [24], increasing the stresses acting on the reinforcement-matrix interface and thus aiding the development of interfacial debonding [25,26]. In addition, the plastic strains in the matrix are also at maxima in these areas, leading to the nucleation of voids in the matrix from dispersoids and other submicron-sized particles. These microvoids grow and coalesce quickly under the influence of the tensile hydrostatic stress fields, leading to failure in the matrix near the interface [27]. The preponderant damage mechanism (failure either at the interface or in the matrix near the interface) depends on complex microstructural factors, such as the interfacial strength and the fraction of void nucleating particles in the matrix. For instance, Lee et al. [11] found a change from void nucleating particles in the matrix in two powder metallurgy A1-Cu-Mg alloys reinforced with SiC whiskers as the volume fraction of Fe- and Mn-containing dispersoids increased. Similarly, Somerday et al. [12] reported a transition in the damage initiation sites from the interface to within the matrix as the temperature increased, although the precise void nucleation sites in the matrix were not determined. In our materials, void nucleation at the interface (Fig. 7(c)) and in the matrix around the ceramic reinforcements (Fig. 6) coexisted at elevated temperature. Finally, it is interesting to study whether the degradation of the fracture toughness above 200 °C was mainly due to changes in the microstructure (as occurs after overaging [4,6]) or to the reduction in the mechanical properties with temperature. From the continuum mechanics viewpoint, Rice and Johnson [28] proposed that crack extension proceeds when the extent of the heavily deformed region ahead of the crack tip was comparable
191
P. Poza, J. Llorca /Materials Science and Engineering A206 (1996) 183 193
with a critical microstructural distance 2~. The length of this region was equal to the critical crack tip opening displacement gc at crack propagation which for smallscale yielding conditions and plane strain can be calculated as
~c ~
0.5/(L
= 2~
{2)
E6y
(a)
=--~-'-AI 6061 + 15% AI203
500
Q--
--"
3
400
¢-
where E and ay denote for the Young's modulus and the yield strength of the composite respectively. Hahn and Rosenfield [29] assumed that 2~ could be taken as the width of the unbroken matrix ligaments separating cracked particles and applied this model successfully to a number of A1 alloys. The Hahn and Rosenfield model was subsequently applied to particulate- and whiskerreinforced MMCs by considering that 2c was equal to the interparticulate distance [3,5,12,30]. The available experimental results have been reviewed recently in the light of this approach [30], and they reveal a general trend of increasing 2~ with the interparticulate distance, although the exact kind of relationship was not clear. This was to be expected because the fracture processes depend on several factors which can be very different among similar composites. Specifically, not all the reinforcements around the crack tip nucleate voids (Fig. 8) and changes in the reinforcement size and spatial distribution can modify significantly the actual values of interparticulate distance [31]. Similarly, the volume fraction and size of second phase and constituent particles in the matrix, which control the final coalescence of voids in the composites, can lead to very different values of 2~ for a given reinforcement distribution [l 1]. Regardless of the model limitations, it is sensible to assume that the magnitude of 2: should remain constant over the temperature range studied if the changes in K~. were due to mechanical effects. To calculate 2~, from Eq. (2), it is necessary to know the influence of the temperature on ay and E for the composites. The variation in r~y for both composites in the range 20 °C-370 °C after 30 min of exposure was reported by the manufacturer [21] and is shown in Fig. 10(a). The modulus dependence with temperature was studied by Lloyd et al. [32]. They concluded that, for composites in a stable temper, the Young's modulus was given by a linear relationship of the form E = E o - r e ( T = To)
600
(3)
where E0 was the modulus at ambient temperature To, which was measured previously for each composite [18]. The coefficient m was found to be close to 0.06 for several Al-based MMCs [32]. The values of 2c as a function of the temperature are shown in Fig. 10(b) for both composites. They are far from the theoretical value of 18.2gm, obtained by assuming that 2: was equal to the interparticulate distance and that all the reinforcements were spherical and
¢t:~
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200
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Temperature (°C) 20
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300
(°C)
Fig. 10. (a) Influenceof the temperature (30 min soak time) on the yield strength of the composites [21]. (b) Evolution of the critical microstructural distance 2~ as a function of the temperature for both composites. homogeneously distributed. In addition, the magnitude of 2~ was consistently lower for the 6061 A1 based composite, probably because of the smaller fraction of inclusions and void nucleating particles in the matrix. However, 2~ remained fairly constant over the temperature range studied, especially if it is taken into account that the actual values of 2c at 300 °C should be slightly lower than those shown in Fig. 10(b) because of the partial plane stress conditions at the crack tip. Thus mechanical effects seemed to be the dominant factors controlling the changes in the composite toughness at elevated temperatures. This conclusion is in agreement with other observations. For instance, the total fraction
192
P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
of void nucleating particulates was independent of the temperature for each material (although the nucleation mechanism was different) and no significant changes in the microstructure were observed in the scanning electron microscope after high temperature exposure. It should be noted here that the increment in fracture toughness with the yield strength, predicted by the Rice and Johnson model, is contrary to many experimental observations in high strength AI alloys and composites. For instance, artificial aging of A1 alloys leads to higher strength and lower toughness until the peak-aged condition is attained. Further aging beyond this point reduced the alloy strength while the fracture toughness increased [29]. Data for the composites are less conclusive, although the increment is strength caused by artificial aging also produced a marked reduction in toughness [3,4,6,13]. In general, the differences between the model predictions and the experimental results were attributed to the development and coarsening of the precipitates during aging, which modified the magnitude of 2c and the mechanisms of void coalescence [6,29]. In our case, the latter factors could only be important in reducing the fracture toughness at 300 °C, where significant overaging is likely to have occurred after 20 rain of exposure at this temperature.
6. Conclusions The fracture toughness of two AI alloys reinforced with 15 vol.% A1203 particulates were measured in the temperature range from - 150 °C to 300 °C. Both composites maintained acceptable levels of fracture toughness up to 190°C. On the contrary, the fracture toughness decreased by 30%-40% from 190°C to 300 °C. The fracture processes ahead of the crack tip involved the nucleation of voids from the ceramic particulates, the progressive growth of these voids, and the final coalescence through the matrix. During the final coalescence, microvoids were also nucleated from inclusions and dispersoids present in the matrix. The main difference in the fracture processes between low and high temperatures was the void nucleation mechanism, which changed progressively from reinforcement fracture to interfacial decohesion and to failure in the matrix near the interface as the temperature increased. However, the fraction of void nucleating particulates remained fairly constant (around 15%-20%) throughout the whole temperature range studied. The change from reinforcement fracture to interfacial decohesion was explained by the reduction in the reinforcement stresses as a result of matrix softening at elevated temperatures. Finally, the influence of the temperature on the fracture toughness was analysed through the Rice and
Johnson model [28], which assumed that crack propagation took place when the extent of the heavily deformed region ahead of the crack tip was comparable with a critical microstructural distance. It was shown that this distance remained almost constant for the temperature range studied. Thus the degradation of the fracture toughness at elevated temperatures was mainly due to the deterioration of the mechanical properties with temperature rather than to changes in the composite microstructure.
Acknowledgements This research was supported by Tecnologia y Gesti6n de la Innovaci6n SA, through Contract 910-117, and by Comisi6n Interministerial de Ciencia y Tecnologia (CICYT), Spain, under Grant MAT92-32. The help of J.M. Martinez during the mechanical tests is gratefully acknowledged.
References [1] B. Roebuck and J.D. Lord, Mater. Sci. Technol., 6 (1990) 1199 - 1209. [2] S.B. Wu and R.J. Arsenault, Mater. Sci. Eng. ,4, 138 (1991) 227-235. [3] S. Kamat, J.P. Hirt and R. Mehrabian, Acta Metall., 37 (1989) 2395-2402. [4] J.J. Lewandowski, C. Liu and W.H. Hunt, Jr., Mater. Sci. Eng. ,4, 107 (1989) 241-255. [5] M. Manoharan and J.J. Lewandowski, Acta Metall. Mater., 38 (1990) 489-496. [6] T.F. Klimowicz and K.S. Vecchio, in P.K. Liaw and M.N. Gungor (eds.), Fundamental Relationships Between Microstructure & Mechanical Properties o f Metal-Matrix Composites, The Mineral, Metals and Materials Society, Warrendale, PA, 1989, pp. 255-267. [7] D.L. Davidson, Metall. Trans. A, 22 (1991) 113-123. [8] T.J. Downes and J.E. King, in N. Hansen et al. (eds.), Metal Matrix Composites--Processing, Microstructure and Properties, Riso National Laboratory, Roskilde, 1991, pp. 305-310. [9] M. Manoharan, J.J. Lewandowski and W.H. Hunt, Jr., Mater. Sci. Eng. A, 172 (1993) 63-69. [10] M.J. Hadianfard, G. Heness, J.C. Healy and Y.-W. Mai, Fatigue Fract. Eng. Mater. Struct., 17 (1994) 253-263. [11] S. Lee, T.H. Kim and D. Kwon, MetalL Mater. Trans. ,4, 25 (1994) 2213-2223. [12] B.P. Somerday, Y. Leng, F.E. wawner and R.P. Gangloff, in M.N. Gungor (ed.), Advanced Metal Matrix Composites ,for Elevated Temperatures, ASM International, Cincinnati, OH, 1991, pp. 167 182. [13] J. Llorca and P. Poza, Scr. Metall. Mater., 29 (1993) 261-266. [14] M.J. Hadianfard, J.C. Healy and Y.-W. Mai, ,4ppl. Compos. Mater., 1 (1994) 93-113. [15] G. Gonzfilez-Doncel and O.D. Sherby, ,4cta Metall. Mater., 41 (1993) 2797-2805. [16] A. Martin, M.A. Martinez and J. Llorca, Scr. Metall. Mater., 28 (1993) 207-212. [17] R.P. Reed and A.F. Clark (eds.), Materials at Low Temperatures, ASM International, Cincinnati, OH, 1983.
P. Poza, J. Llorca / Materials" Seienee and Engineering A206 (1996) 183 193 [18] J. Llorca and P. Poza, Mater. Sci. Eng. A, 185 (1994) 25 37. [19] J.Y. Pastor, J. Llorca, J. Planas and M. Elices, J. Eng. Mater. Technol., 115 (1993) 281-285. [20] D.F. Carroll, S,M. Wiederhorn and D.E. Roberts, J. Am. Ceram. Soe., 72 (1989) 1610-1614. [21] Duralcan, Mechanical and Physical Property Data for Wrought Composites, San Diego, CA, 1991. [22] D. Zhao, F.R. Tuler and D.J. Lloyd, Acta Metall. Mater., 42 (1994) 2525 2533. [23] I. Justice, J.L. Martinez, P. Poza and J. Llorca, Metall. Mater. Trans. A, 26 (1995) in press. [24] Z. Wang, T.-K. Cheng and D.J. Lloyd, Metall. Trans. A, 24 (1993) 197-207. [25] S.R. Nutt and A. Needleman, Scr. Metall., 21 (1987) 705 710. [26] V. Tvergaard, Mater. Sci. Eng. A, 125 (1990)203-213.
193
[27] J. Llorca, A. Needleman and S. Suresh, Aeta Metall. Mater., 39 (1991) 2317 2335. [28] J.R. Rice and M.A. Johnson, in M.F. Kanninen et al. (eds.), Inelastic Behavior of Solids, McGraw-Hill, New York, 1970, pp. 641-672. [29] G.T. Hahn and A.R. Rosenfield, Metall. Trans. A, 6 (1975) 653 670. [30] J.J. Lewandowski and P.M. Singh, in J.J. Lewandowski and W.H. Hunt, Jr., (eds.), Intrinsic and Extrinsic Fracture Mechanisms in Inorganic Composite Systems, The Mineral, Metals and Materials Society, Warrendale, PA, 1995, pp. 129 146. [31] S. Tao and D.J. Boyd, in R.B. Bhagat, S.G. Fishman and R.J. Arsenault (eds.), Mechanisms and Mechanics ~71 Composite Fracutre, ASM International, Cincinnati, OH, 1993, pp. 29 40. [32] D.J. Lloyd, H. Burke and B. Farahbakhsh, Mater. Sci. Teehnol., 10 (1994) 257 262.
Materials Science and Engineering A206 (19961 183 193
A
Fracture toughness and fracture mechanisms of A1-A1203 composites at cryogenic and elevated temperatures P. Poza, J. Llorca Department q[' Materials Science, Polytechnic Universitj' of Madrid, E.T.S. de lngenieros de Caminos, 28040 Madrid, Spain Received 31 May 1995: in revised form 13 July 1995
Abstract
The fracture toughness of two Al alloys (2014 and 6061) reinforced with 15 vol.% AI20 3 particulates was measured in the temperature range from - 150 °C to 300 °C. Both composites maintained acceptable levels of toughness (over 20 MPa m ~';2)up to 200 °C, but the fracture toughness dropped very quickly above this temperature. Fracture always took place by a ductile mechanism involving the nucleation of voids from the ceramic particulates, the progressive growth of these voids, and the final coalescence through the matrix. Quantitative microscopy analyses demonstrated that voids were initiated from reinforcement fracture at low temperatures and by failure at the interface or in the matrix near the interface at elevated temperatures. Regardless of these differences, the changes in toughness could be explained by the deterioration in the mechanical properties with temperature rather than by the microstructural changes induced by the exposure to elevated temperatures. Keywords: Fracture; Composites; High temperature; Cryogenic temperatures; Toughness
1. Introduction
The low fracture toughness of discontinuously reinforced metal-matrix composites (MMCs) has been one of the main drawbacks to their industrial application. The plane strain fracture toughness of Al-based composites is usually in the range 10-25 M P a m ~/2 [1] and this factor, together with their reduced tensile ductility, has precluded their use in structural elements in the aerospace industry. As a result, the fracture behaviour of particulate- and whisker-reinforced M M C s has received considerable attention from the research community in quest of the mechanisms responsible of this brittle behaviour [2-11]. The investigations of the fracture toughness and the fracture mechanisms in M M C have been focused on the ambient temperature behaviour, and in fact data in the fracture toughness above 100°C are scarce [12-14]. Somerday et al. [12] measured the fracture resistance of a 2009 AI alloy reinforced with 20 vol.% SiC particulates in the temperature range 25-316 °C. They found that the fracture toughness increased slightly from 20 MPa m 12 at ambient temperature to a m a x i m u m of 29 MPa m ~'2 at 220°C. Beyond this temperature, the 0921-5093/96/$15.00 ,~2 1996 S S D I 0921-5093(95)09999-9
Elsevier Science S.A. All rights reserved
composite toughness decreased to a minimum of 14 M P a m L:2 at 316 °C. In addition, a transition from unstable to stable crack propagation was found at 175 °C. Fractographic observations indicated that crack intiation and propagation took place by microvoid nucleation, growth and coalescence in the matrix at all temperatures. The preferred void initiation site changed from the S i C - m a t r i x interface at low temperatures to within the matrix above 200 °C. A micromechanical model based on a critical strain concept was used to predict the changes in the initiation toughness and tearing modulus with the temperature, with the local fracture strain being inferred from the tensile data. In another investigation, we also measured the fracture toughness of another A I - S i C composite in the temperature range from - 1 3 6 °C to 190 °C [13]. The fracture toughness of this material was fairly constant (between 17 and 2 0 M P a m ~:2) in this temperature range and crack propagation was unstable in all cases. The mechanism of void initiation in this composite was reinforcement fracture, regardless of the temperature and aging condition, and the volume fraction of broken reinforcements ahead of the crack tip was between 10% and 15%.
P, Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
184
These preliminary results seem to indicate that the fracture properties of discontinuously reinforced MMCs are maintained up to 200 °C, but it is evident that more research is needed to ascertain the influence of the temperature on the fracture toughness and fracture mechanisms in these materials. This need is more critical when it is considered that MMCs are likely to be used in high temperature applications owing to their excellent creep [15] and wear [16] resistance. In addition, precipitation- and solution-hardened A1 alloys are currently used for the storage and handling of cryogenic liquids (such as nitrogen, oxygen, natural gas, and carbon dioxide) and in other low temperature applications owing to their non-magnetic behaviour, stable microstructure, and retention of strength, ductility and toughness at very low temperatures [17]. Potential savings are also envisaged in the use of A1-SiC an AIA1203 composites for cryogenic applications but it is necessary to determine the mechanical properties of these materials under such conditions. This investigation was aimed at determining the effect of the temperature (in the range from - 150 °C to 300 °C) on the fracture resistance and the fracture mechanisms in two precipitation-hardened A1 alloys reinforced with A1203 particulates. The results of the fracture tests, together with detailed analyses of the failure mechanisms by means of quantitative microscopy, were used to elucidate the relationships between the macroscopic behaviour and the microstructural features.
2. M a t e r i a l s
Two composite materials were used in this investigation: a 2014 A1 alloy reinforced with 14.4 vol.% A1203 particulates, and a 6061 A1 alloy reinforced with 14.9 vol.% A1203 particulates. The cast materials w e r e supplied by the manufacturer (Duralcan, San Diego, CA, USA) in the form of extruded circular bars 50 mm in diameter in the peak-aged condition (T6). The heat treatments to achieve this temper in the 2014 A1 based composite included a solution heat treatment at 502 °C for 2 h, water quenching, and aging at 160 °C for 16 h. The heat treatments in the 6061 A1 based material w e r e the same, although the solution heat treatment temperature was 560 °C. The chemical compositions of the
Table 2 Geometric characteristics of the AI203 particulates in the longitudinal direction Composite
AI 2014 + AI203 AI 6061 + A1203
Areaa (I.tm2)
Aspect ratio a
145 + 81 144 + 70
0.44 + 0.16 0.51 + 0.19
Dmin/Dmax
a+ standard deviation.
matrices in both composites were provided by the manufacturer and are given in Table 1. The microstructural characteristics of both composites w e r e detailed in a previous publication [18]. The sizes and morphologies of the ceramic particulates were very similar in the two materials, and the average particulate area and aspect ratio are given in Table 2. The particulates were oriented with their longer axis parallel to the extrusion direction (Fig. 1), and the aspect ratio is approximately the quotient between the particulate width (perpendicular to the extrusion axis) and length (parallel to the extrusion axis). The area and aspect ratio distributions were wide, as indicated by the standard deviations in Table 2. The area of 80% of the particulates was in the range 60-225 gm 2, although reinforcements of up to 500 ~tm2 were occasionally found. It is worth noting that broken A1203 particulates w e r e found in both composites prior to testing. Most of the broken reinforcements were observed in regions w h e r e the ceramic particulates were very close together in a cluster and they were shattered (Fig. 2) rather than broken with a single crack. Porosity (black regions in Fig. 2) was often associated with clusters of shattered particulates. This fragmentation of the ceramic particulates probably took place during extrusion in regions w h e r e the local volume fraction of reinforcement was high. In addition to shattered reinforcements, some isolated particulates which were not associated with a cluster were also broken. The fracture mode was differ-
Table 1 Chemical composition (weight per cent) of the matrices Si
Cu
Mg
Mn
Fe
Cr
Zn
Ti
AI
2014 A1 0.76 4.70 0.47 0.78 0.06 0,01 0,06 0.03 Remainder 6061 AI 0.62 0.24 0.94 0.04 0.07 0.10 0.02 0.01 Remainder
Fig. 1. Low magnification micrograph of the 2014 A1 based composite. Notice that the AI203 particulates are oriented with their longer axis in the extrusion direction (vertical).
P. Poza, J, LIorca / Materials' Science and Engineering A206 (1996) 18.7 193
Fig. 2. Backscattered electron micrograph of a cluster of shattered particulates in the 2014 A1 alloy reinforced with 15v01.% AI203 particulates. AI203 particulates are seen as dark grey zones, whereas AI C u - M n inclusions are white and the black areas are pores. Extrusion axis is vertical.
ent in these cases, the reinforcement being fractured by a single crack perpendicular to the extrusion direction. The fraction of isolated particulates was, however, much less important than damage associated with clusters, as indicated in Table 3, where the percentage of isolated and clustered broken AI203 particulates is given for each composite. Finally, it should be indicated that the ceramic reinforcements were well bonded to the A1 matrix, and interfacial decohesion was rarely observed. In addition to the A1203 particulates, inclusions of irregular shape and about 1 jam in diameter were distributed within the matrix of the 2014 A1 based composites (Fig. 2). They appeared to be intermetallic inclusions, whose main constituents (determined by energy dispersive X-ray microanalysis) were A1, Cu and Mn. The study by means of backscattered electrons of the 6061 A1 based material indicated that the volume fraction of inclusions in this material was very small, and areas with a chemical composition different from the matrix or the particulates were scarce. Small dispersoids (around 0.1 jam) were found in the matrix, and they were identified as Mn- and Cr-rich compounds, which inhibit recrystallization during cooling.
3. Experimental techniques Compact tension specimens were machined from the bars for the fracture tests. Specimen thickness was 18 ram, the crack being oriented in the L-T direction. Specimens were precracked in fatigue, in accordance with the recommendations of ASTM Standard E 399, in a servohydraulic mechanical testing machine. To reduce the number of fatigue cycles to initiate the fatigue crack, a through-thickness cut 1 mm in depth was given at the notch root with a very thin (0.3 ram)
185
diamond saw. Fracture tests were performed in stroke control, and the displacement rate was set to obtain a stress intensity factor rate of 0.6 M P a m 1 2 s ~. Load and crack mouth opening displacement (CMOD) were continuously monitored through a computer-controlled data acquisition system. For high and low temperature tests, the specimen was placed in a temperature chamber. Heating was by an electric resistance and liquid nitrogen was used for cooling. The specimen temperature was measured with a K-type thermocouple, which was introduced in a 2 mm deep hole machined in the specimen. Heating (or cooling) rate was always slower than 5 °C rain -~. The specimen was held for 20 min at the test temperature (_+ 2 °C) prior to testing, to ensure a homogeneous temperature distribution through the thickness. The C M O D was measured with a standard clip displacement gauge in the tests carried out below 200°C. As this gauge could not work beyond this temperature, a laser extensometer developed to measure the C M O D at high temperatures was used [19]. This extensometer is essentially similar to that reported by Carroll et al. [20]. An He Ne laser emitter sends a scanning laser beam through two silica windows across the furnace walls. Two alumina pins glued on the points which determine the C M O D are interposed in the laser beam and the resulting intensity distribution is measured by a detector placed on the opposite side of the furnace. A microcomputer processes the intensity distribution to determine the distance between the pins. Although the resolution of the system is very high, the accuracy is limited mainly by two factors: the air seepage around the loading rods and the thermal gradients along the optical paths. Preliminary calibrations of the extensometer indicated that air seepage was dominant as compared with the thermal gradients in our system, perhaps because the large size of the furnace led to smaller temperature gradients along the laser path. Thus special care was taken to reduce the chimney effect as much as possible by closing the gaps between the loading rods and the furnace. In addition, the specimen remained at the test temperature for 20 min before testing to achieve a stationary distribution of temperatures within the furnace. Moreover, the good thermal insulation of the furnace allowed the laser source and detector to be located very close, to the silica glass without danger of overheating, reducing the fluctuations produced outside the furnace. In addition to all these precautions, the measurements were averaged over 150 scans for each displacement to minimize the noise. Metallographic samples for both materials were prepared in the longitudinal (extrusion) direction from the specimens after testing to analyse the failure micromechanisms. The specimens were cut through the middle with a low speed diamond blade wheel to minimize
P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
186
damage, using oil as a lubricant, and initially polished on SiC abrasive paper to 600 grit finish. This was followed by polishing on diamond slurry (9, 3 and 1 ~tm) and finally on magnesia. The polished samples were first cleaned in deionized water and afterwards by ultrasound in acetone. Scanning electron microscope micrographs were taken randomly from the polished surfaces at less than 100 tam from the fracture surface. This distance was always less than the size of the plastic zone around the crack tip for all the materials and temperatures. These micrographs were processed by means of a computer-controlled image analysis system to determine the fraction of void nucleating particles in the composite and the origin of the voids (either particulae fracture or matrix-reinforcement decohesion). In addition, the fracture surfaces were examined by scanning electron microscopy to ascertain the main failure mechanisms for each material and temperature.
30
. . . .
,
'
'
'
'
,
. . . .
,
. . . .
,
,
,
,
,
I
,
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25
ffl e.1:: 2 o O b-
~
15
It.
•
-200
-100
AI 2014 + 1 5 * AI203
0
100
200
300
Temperature (°C) Fig. 3. Fracture toughness of the composite materials as a function of the test temperature.
4. R e s u l t s
4.1. Fracture toughness The results of the fracture tests were examined according to the procedures indicated in the ASTM Standard E-399. All the tests fulfilled the conditions prescribed for fatigue precracking, crack propagation during fatigue, crack front curvature, and ratio between the load for crack initiation and the maximum load during the test. The specimen thickness also met the requirements indicated in the standard for all the tests except those carried out at 300 °C. Given the values of the yield strength reported for these composites at 300 °C by the manufacturer [21], the specimen thickness should have been 26 mm for the 2014 A1 based composite and 38mm for the 6061 A1 based material. (According to Ref. [21], the yield strength of the composites at different temperatures was measured after 30min soak time. The actual values of the yield strength as a function of the temperature are plotted in Fig. 10(a)). Thus the results presented in Fig. 3 represent the plane strain fracture toughness of the composites up to 190 °C. The data obtained at 300 °C cannot be regarded as Ktc values, properly speaking, because plane strain conditions did not prevail at the crack tip. In fact, while the fracture surfaces up to
190 °C were flat, shear lips oriented on +45 ° bands to the specimen thickness were found at 300 °C. These shear lips, indicative of plane stress conditions at the crack tip, covered around 35% of the fracture surface in the 2014 A1 based composite and around 55% in the 6061 A1 based material. Even so, it is evident that both composites maintained acceptable levels of fracture toughness in the temperature range from - 1 5 0 °C to 190 °C. On the contrary, the fracture toughness decreased by 30%40% from 190 °C to 300 °C. This result is more dramatic when it is considered that the actual values of the plane strain fracture toughness at 300 °C should be lower than those obtained from our fracture tests. It should be noted that the reduction in the resistance to crack initiation was accompanied by an increase in the resistance to crack propagation, as is shown by the load-CMOD curves presented in Fig. 4 for the 6061 A1 based composite. The crack began to propagate unstably near to the maximum load at low temperatures, while a substantial amount of stable crack growth was observed at 300°C. The tests carried out at 190°C revealed an intermediate behaviour. Similar results were observed in the other composite and were reported by other researchers [12].
Table 3 Fraction of broken A1203 particulates in the as-received materials
4.2. Fracture mechanL~ms
Composite
AI 2014 + A1203 A1 6061 + A1203
Isolated particulates
Clustered particulates
5.5 2.7
8.5 13.0
(%)
(%)
The analysis of the fracture mechanisms began with the study of the fracture surfaces in the scanning electron microscope. All the fracture surfaces presented the typical morphology of failure by a ductile mechanism involving the nucleation, growth and coalescence of
P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183 193
voids (Figs. 5 and 6). At low temperatures (between - 1 5 0 °C and 20 °C), the principal feature found in both materials was the presence of a large number of A1203 particulates on the fracture surfaces (Fig. 5). The ceramic particulates were clean and no traces of the A1 matrix were observed on them. The A1 matrix around the reinforcements was heavily deformed and another microvoid population was observed in the matrix. Microvoids were formed around the A 1 - C u - M n inclusions and smaller dispersoids and precipitates in the 2014 A1 based composite, while the Mn- and Cr-rich dispersoids were identified as microvoid nucleation sites in the 6061 A1 based material. It should also be noted that it was not possible to distinguish the fracture surfaces created at ambient temperature from those formed at - 150 °C. The fracture surfaces of the specimens tested at 300 °C also presented the dual population of voids formed around the ceramic reinforcements and the inclusions and dispersoids in the matrix. However, the majority of the A1203 particulates (whose size and morphology were clearly visible on the fracture surface) were covered with AI matrix (Fig. 6), while they were free from matrix in the fracture surfaces created at ambient and cryogenic temperatures. The fracture surfaces at 190 °C (not shown here) exhibited an intermediate behaviour where ceramic reinforcements both clean and covered with matrix could be seen. It should be noted that the average void sizes were very similar at all temperatures. The previous observations seemed to indicate a change in the void nucleation sites from the reinforce-
AI 6061 + 15 % AI203 10--,
,
,
,
~
,-,
,
,
,
. . . .
(°C)
Temperature -150
8 ---=---
Z
6
o .._1
4
0
0.5
1
190
1.5
CMOD (mm) Fig. 4. Load C M O D curves for the 6061 AI composite reinforced with AI~O 3 particulates at different temperatures.
187
Fig. 5. Ambient temperature fracture surfaces. (a) 2014 AI alloy reinforced with AI203 particulates. (b) 6061 AI alloy reinforced with AI203 particulates. The ceramic reinforcements on the fracture surfaces were free of matrix.
ment at low temperatures to those near the interface at elevated temperatures. To check this hypothesis, the specimens tested at - 1 5 0 °C, 190 °C and 300 °C were sliced through the middle and polished, and the region around the crack path was observed in the scanning electron microscope. Three representative micrographs of the mechanisms seen at the three temperatures are shown in Fig. 7. The dominant void initiation mechanism at - 1 5 0 °C was reinforcement fracture by one crack oriented perpendicularly to the loading direction, as shown by the arrows in Fig. 7(a). Afterwards, the voids grew in the longitudinal but not in the radial direction. These observations are in agreement with the morphology of the fracture surfaces, where the most typical feature was a matrix void containing one clean A1203 particulate, the void diameter and shape being similar to those of the reinforcement. At 300 °C, on the contrary, interracial decohesion became the main void nucleation mechanism (Fig. 7(c)). Matrix-reinforcement decohesion usually took place at the end of elongated reinforcements or in regions where the ends of two reinforcements were very close in the longitudinal direction. The specimens tested at 190 °C presented both kinds of failure mechanisms, and rein-
188
P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
forcement fracture and matrix-reinforcement decohesion were observed around the fracture surfaces (Fig. 7(b)). To quantify the importance o f each mechanism, the micrographs were processed by means of a computer-controlled image analysis system to determine the fraction of void nucleating particulates in the composite and the origin of the voids (either particulate fracture or matrix-reinforcement decohesion). This study did not include the ceramic particulates shattered in clusters, because they were already broken prior to testing.
Fig. 7. Void nucleation mechanisms as a function of the test temperature. (a) 6061 AI based composite tested at -150 °C. (b) 2014 AI based material tested at 190 °C. (c) 6061 AI based composite tested at 300 °C.
Fig. 6. Fracture surfaces created at 300 °C. (a) 2014 AI alloy reinforced with AI203 particulates. (b), (c) 6061 A1 alloy reinforced with AI203 particulates. Matrix voids were observed on the ceramic reinforcements, which are easily recognized by their size and shape.
The results of this analysis for both composites are shown in Fig. 8. It shows that only a small fraction (between 15% and 20%) of the ceramic particulates acted as void nucleation sites, regardless of the matrix alloy and test temperature. In addition, the void nucleation mechanism changed from reinforcement fracture at - 1 9 0 °C to interfacial decohesion at 300 °C. Ap-
P. Poza, J, Llorca / Materials Science and Engineering A 206 (1996) 183-193
proximately the same numbers of voids were nucleated by reinforcement fracture and by interfacial decohesion at 190 °C. It is worth noting that the fracture of broken reinforcements at 300 °C was very similar to the fraction of isolated A120~ particulates fractured before testing (Table 3). Thus reinforcement fracture had practically disappeared at this temperature.
5. Discussion The experimental observations showed that void nucleation by reinforcement fracture was progressively substituted by failure at or near the reinforcement-matrix interface as the temperature increased. This effect was also reported by Zhao et al. [22] who analysed the influence of the temperature on the tensile behaviour of a 2014 AI alloy reinforced with AI203 particulates, In agreement with our results, reinforcement fracture was dominant at ambient temperature while interfacial decohesion and matrix voiding were preponderant at 300 °C. Around 200 °C, they found that the number of voids was evenly divided between both mechanisms, in perfect agreement with the results plotted in Fig. 8. The reduction in the number of broken reinforcements as a function of the temperature can be explained through the changes in the matrix mechanical properties. The stresses on the reinforcements during monotonic deformation were recently calculated by Justice et al. [23] who carried out a parametric study of the influence of the matrix properties on the reinforcement stresses by means of the finite element analysis of a unit cell, representative of the composite. The ceramic reinforcements were located at the centre of each cell. Unit cylinders (the cylinder length was equal to the diameter) were used to model reinforcements with irregular shape, sharp corners and low aspect ratios, such as those used in particulate-reinforced Al-based composites. For a given reinforcement volume fraction and shape, the numerical results were fitted to polynomial expressions of the matrix properties. The tensile properties of the 2014 unreinforced alloy in the peak-aged condition, processed in the same way, were obtained by Zhao et al. [22] up to 300 °C after 30 min of exposure and they are shown in Fig. 9(a). The matrix flow stress o-m can be expressed by a power law of the plastic strain ~ as am=
Ag "
(1)
where the coefficients A and n, which characterize the matrix stress strain behaviour, can be obtained from the results in Fig. 9(a) and they are shown in Table 4. The stresses acting on the ceramic particulates as a function of the plastic strain can then be calculated by introducing these values of A and n in the polynomial expressions reported in Ref. [23]. The results are shown in Fig. 9(b),
189
where the stresses on the reinforcements are plotted as a function of the composite plastic strain. They clearly show that the stresses on the reinforcements remain almost constant up to 100 °C, decrease slightly at 200 °C and drop to very low levels at 300 °C as a consequence of the reduction in the matrix flow stress. The evolution of the stresses on the reinforcements around the crack tip during the fracture test cannot be inferred directly from the results in Fig. 9(b). The differences in triaxiality between a tensile and a fracture
2014 AI + 15% AI203 -'--25
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200
300
T e m p e r a t u r e (°C)
6061 AI + 15% AI203 ~
o~
25
....
, ....
, ....
~ ....
, ....
,,
(b)
v
20 (.)
(:~ c
15
m o E
10
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5
o
U-
n0 -200
'
-100
0
100
200
300
Temporature (°C) Fig. 8. Fraction of void nucleating reinforcements as a function of the test temperature. (a) 2014 AI alloy reinforced with A1203 particulates. (b) 6061 AI alloy reinforced with AI203 particulates.
P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
190
500
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i
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i
'
i
'
i
,
i
,
i
(a)
0.3
400 0.2 o_ 300
O9
v
o') U)
2 200 O9
0.1
100
Ultimate strength Elongation
•
0
0
50
100
200
150
250
300
Temperature (°C) 1500
'
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I
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'
'
I
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'
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I
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v
1000 (j,)
E E
q) o o ._ n-"
500
(1)
0 0
0.02
0.04
0.06
0.08
0.1
Plastic strain
Fig. 9. (a) Mechanical properties of the unreinforced 2014 AI alloy as a function of the temperature [22]. (b) Stresses acting on the A1203 particulates during monotonic tensile deformation of the 2014 AIbased composite at different temperatures. Table 4 Coefficients A and n for the 2014 A1 matrix as a function of the temperature Temperature (°C)
20
100
200
300
A (MPa) n
1825 0.1063
1848 0.1230
1572 0.1259
346 0.0498
specimen, and the high stress gradients around the crack tip, modify the amount of load transferred from the matrix to the reinforcements. However, the changes in the particulate stresses presented in Fig. 9(b) are in perfect qualitative agreement with the levels of reinforcement fracture observed just below the fracture surfaces (Fig. 8(a)). The fraction of broken reinforce-
ments decreased only slightly up to 200 °C, as did the stresses on the reinforcements. On the contrary, the stresses on the reinforcements were reduced by factor of 4-5 between 200 °C and 300 °C and the fraction of broken reinforcements at 300 °C was almost negligible. Thus void nucleation around the crack tip took place at low temperatures by the progressive fracture of the reinforcements, as indicated by in situ observations [4,7,9]. Once a reinforcement has cracked, any additional strain in that region can be accommodated by the opening of an intraparticle void, as shown in Fig. 7(a). At high temperatures, however, the stresses transferred from the matrix to the reinforcements were not high enough to fracture the ceramic particulates. When reinforcement fracture is inhibited, all the strain has to be accommodated by the ductile matrix, and damage is concentrated at the interface or in the matrix near the interface [2,4,5,7,12]. Two regions of the microstructure are particularly susceptible to void nucleation: individual particulate ends and the region between closely spaced particulates aligned in the loading direction (Fig. 7(c)). Finite element analyses have demonstrated that the hydrostatic stresses at maxima in these areas [24], increasing the stresses acting on the reinforcement-matrix interface and thus aiding the development of interfacial debonding [25,26]. In addition, the plastic strains in the matrix are also at maxima in these areas, leading to the nucleation of voids in the matrix from dispersoids and other submicron-sized particles. These microvoids grow and coalesce quickly under the influence of the tensile hydrostatic stress fields, leading to failure in the matrix near the interface [27]. The preponderant damage mechanism (failure either at the interface or in the matrix near the interface) depends on complex microstructural factors, such as the interfacial strength and the fraction of void nucleating particles in the matrix. For instance, Lee et al. [11] found a change from void nucleating particles in the matrix in two powder metallurgy A1-Cu-Mg alloys reinforced with SiC whiskers as the volume fraction of Fe- and Mn-containing dispersoids increased. Similarly, Somerday et al. [12] reported a transition in the damage initiation sites from the interface to within the matrix as the temperature increased, although the precise void nucleation sites in the matrix were not determined. In our materials, void nucleation at the interface (Fig. 7(c)) and in the matrix around the ceramic reinforcements (Fig. 6) coexisted at elevated temperature. Finally, it is interesting to study whether the degradation of the fracture toughness above 200 °C was mainly due to changes in the microstructure (as occurs after overaging [4,6]) or to the reduction in the mechanical properties with temperature. From the continuum mechanics viewpoint, Rice and Johnson [28] proposed that crack extension proceeds when the extent of the heavily deformed region ahead of the crack tip was comparable
191
P. Poza, J. Llorca /Materials Science and Engineering A206 (1996) 183 193
with a critical microstructural distance 2~. The length of this region was equal to the critical crack tip opening displacement gc at crack propagation which for smallscale yielding conditions and plane strain can be calculated as
~c ~
0.5/(L
= 2~
{2)
E6y
(a)
=--~-'-AI 6061 + 15% AI203
500
Q--
--"
3
400
¢-
where E and ay denote for the Young's modulus and the yield strength of the composite respectively. Hahn and Rosenfield [29] assumed that 2~ could be taken as the width of the unbroken matrix ligaments separating cracked particles and applied this model successfully to a number of A1 alloys. The Hahn and Rosenfield model was subsequently applied to particulate- and whiskerreinforced MMCs by considering that 2c was equal to the interparticulate distance [3,5,12,30]. The available experimental results have been reviewed recently in the light of this approach [30], and they reveal a general trend of increasing 2~ with the interparticulate distance, although the exact kind of relationship was not clear. This was to be expected because the fracture processes depend on several factors which can be very different among similar composites. Specifically, not all the reinforcements around the crack tip nucleate voids (Fig. 8) and changes in the reinforcement size and spatial distribution can modify significantly the actual values of interparticulate distance [31]. Similarly, the volume fraction and size of second phase and constituent particles in the matrix, which control the final coalescence of voids in the composites, can lead to very different values of 2~ for a given reinforcement distribution [l 1]. Regardless of the model limitations, it is sensible to assume that the magnitude of 2: should remain constant over the temperature range studied if the changes in K~. were due to mechanical effects. To calculate 2~, from Eq. (2), it is necessary to know the influence of the temperature on ay and E for the composites. The variation in r~y for both composites in the range 20 °C-370 °C after 30 min of exposure was reported by the manufacturer [21] and is shown in Fig. 10(a). The modulus dependence with temperature was studied by Lloyd et al. [32]. They concluded that, for composites in a stable temper, the Young's modulus was given by a linear relationship of the form E = E o - r e ( T = To)
600
(3)
where E0 was the modulus at ambient temperature To, which was measured previously for each composite [18]. The coefficient m was found to be close to 0.06 for several Al-based MMCs [32]. The values of 2c as a function of the temperature are shown in Fig. 10(b) for both composites. They are far from the theoretical value of 18.2gm, obtained by assuming that 2: was equal to the interparticulate distance and that all the reinforcements were spherical and
¢t:~
300
(./)
.__ >-
200
100
0 0
50
I
I
I
I
100
150
200
250
1
I
300
350
400
Temperature (°C) 20
I
I
I
(b)
I--'
I
I
- - o - - A I 6061 + 15% AlaO 3 • AI 2014 + 15% Al=Oz
15 -------O
E :::L
v ...o
10
S
J
0
i
0
I
I
50
100
i
I
150
i
[
200
Temperature
i
I
250
L
I
300
(°C)
Fig. 10. (a) Influenceof the temperature (30 min soak time) on the yield strength of the composites [21]. (b) Evolution of the critical microstructural distance 2~ as a function of the temperature for both composites. homogeneously distributed. In addition, the magnitude of 2~ was consistently lower for the 6061 A1 based composite, probably because of the smaller fraction of inclusions and void nucleating particles in the matrix. However, 2~ remained fairly constant over the temperature range studied, especially if it is taken into account that the actual values of 2c at 300 °C should be slightly lower than those shown in Fig. 10(b) because of the partial plane stress conditions at the crack tip. Thus mechanical effects seemed to be the dominant factors controlling the changes in the composite toughness at elevated temperatures. This conclusion is in agreement with other observations. For instance, the total fraction
192
P. Poza, J. Llorca / Materials Science and Engineering A206 (1996) 183-193
of void nucleating particulates was independent of the temperature for each material (although the nucleation mechanism was different) and no significant changes in the microstructure were observed in the scanning electron microscope after high temperature exposure. It should be noted here that the increment in fracture toughness with the yield strength, predicted by the Rice and Johnson model, is contrary to many experimental observations in high strength AI alloys and composites. For instance, artificial aging of A1 alloys leads to higher strength and lower toughness until the peak-aged condition is attained. Further aging beyond this point reduced the alloy strength while the fracture toughness increased [29]. Data for the composites are less conclusive, although the increment is strength caused by artificial aging also produced a marked reduction in toughness [3,4,6,13]. In general, the differences between the model predictions and the experimental results were attributed to the development and coarsening of the precipitates during aging, which modified the magnitude of 2c and the mechanisms of void coalescence [6,29]. In our case, the latter factors could only be important in reducing the fracture toughness at 300 °C, where significant overaging is likely to have occurred after 20 rain of exposure at this temperature.
6. Conclusions The fracture toughness of two AI alloys reinforced with 15 vol.% A1203 particulates were measured in the temperature range from - 150 °C to 300 °C. Both composites maintained acceptable levels of fracture toughness up to 190°C. On the contrary, the fracture toughness decreased by 30%-40% from 190°C to 300 °C. The fracture processes ahead of the crack tip involved the nucleation of voids from the ceramic particulates, the progressive growth of these voids, and the final coalescence through the matrix. During the final coalescence, microvoids were also nucleated from inclusions and dispersoids present in the matrix. The main difference in the fracture processes between low and high temperatures was the void nucleation mechanism, which changed progressively from reinforcement fracture to interfacial decohesion and to failure in the matrix near the interface as the temperature increased. However, the fraction of void nucleating particulates remained fairly constant (around 15%-20%) throughout the whole temperature range studied. The change from reinforcement fracture to interfacial decohesion was explained by the reduction in the reinforcement stresses as a result of matrix softening at elevated temperatures. Finally, the influence of the temperature on the fracture toughness was analysed through the Rice and
Johnson model [28], which assumed that crack propagation took place when the extent of the heavily deformed region ahead of the crack tip was comparable with a critical microstructural distance. It was shown that this distance remained almost constant for the temperature range studied. Thus the degradation of the fracture toughness at elevated temperatures was mainly due to the deterioration of the mechanical properties with temperature rather than to changes in the composite microstructure.
Acknowledgements This research was supported by Tecnologia y Gesti6n de la Innovaci6n SA, through Contract 910-117, and by Comisi6n Interministerial de Ciencia y Tecnologia (CICYT), Spain, under Grant MAT92-32. The help of J.M. Martinez during the mechanical tests is gratefully acknowledged.
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[27] J. Llorca, A. Needleman and S. Suresh, Aeta Metall. Mater., 39 (1991) 2317 2335. [28] J.R. Rice and M.A. Johnson, in M.F. Kanninen et al. (eds.), Inelastic Behavior of Solids, McGraw-Hill, New York, 1970, pp. 641-672. [29] G.T. Hahn and A.R. Rosenfield, Metall. Trans. A, 6 (1975) 653 670. [30] J.J. Lewandowski and P.M. Singh, in J.J. Lewandowski and W.H. Hunt, Jr., (eds.), Intrinsic and Extrinsic Fracture Mechanisms in Inorganic Composite Systems, The Mineral, Metals and Materials Society, Warrendale, PA, 1995, pp. 129 146. [31] S. Tao and D.J. Boyd, in R.B. Bhagat, S.G. Fishman and R.J. Arsenault (eds.), Mechanisms and Mechanics ~71 Composite Fracutre, ASM International, Cincinnati, OH, 1993, pp. 29 40. [32] D.J. Lloyd, H. Burke and B. Farahbakhsh, Mater. Sci. Teehnol., 10 (1994) 257 262.