Internal stress and microstructure of SiC reinforced aluminium alloy 2014

PII:

Acta mater. Vol. 46, No. 15, pp. 5271±5281, 1998 # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain S1359-6454(98)00212-2 1359-6454/98 $19.00 + 0.00

INTERNAL STRESS AND MICROSTRUCTURE OF SiC REINFORCED ALUMINIUM ALLOY 2014 A. WEILAND{1{, L. HULTMAN2, U. WAHLSTROÈM3, C. PERSSON}1 and T. JOHANNESSON}1 1 Division of Engineering Materials, Department of Mechanical Engineering, LinkoÈping University, S-581 83 LinkoÈping, Sweden, 2Thin Film Physics Division, Department of Physics and Measurement Technology, LinkoÈping University, S-581 83 LinkoÈping, Sweden and 3Industrial Microelectronics Center, Department of Physics and Measurement Technology, LinkoÈping University, S-581 83 LinkoÈping, Sweden

(Received 20 January 1998; accepted 9 June 1998) AbstractÐPrecipitates and interfaces in a SiC particulate reinforced AA 2014 alloy have been characterized by high-resolution transmission electron microscopy and X-ray microanalysis. The kinetics of precipitation and the lattice parameter and microstress variations during ageing of the composite were studied by X-ray di€raction at elevated temperature. The kinetics of precipitation in a composite sample and a stress-free reference powder di€er substantially. Thermal mis®t stresses induced by SiC particulates increase the rate of precipitation by approximately a factor three, but do not in¯uence the type or total amount of precipitates. Due to the di€erent precipitation rates, X-ray di€raction measurements of three-dimensional stresses must be made with caution; discrepancies between the sample and the reference powder compositions may give erroneous results. There is no evidence of interfacial reactions between the particulates and the matrix alloy, or of pore formation or preferential precipitation at the interfaces. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. ReÂsumeÂÐLes preÂcipiteÂs et interfaces d'un alliage AA 2014 renforce par des particules de SiC ont eÂte analyseÂs par microscopie eÂlectronique aÁ transmission sous haute reÂsolution et par microanalyse aux rayons X. Les cineÂtiques de preÂcipitation ainsi que les variations de distance interreÂticulaire et de microcontraintes au cours du vieillissement ont eÂte eÂtudieÂes par di€raction X sous haute tempeÂrature. Les cineÂtiques de preÂcipitation du composite et d'une poudre de reÂfeÂrence, libre de contraintes, di€eÁrent consideÂrablement. Les contraintes d'origine thermique induites par les particules de SiC accroissent la vitesse de preÂcipitation d'un facteur 3, mais n'in¯uencent ni le type ni la quantite totale de preÂcipiteÂs. A cause des di€eÂrentes vitesses de preÂcipitation les mesures par di€raction X des contraintes en 3D doivent eÃtre e€ectueÂes avec preÂcaution. Des di€eÂrences de compositions entre l'eÂchantillon et la poudre de reÂfeÂrence peuvent engendrer des reÂsultats erroneÂs. Aucune observation de reÂaction aÁ l'interface matrice±particule n'a eÂte faite, ni de formation de pore ou de preÂcipitation preÂfeÂrentielle aÁ l'interface. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. ZusammenfassungÐIn einem mit SiC Kurzfaser-verstaÈrkten AA 2014 Aluminium sind Ausscheidungen und Grenz¯aÈchen mit Hilfe von roÈntgenographischer Analyse und Elektronenmikroskopie untersucht worden. Ausscheidungsgeschwindigkeiten und AÊnderungen der Gitterstrukturparameter und der Mikroeigenspannungen bei der Alterung dieser Verbundwerksto€ sind mittels RoÈntgenbeugung bei gehobener Temperatur gemessen worden. Das Ausscheidungsverhalten des Verbundwerksto€es und eines beanspruchungsfreien Bezugspulvers unterscheiden sich grundlegend. Thermische Eigenspannungen, wie sie durch SiC Partikel hervorgerufen werden, verdreifachen die Ausscheidungsgeschwindigkeit nahezu, beein¯ussen aber weder Typ noch Anzahl der Ausscheidungen. Aufgrund verschiedener Ausscheidungsraten muÈssen RoÈntgenbeugungsmessungen von drei-dimensional beanspruchten Proben mit Vorsicht gedeutet werden; Unterschiede zwischen Probe und Bezugspulver koÈnnen zu irrefuÈhrenden Ergebnissen fuÈhren. Es gibt keine Anzeichen fuÈr Grenz¯aÈchenreaktionen zwischen den Partikeln und der Aluminium-Matrix, fuÈr Porenformationen oder vorzeitige Ausscheidungen an den Grenz¯aÈchen. # 1998 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.

{Present address: Epact Technology AB, Teknikringen 8, S-583 30 LinkoÈping, Sweden {To whom all correspondence should be addressed. }Present address: Division of Materials Engineering, Department of Solid Mechanics, Lund University, S-221 00 Lund, Sweden }Present address: Division of Materials Engineering, Department of Solid Mechanics, Lund University, S-221 00 Lund, Sweden

1. INTRODUCTION

A thorough understanding of mechanisms involved in the microstructural changes of aluminium matrix composites due to the presence of a reinforcing phase is especially important since changes in matrix microstructure are not always re¯ected in macroscopic yield strength or ultimate strength,

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although their in¯uence on ductility, fracture resistance, and corrosion properties may be very signi®cant [1, 2]. In composites, it is recognized that the matrix alloy does not necessarily respond to heat treatment in a manner identical to that of the unreinforced alloy, and that consideration should therefore be given to the possible e€ects of the ®bres, particulates, or whiskers on the structure of the heat-treated matrix and its properties [2]. Several workers [2±8] have found a substantial e€ect of reinforcement on, e.g. ageing, precipitation, and dissolution kinetics. It has clearly been shown that the heat-treatment response of a reinforced alloy should be signi®cantly di€erent from that of the unreinforced material. The most commonly proposed reason for di€erent ageing kinetics in metal matrix composites (MMCs) is that they contain high dislocation densities compared with unreinforced alloys, generated to relieve the thermal stresses produced during cooling. High dislocation densities may contribute positively to ageing rates by either enhancing the di€usion rates of elements within the matrix or by providing preferential sites for precipitate nucleation [3, 7]. However, in some cases such high dislocation densities may reduce the age-hardening rates of the MMC by e€ectively removing the quenched-in vacancies which are an important constituent in the early stages of Guinier±Preston (GP) zone formation [9]. In the present work, the response of a spray-cast SiC particulate reinforced AA 2014 alloy to heat treatment has been investigated by transmission electron microscopy and X-ray di€raction measurements, both in situ during heat treatment. The objective has been to characterize the precipitates and the matrix/particulate interface and to clarify the in¯uence of particulates and particulate-induced stresses on the hardening and ageing behaviour.

2. MATERIAL AND METHODS

2.1. Material The composite investigated was the aluminium alloy AA 2014 reinforced with 15 vol.% SiC particulates of the SIKA F600 grade. The composition of the matrix alloy is given in Table 1. The composite was manufactured by Cospray Advanced Aluminium Materials in the U.K. using the spray casting technique developed by Osprey Metals Ltd. The material has been hot-extruded at 4158C with an extrusion ratio of 22.8:1. Before extrusion, the nominal particulate size (median 50% value) is 9.3 2 1.0 mm [10]. The standard deviation is

increased by the extrusion, while the absolute particulate size is just marginally decreased. The aspect ratio of the particulates after extrusion is about 2.2:1 and they are aligned in the extrusion direction, giving an e€ective aspect ratio of 1.9:1 [11]. The micrograph in Fig. 1 shows a section cut parallel to the extrusion which has been etched with Keller's reagent. The small grains in the aluminium matrix and the moderate alignment of particulates in the extrusion direction are clearly seen. The particulate distribution is very homogeneous. The grain size is 5.0 2 1.9 mm. 2.2. X-ray di€raction measurements Samples for X-ray di€raction (XRD) measurements were made by cutting the as-received composite in 3 mm thick plates with a diamond cut-o€ wheel. Then, approximately 25 mm was removed by electropolishing to remove any surface e€ects from the cutting. A reference powder including matrix alloy as well as SiC particulates was made from the same material by grinding using SiC or ZrO2 grinding paper. The powder used in this investigation was not annealed after grinding, but, for reference, identical powder has previously been given a stressrelieving heat treatment at 3008C for 30 min before X-ray measurements in order to study possible e€ects of the grinding on the peak position and shape. The average grain size of the powder was less than 1 mm, i.e. well below the matrix alloy grain size or particulate size. For X-ray measurements, an 00.5 mm thick layer of powder was put in a 10 mm diameter cavity in a substrate. The powder surface level coincided with the surrounding substrate surface, and correct placement of the powder was assured by alignment checks above as well as below the cavity (checks cannot be made on the powder surface) using the same device as for the composite samples. The measurement procedure commenced with a solution treatment at 5028C for two hours, i.e. the standard solutionizing for unreinforced 2014 alloy. The samples were then immediately (without intermediate cooling or quenching) moved to the heating stage on the goniometer, which was set to the measurement temperature, 2008C. Di€raction measurements were con®ned to the Al alloy matrix. They were performed using a Jeol Ogoniometer with a position-sensitive detector (PSD). The X-ray radiation used was CuKa, where the diffraction angle for the Al (422) peak is 2y = 137.458. For a thorough description of the method, see, e.g. Ref. [12]. The di€racted peak position as a function of time at constant temperature was measured using dif-

Table 1. Chemical composition of the 2014 matrix alloy (wt%) [10] Cu 4.32

Mn 0.71

Si 0.7

Mg 0.5

Fe 0.24

Zn 0.09

Ti 0.028

Cr 0.02

Ni 0.01

H (ppm) 0.5

Al Balance

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Fig. 1. Optical micrograph of a sample cut parallel to the extrusion direction. Etching with Keller's reagent reveals the grain boundaries in the matrix. The particulates are moderately aligned in the extrusion direction and very homogeneously distributed.

fracting lattice planes oriented parallel to the sample surface. This measurement gives information about lattice spacing variations with time, changes in the hydrostatic component of the stress as well as on changes in lattice parameter due to changes in matrix composition during formation of precipitates. The di€ractometer was not calibrated for exact lattice parameter determination, since changes rather than absolute values were of primary concern in the investigation. The di€raction peak positions were determined by least-squares ®tting of a pseudo-Voigt function to the measured peaks [13]. The Ka1 and Ka2 peaks were both used in the ®tting procedure, and the ®tting parameter for the Ka1 peak was taken as the peak position. The pseudo-Voigt function is the sum of a Gaussian and a Lorentzian function, where the fraction of each function type is a parameter. In addition to peak position and intensity it also has the peak width as a parameter. To facilitate XRD measurements at elevated temperatures, a special sample holder with a conduction heater has been constructed [14]. The temperature was controlled with a Eurotherm Controller Type 070. The thermocouple monitoring the temperature was positioned between the heating conductor and the sample. A second thermocouple attached to the sample surface and connected to a single-channel recorder showed temperature variations within 228C. 2.3. Transmission electron microscopy Transmission electron microscopy (TEM) samples were prepared by cutting 1 mm slices of the as-

extruded material using a diamond cut-o€ wheel. From these slices 3 mm diameter discs were cut out using an ultrasonic disc cutter. The discs were then mechanically ground and polished to a thickness of 050 mm using a 16 mm diamond grinding disc and 3 mm and 1 mm diamond paste in succession. Finally, ion milling was used to obtain electron transparent foils for transmission electron microscopy. The considerable hardness di€erence between the matrix and the SiC particulates necessitated a very low ion incident angle (7.58) to obtain a smooth foil surface. It has been concluded by Vogelsang et al. [15] that ion milling itself does not alter the dislocation density in moderately thick TEM foils. The thinned samples were investigated using Philips EM 400 TEM and high-resolution Philips CM 20 UT TEM microscopes operated at 120 and 200 kV, respectively. The former microscope was equipped with a furnace-type heating stage, which was used to study the microstructure of the samples in situ at elevated temperatures. The high-resolution instrument was equipped with a LINK QX 2000 Xray microanalysis system, by which qualitative and quantitative element analysis can be performed with an e€ective probe diameter as small as 5 nm. 3. RESULTS

3.1. X-ray di€raction characterization The lattice parameter in the aluminium alloy matrix as a function of time t at 2008C after cooling down from the solution temperature is shown in Fig. 2(a) for a composite sample (ac) as well as the

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measure of the peak width is the full width at half maximum (FWHM) which is the angular width at 50% of the maximum intensity of the peak. In Fig. 3 the FWHM of the Al alloy matrix (422) Ka1 peak of the composite is shown as function of time at 2008C; the peak broadening decreased signi®cantly with time, indicating a decreased variation in microstress. A similar but much less pronounced peak width decrease was seen comparing (at room temperature) the powder heat treated at 3008C with the as-ground powder, indicating that some plastic damage emanating from the grinding has been relieved by the heat treatment. The peak positions were, however, identical and also independent of grinding paper type. 3.2. Transmission electron microscopy characterization

Fig. 2. (a) Lattice parameter for composite sample (ac) and reference powder (ap) as a function of time at 2008C after cooling from 5028C. (b) Lattice parameter di€erence Da = acÿap between the curves in (a). It should be noted that Da = Da0 as indicated by the horizontal dashed line.

reference powder (ap). At time t = 0 the temperature 2008C is reached upon cooling. Both the sample and the powder exhibit an initially rapid increase in plane spacing, but the derivative of the sample curve decreases after about 15 h to practically zero. For the powder, a similar decrease is found after about 10 h, but the curve is not horizontal until 050 h have passed. Stress measurements on the powder using the sin2c-method [12] shows that it is stress-free at all measured temperatures and three-dimensional stress analysis [12] on the composite samples indicates that an entirely hydrostatic stress state with only marginal exceptions prevails in the composite samples. Figure 2(b) shows the di€erence in lattice parameter Da = acÿap obtained from Fig. 2(a). Da describes a hump with a rather steep increase towards a maximum at t = 10 h and a subsequent more moderate decrease, reaching a plateau after about 50 h. For tr = 50 h, the ®nal di€erence Da is equal to the initial di€erence Da0 at t = 0. The width of the di€racted peak re¯ects the variation in microstress, i.e. the variation in stress between di€erent points within a grain of the diffracting phase. A commonly used quantitative

3.2.1. Ambient temperature observations. The TEM micrographs in Figs 4(a) and (b) show the material at di€erent magni®cations at room temperature. The vertical lines in the SiC particulate in Fig. 4(a) indicate stacking faults giving polytype variations typical for SiC. Figure 4(b) shows typical CuAl2 precipitates in the matrix of y phase (large spherical) and y' phase (lath-like). The dominant polytype in SiC was 6H-SiC. The precipitates are more clearly seen in the lattice-resolution images in Fig. 5 which shows (a) the y' precipitate and (b) the y precipitate. The laths are typically 50±150 AÊ wide and between 500 and 1000 AÊ long. A typical diameter of the y precipitates is 800 AÊ. Nano-probe X-ray microanalysis of the matrix alloy showed that the matrix is depleted of Cu (the major precipitate-forming element) in the precipitate-free regions. The large, spherical precipitates contained the highest relative concentration of Cu, while some of the smaller rectangular precipitates also contained Mg. A few lath-formed precipitates also indicated presence of Si and Mn. Phase identi®cation of precipitates to y and y' was made by measuring lattice fringe spacings in precipitates

Fig. 3. The full width at half maximum (FWHM) as a function of time at 2008C for the matrix alloy (422) Ka1 peak of the composite.

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Fig. 4. Transmission electron micrographs of as-prepared SiC particulate reinforced AA 2014 showing (a) overview with a SiC particulate and an Al matrix grain, and (b) precipitates in Al matrix of y phase and y' phase with composition CuAl2.

similar to the ones shown in Figs 5(a) and (b). Based on the expected composition of y and y' in Al±Cu alloys [16] and the above chemical analysis, the expected composition CuAl2 was inferred. The matrix±particulate interface was also studied using TEM, which showed no evidence of pore formation or interphasial layers. Figures 6(a) and (b) are edge-on high-resolution images from two typical interface regions showing an atomically sharp interface between matrix and SiC phases viewed close to edge-on. 3.2.2. Observations during solutionizing and ageing. The microstructural development of the TEM foils during slow heating and cooling was followed in situ between room temperature and solution temperature. The heating rate was 088C/min with short interruptions for photographing in order to avoid thermal drift of the sample when heated. Figures 7(a)±(c) shows the matrix (a) at room temperature, and (b) and (c) during heating at 2308C and 3858C, respectively. During heating some small precipitates disappear before the solution temperature is reached, while others grow slowly as di€usion is enhanced, which can be seen by comparing Fig. 7(a) with Fig. 7(b)

and (c). Reaching 4858C the remaining precipitates dissolved almost instantly, simultaneous with extensive buckling of the foil. The nucleation stage of precipitation at 2008C after solutionizing is depicted in Figs 8(a)±(c), which shows a region in the matrix after di€erent TEM holding times in situ at 2008C. At the high magni®cations employed, thermal drift was a severe obstacle to sharp images. The tiny dark spots are probably y0, sometimes also referred to as GP2 zones [16], typically 50±100 AÊ in diameter. Most precipitates present in Fig. 8 have formed already at 24 min of ageing [see Fig. 8(a)], but ®rst appear ill-de®ned. After prolonged ageing the precipitates formed sharper edges and showed a stronger contrast [see regions marked with arrow heads in Figs 8(b) and (c)]. No microstructural changes were observed in TEM between 61 and 135 min of ageing. 4. DISCUSSION

4.1. Stress state As opposed to the bulk composite samples, the very ®ne powder of the same material should not

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Fig. 5. High-resolution transmission electron micrographs of (a) lath-formed CuAl2 (y') precipitate and (b) spherical-like CuAl2 (y) precipitate embedded in Al matrix.

contain any stresses due to constraints between the di€erent phases. The precipitation-induced stresses in the powder are believed to be negligible and, indeed, stress measurements on the powder indicated zero stress regardless of temperature. It follows that changes in the lattice parameter of the powder are solely a result of compositional changes due to precipitation. Regarding the composite samples, the di€erence in coecient of thermal expansion (CTE) between the matrix and the reinforcement, which for the present system is around 20  10ÿ6/K, may induce large mis®t strains (and stresses) during temperature changes. The lattice parameter of the samples is hence dependent on thermal and precipitation-induced stresses and on compositional changes as a consequence of precipitation. Comparing the composite and powder curves in Fig. 2(a) it is obvious that the precipitation kinetics di€ers. The initially higher rate of lattice parameter change in the composite and the shorter time for the composite to reach steady state indicate that the nucleation rate as well as the growth rate for precipitates are higher. At a given temperature, the state of residual stress and the dislocation density exert in¯uence on these rates. Calculations by Persson [17] using the Eshelby equivalent inclusion

model (see, e.g. Withers et al. [18] for a summary) have shown that high thermal mis®t stresses may exist locally around the particulates. The dislocations may act as nucleation sites for the semicoherent y' precipitates (cf. Section 4.2), and the stresses provide a driving force for di€usion and hence growth of precipitates. Since the powder is stress free, the appearance of the powder curve in Fig. 2(a) is the result of depletion of Cu solute atoms from the Al-rich (matrix) phase during precipitation, while the composite sample curve is in addition a€ected by formation and relaxation of stresses. According to Mondolfo [16] the Al lattice parameter decreases by 0.002 AÊ for each per cent of Cu addition. Hence, the lattice parameter increase seen in Fig. 2(a) was the e€ect of a reduction in Cu content in the matrix from 4.32% to 02.6% during precipitate formation. Assuming that the volume of precipitates and consequently the compositional changesÐdespite di€erences in reaction ratesÐare comparable in the composite bulk samples and the powder, Fig. 2(b) shows the behaviour of the samples compensated for e€ects of precipitates. The formation and subsequent relaxation of precipitate-induced stresses cause a hump in the lattice parameter vs time curve as shown in Fig. 2(b), where the e€ects of compo-

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Fig. 6. Edge-on high-resolution transmission electron micrographs (a) and (b) from two typical matrix/ SiC-particulate interfaces showing an atomically sharp interface.

sitional changes due to precipitation have been removed from the sample behaviour by subtraction of the powder in¯uence. Caution must be used during evaluation of threedimensional stress states measured by X-ray di€raction. Such measurements require a stress-free reference, e.g. a powder, and especially during elevated temperature measurements there may be discrepancies between the powder and sample compositions, giving erroneous three-dimensional stress tensor results. At non-steady state, i.e. for 0 < tR = 50 h, the di€erent kinetics of precipitation render stress calculations by comparison with a stress-free powder impossible. At steady state, the di€erence between the curves in Fig. 2(b) is a measure of the hydrostatic component of the stress in the matrix of the samples, and hence Da = acÿap depicts the e€ects of thermal constraints between the matrix material and the reinforcement on the matrix lattice parameter of the composite. From Fig. 2(b) it can be deduced that the hydrostatic component of the stress remains una€ected by precipitation, since Da = Da0. This equality furthermore validates the

above assumption regarding the precipitate volume; the samples and the powder respond to the heat treatment in the same way as far as the amount of precipitates is concerned. The values of Da correspond to a rather high tensile strain in the matrix, typically 1  10ÿ3. However, measurements of strain at 2008C after slow cooling from 5008C on an Al/SiC composite by Withers et al. [19] show similar but slightly lower strain values than the present work, but the volume fraction of reinforcement in the former work was lower by a factor of three. Withers et al. also note that the two phases of the composite retain signi®cant elastic strains even after 2 h at 5008C, and that the strains in the Al matrix are generally larger than expected from Eshelby modelling. For the present composite, Eshelby calculations similarly predict lower thermal stresses and strains in the matrix than measured. Earlier work by Weiland and Ericsson [20] indicates that low stress values should be expected at 2008C. The Al matrix XRD peak width depicted in Fig. 3 indicates a variation in microstress with time in the

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Fig. 7. Transmission electron micrographs of matrix with precipitates at (a) room temperature before heating, (b) 2308C during heating, and (c) 3858C during heating. The locations of dissolving y' precipitates are marked by arrow heads.

matrix. Microstresses are usually inhomogeneous on a macroscopic scale, and may be inhomogeneous on a microscopic scale as well. The decreasing peak width with time implies that the microstress state becomes more homogeneous during annealing. Microstresses may be relieved by di€usional relaxation and formation of dislocations around particulates and precipitates. Further stress reduction is obtained as the coherency of precipitates is lost by formation of, e.g. y phase as shown in Fig. 5(b).

Stress relaxation is also greatly facilitated by the large amount of grain boundaries in the smallgrained composite (cf. Fig. 1). 4.2. Precipitation Precipitation in Al±Cu alloys commences with formation of Guinier±Preston (GP) zones as thin platelets. In alloys where the Cu:Mg ratio is higher than 8:1, the main hardening agent is CuAl2 formed from the GP zones via the coherent y0 phase, the

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Fig. 8. Transmission electron micrographs of matrix with alleged y0 precipitates at 2008C after (a) 24 min, (b) 61 min, and (c) 135 min at temperature after cooling from solution temperature. Locations of nucleated precipitates are marked by arrow heads.

semi-coherent y', and subsequently incoherent y precipitates [21]. For ratios between 8:1 and 4:1 both CuAl2 and CuMgAl2 are active [16]. The composition of the present matrix alloy (Table 1) gives a Cu:Mg ratio of 8.64:1. Thus, the observed CuAl2 precipitates are to be expected. Figure 5(a) shows that the y0 precipitate is coherent with the matrix in the h100i direction. The semicoherency and the plane spacing indicate a close similarity with y' precipitates formed in unrein-

forced AA 2014 alloys. The y' phase has lattice parameters a = 4.04 AÊ and c = 5.80 AÊ, i.e. the a value is almost identical to that of pure Al. Consecutive {200} planes are resolved in the matrix as well as in the precipitate in the micrograph. These precipitates most probably nucleate at dislocations, and grow along a h100i direction [21]. It has been argued by Vogelsang et al. [15] that dislocations may be lost through the surfaces of a thin TEM foil. Hence, a dislocation density lower

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than the actual density can be expected in the TEM samples. Previous investigations of similar composites show dislocation accumulation in the vicinity of the reinforcement [7, 15], and Dutta and Bourell [7] point out that precipitation is associated with dislocations. In the present material only a few stray dislocations have been observed in the numerous micrographs taken. However, during in situ heating and cooling as in the present study, elastic relaxation by the observed buckling of the TEM foil may relieve thermal stresses, thus suppressing dislocation generation [15]. Furthermore, as seen in Figs 1 and 4, the distribution of precipitates is very homogeneous, with no sign of preferential nucleation near SiC particulates, which may support the view that no extensive dislocation formation has occurred around particulates. Homogeneous distribution of dislocations and precipitates in SiC reinforced Al alloys has been reported by Christman and Suresh [2] as well as Cottu et al. [22]. In Fig. 5(b) a large, spherical-like precipitate is imaged in high-resolution mode. The incoherent precipitate was identi®ed as y, i.e. a body-centred tetragonal CuAl2 phase with lattice parameters a = 6.07 AÊ and c = 4.87 AÊ. The y phase lattice planes resolved in the ®gure are {110}, with a planar distance of 4.29 AÊ, and {200} planes of the matrix are also seen. In the precipitate, the direction h111i is parallel to the matrix h100i, indicating that the precipitate has formed directly from a matrix defect rather than from a semi-coherent y' plate [16]. Preferred nucleation sites in this case are grain boundaries and y'/matrix interfaces [23], as can be seen in Fig. 7. Chemical and structural analysis showed no evidence of precipitates containing large amounts of Mg, Si, or Mn. It is probable that the traces of these elements emanate from the matrix; if the precipitate thickness is less than that of the foil or a part of the probe has drifted o€ the precipitate during analysis, the matrix composition will a€ect the result. From the microstructural characterization it is hence clear that the material exhibits the same precipitates as does unreinforced AA 2014. However, the X-ray di€raction results show that the rate at which these precipitates nucleate and grow is increased by reinforcement-induced thermal mis®t stresses. To date, the commonly used values for ageing time and temperature have been identical to those for unreinforced material. The present results indicate that recommendations regarding times and temperatures for heat treatment of reinforced Al alloys are desirable to further improve material properties. From Fig. 8 it is obvious that a longer exposure at 2008C would have been required for the precipitation in the TEM foil to proceed past the y0 stage. A comparison with Fig. 2(a) shows that initially the transfer of Cu from the matrix to the precipitates was expeditious. Still, as an example, about 15 h

ageing was needed for 1% Cu (out of the total 4.3% originally present in the matrix) to transfer from the matrix to the precipitates, corresponding to a lattice parameter increase by 0.002 AÊ. The particulate±matrix interface is well de®ned (see Fig. 6). There were no signs of pore formation, and no reactions between matrix and reinforcement have occurred during production or subsequent heat treatment. Moreover, since no increased precipitate concentration is observed at the interface, it can be concluded that the interface itself did not contribute appreciably to the ageing behaviour of the composite. The good bonding expected from the well-de®ned interface would indicate a material with excellent load-transferring properties and hence bene®cial sti€ness and strength. Assessment of mechanical properties and their relations to residual stresses and microstructure is the objective of a study in progress.

5. CONCLUSIONS

It is demonstrated that the kinetics of precipitation in a SiC particulate reinforced AA 2014 alloy and a reference powder of the same composition di€er substantially. The di€erence is due to thermal mis®t stresses in the composite, which are not present in the powder. The stresses do not in¯uence the ®nal fraction of precipitates, but increase the rate of precipitation. The time to reach equilibrium in the composite is about one third of the time required in the stress-free condition. Obviously, standards for heat treatment of reinforced aluminium alloys are required. The presence of SiC particulates does not in¯uence the type of precipitates formed, since the y and y'-phase CuAl2 precipitates observed are identical to those forming in unreinforced AA 2014. The SiC particulates are stable in the temperature range investigated, and there is no evidence of interfacial reactions between the particulates and the matrix alloy. The improved properties of aluminium-based composites cannot be attributed to changes in precipitation behaviour compared to unreinforced material. The increased strength is rather an e€ect of load transfer from the matrix to the reinforcing phase. The SiC particulates in the investigated material are well-bonded to the matrix, thus promoting load-sharing between the phases. The increased strength is suggested to be an e€ect of load transfer from the matrix to the reinforcing phase.

AcknowledgementsÐThe investigated composite was supplied by Renault in France through Eric Hanus, to whom thanks are also directed for valuable discussions and for providing some optical micrographs.

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