Al–Si metal matrix composites

Materials Science and Engineering A 486 (2008) 427–432

Study on the yield behavior of Al2O3–SiO2(sf)/Al–Si metal matrix composites Jihua Peng ∗ , Dehong Han Wenfang Li, Jun Du Yong Xie, Guanjun Liu College of Mechanical Engineering, South China University of Technology, Guangzhou 510640, PR China Received 22 March 2007; received in revised form 9 September 2007; accepted 10 September 2007

Abstract As the yield behavior of Al2 O3 –SiO2 (sf)/Al–Si MMCS is concerned, effects of heat treatment and parameters of short fiber, including volume fraction, size, distribution mode, were investigated. Dislocation configurations adjacent to interface between matrix and fiber were observed by TEM. Macro-yield stress (σ 0.2 ) and micro-yield stress (σ MYS ) vary with parameters of short fiber, and effects of these parameters on σ 0.2 appear to be opposite to those on σ MYS . This phenomenon was interpreted by thermal residual stress in matrix and dislocation configuration. Suitable quenching followed aging treatment is an effective method to enhance the σ MYS and the σ 0.2 simultaneously. For the specimen with heat treatment of 550 ◦ C/1 h WQ + 170 ◦ C/6 h (T6) AC, σ 0.2 and σ MYS reach 200 and 58 MPa, respectively, and they are almost as twice as those as-cast. © 2007 Published by Elsevier B.V. Keywords: Al–Si metal matrix composite; Micro-yield strength; Alumina silicate short fibers; Heat treatment

1. Introduction Metal matrix composites (MMCs) have attracted great attention due to their higher specific strength and Young’s modulus, higher micro-plastic strain resistance, adjustable heat conductivity and thermal expansibility. They have been applied in precision instruments and optical systems such as guide for navigation, infrared detector, astronomical telescope, where better dimensional stability is required. SiCp/Al composites were considered as promising instrument materials in America [1]. It is important to understand the deformation behavior and mechanisms of such composites at low stress level, because precision devices require severe resistance to micro-deformation. Usually, micro-yield strength, the stress value corresponding to the material plastic strain of 1 × 10−6 , is identified as one of indicators of the dimensional stability. In 1961, Brown and Lukens (B–L) proposed out a parabola relationship between micro-yield stress σ and micro-plastic deformation εp of the metal materials [2]. This relationship was derived from the speculation of grain boundary barrier to dislocation movement in pure metal. Study on micro-deformation behavior of SiCp (or Al2 O3 p)/6061Al showed that the B–L relationship is suit-

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0921-5093/$ – see front matter © 2007 Published by Elsevier B.V. doi:10.1016/j.msea.2007.09.016

able to describe the micro-yield behavior of some MMCs [3]. Compared to the unreinforced alloys, reinforcements result in higher density of dislocations and larger residual internal stresses. It is shown that micro-deformation is dependent on thermal residual stress, dislocation structure, also morphology and size of reinforcements of particles or whiskers in MMCs [3–13]. Some heat treatment and pre-treatment were developed to improve the micro-yield stress of Al-matrix composites, such as thermal-cold cycling, deep colding, and thermal residual stress creeping [4]. Among these processes, quenching [5] and quenching followed aging [6–8] were proved to be beneficial for particle or whisker reinforced Al-matrix MMCs. In sub-micron Al2 O3 p/2024Al-MMC, because the thermal residual stress is less, tiny and uniformly dispersed precipitates S (Al2 CuMg) play a great role to increase the micro-plastic deformation [9]. It is well known that the ultimate strength and creep resistance of short fiber reinforced Al-matrix MMCs could be improved [10,11]. Less study was focused on dimensional stability on them. Crystallized alumina silicate short fiber, as one of the low-cost reinforcement has been used in aluminum-based MMCs [12,13]. In this present study, by means of observing dislocation structure, combining with analysis of thermal residual stress, effects of various factors on the yield behavior Al2 O3 –SiO2 (sf)/Al–Si MMC were investigated, i.e., volume fraction, size, alignment of short fibers and heat treatments.

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Table 1 Elemental composition of Al–Si alloy (wt%) Element

Fraction (wt%)

Si Cu Fe Mg Mn Others Al

12.60 0.15 0.30 0.24 0.22 0.72 Balance

Table 3 List of conditions of short fibers used in MMCs fabrication and followed treatment Heat treat type

Specification

Annealing Quenching Aging

300 ◦ C/3 h FC 550 ◦ C/1 h WQ 550 ◦ C/1 h WQ + 170 ◦ C/2, 4, 6 h (T6) AC

Note: FC, furnace cooling; WQ, water quenching; AC, air cooling.

2. Materials and experimental procedure The Al2 O3 –SiO2 (sf)/Al–Si MMCs were fabricated by infiltration squeeze method. The matrix material used was eutectic Al–Si alloy containing 12.6 wt% Si with elemental composition seen in Table 1. The matrix alloy was obtained by melting A00 pure aluminum and aluminum–silicon alloy containing 20 wt% Si together at definite proportion. The reinforcement was lowcost alumina silicate (Al2 O3 –SiO2 ) short fibers crystallized at 800–1100 ◦ C, with diameter of Ø8–10 ␮m, which were grinded, washed and size screened. Table 2 lists parameters of short fiber during fabrication and followed treatment. In this table, in order to investigate the effect of one certain factor on yield behavior of MMCs, only this factor varies (the range is shown in the parentheses), while other factors are fixed. When effects of fiber volume fraction (Vf ) and heat treatments on σ MYS of the MMCs were studied, fibers with length ranging from 60 to 120 ␮m were mixed and distributed randomly in Al–Si alloy matrix. Typical heat treatments conducted were listed in Table 3. To study effect of distribution mode of fibers, two pre-forms with Vf 30% were prepared. One pre-form had fibers alignment ratio of about 75% and the other one had random fiber distribution. Samples as-cast were loaded along the fiber’s axial direction for the MMC with aligned fibers. Dislocation structures near the interface between matrix and short fiber were observed with a HITACH 800 type transmission electron microscope (TEM). The thin foils for TEM observation were prepared from unloading samples by ion milling. Requirements of double beams were met as perfectly as possible during obtaining images of dislocation structures. The continuous uniaxial tension tests were carried out on an Instron-5569 test machine with data processing software Merlin at room temperature and in air environment. A special extensometer with an accuracy of 10−7 was used for deformation measurement. Dog-bone shaped coupons with 2 mm thickness, 4 mm width and 8 mm gauge length were machined. Original

curves of loading–displacement and stress–strain (σ–ε) were obtained by Merlin software. The Young’s modulus (E) for each specimen was induced by fitting the σ–ε curve prior to the elastic limit. The residual plastic strain (εp ) was calculated through the formula: ε−σ εp = (1) E The stress corresponding to εp = 1 × 10−6 is referred as the micro-yield strength σ MYS [14]. σ 0.2 is the stress corresponding to εp = 0.2%. From lots of data points (σ and εp ), the following B–L relationships [2] were fitted, and coefficients of K and σ 0 were estimated: σ = σ0 + Kεp 1/2 and K=

(2)

 2Gσ0 ρd 3

(3)

where ρ is the dislocation density, σ 0 the stress to move the first movable dislocation, G the shear modulus of MMC, and d is the grain size. 3. Experimental results Table 4 presents macro-yield stress (σ 0.2 ), micro-yield stress (σ MYS ), and coefficients of K and σ 0 fitted by B–L relationship of all tension tests. For the effect of fiber volume fraction (Vf ), the 1/2 fitness curves of σ–εp were presented in Fig 1a. Variations of K and σ 0 with Vf were shown in Fig. 1b. It is shown that both σ 0.2 and the young’s modulus of Al2 O3 –SiO2 (sf)/Al–Si composites increase with Vf . When the interface between reinforcing fiber and metal matrix is strong, the load applied on matrix could be transferred to the fiber through interface. However, the σ MYS of Al2 O3 –SiO2 (sf)/Al–Si MMCs decreases with Vf . It is obvious that influences of Vf on σ MYS and σ 0.2 are opposite to each other. The stress for moving the first dislocation (σ 0 ) decreases

Table 2 List of parameters in MMC fabrication and treatment Effect of factors studied

Vf (%) L (␮m) Fiber distribution Heat treatment

Volume fraction, Vf

Fiber length, L

Fiber distribution

Heat treatment

(10, 20, 30, 40) Mixture of 60–120 Random Annealed

20 (60, 80, 125) Random Annealed

30 80 (Random, alignment) As-cast

30 Mixture of 60–120 Random See Table 3

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Table 4 Results of all tension tests Factors

No.

σ 0.2 (MPa)

σ MYS (MPa)

σ 0 (MPa)

K (×103 MPa)

E (GPa) 68.9 71.1 81.1 89.8

Volume fraction (%) 10 20 30 40

1 2 3 4

120 130 142 165

78 53 37 24

75 52 34 23

2.5 3.0 1.75 1.5

Fiber length (␮m) 60 80 120

5 6 7

105 117 125

58 42 30

54 35 26

3.0 5.0 2.5

Fiber distribution Random Alignment

8 9

102 132

48 29

46 23.5

3.0 6.0

11 12 13 14 15 16

125 165 142 160 188 200

24 20 37 33 49 58

22.5 19 34 30 46 54

2.0 0.75 1.75 0.85 2.0 3.0

Heat treatment As-cast Quenched Annealed Aged 2 h Aged 4 h Aged 6 h

monotonically with Vf , while K reaches a peak when Vf is 20%. From Table 4, unlike K, it is obvious that σ 0 always keeps pace with micro-yield stress σ MYS . σ MYS and σ 0.2 of Al2 O3 –SiO2 (sf)/Al–Si MMCs are affected remarkably by the length of short fibers (aspect ratio). The longer fibers are, the lower the σ MYS of the MMC is. On other hand, the macro-yield strength of the MMC increases with the length of the fibers. It is indicated that the σ 0.2 of Al2 O3 –SiO2 (sf)/Al–Si MMCs with aligned short fibers is about 30% higher than that with randomly distributed fibers. However, σ MYS of the aligned MMC is much lower than that of the MMC with randomly distributed fibers. It is shown that σ MYS and σ 0.2 of Al2 O3 –SiO2 (sf)/Al–Si MMCs are affected by heat treatments greatly. Among all heat treatment, quenching resulted in the minimum σ MYS , and samples as-cast presented the lowest σ 0.2 . With increase of aging hold time, σ MYS and σ 0.2 increase simultaneously. Aged for 6 h (550 ◦ C/1 h WQ + 170 ◦ C/6 h (T6) AC), σ MYS reaches

58 MPa, and σ 0.2 reaches 200 MPa for Al2 O3 –SiO2 (sf)/Al–Si MMC with Vf 30%. They are as twice as those of samples as-cast. Dislocation structures in Figs. 2 and 3 were from annealed samples with short fiber random distribution. Mixture of fibers with length from 60 to 120 ␮m was used in the sample of Fig. 2, while the fiber length was fixed as 80 ␮m in the sample of Fig 3. Fig. 2a and b shows dislocation structures along fiber’s longitudinal section and cross-section in a sample with Vf 20%, respectively. Fig 2c is from a sample with Vf 40%. The dislocation density in matrix near interface increases with Vf . The interfaces in Fig. 2 appear “clean”, which present no precipitates. In Fig. 3, higher dislocation density in matrix and some micro-crack in the fiber were observed. For the as-cast sample with aligned short fiber, the dislocation configuration in matrix was shown in Fig. 4. In this figure, the highly densified, curved and closely tangled dislocations are observed in the middle part of the matrix between two parallel fibers. Unlike the sample as-cast, low dislocation density and less

Fig. 1. Effects of fiber volume fraction (Vf ) on micro-yield behavior of Al–Si MMCs.

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Fig. 2. Dislocation configurations adjacent to the interfaces in the matrix of composites: (a) along fiber’s longitudinal section; (b) along from cross-section of fiber specimen 2; (c) from specimen 4.

network of dislocation are shown in the sample with treatment of 550 ◦ C/1 h WQ + 170 ◦ C/4 h AC (Fig. 5). Some precipitates are also formed near the interface during aging. 4. Discussion Tensile thermal residual stress in the metal matrix was generated due to the great difference in thermal expansion coefficients (CTE) between metal matrix and reinforcing fibers during the cooling process of the fabrication and followed heat treatment. During MMC fabrication, dislocations in the matrix were unavoidably generated due to the stress relaxation and matrix

Fig. 3. Dislocation configurations adjacent to the interfaces in the matrix of specimen 5.

deformation in the cooling process. Because there exists great stress concentration in the matrix nearby the interfaces or the tip of SiC whisker, the initial stage of the plastic deformation, i.e., the micro-deformation, usually begins from the matrix nearby the tip or interfaces of SiC whiskers [15]. It is known that amount of thermal residual stress (σ TRS ), the dislocation density and state nearby the interfaces have a great influence on the micro-yield behaviors of the composite [3–9]. It is interesting to review previous study on micro-yield behavior of Al-MMCs. For Al2 O3p /6061Al [3], dislocation with the lowest density is almost linear for one cycle, dense and tangled for three cycles, lowest and linear for five cycles. Correspondingly, σ MYS presented the minimum for two cycles, and

Fig. 4. TEM micrographs of dislocation states in the matrix of specimen 9.

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increase of σ TRS when Vf increases, thus σ 0 and σ MYS increase with Vf in Table 4 and Fig. 1. According to Eq. (3), K is limited by shear modulus, grain size, dislocation density and σ 0 . Interaction of all these factors results in complexity of K in yield behavior of MMCs. During the stage of macro-yield, work hardening in matrix derived from dislocation multiplication and load transfer play a significant role in determination of σ 0.2 [5]. As shown in Table 4, σ 0.2 increases with fiber volume fraction. For aluminum matrix reinforced with SiC fibers, Nardone and Prewo [16] investigated the plastic deformation of MMC. Assuming that the load value at the ends of the fibers equals to the average strength of the metal matrix, and basing on the tensile load transfer at the ends of the fibers, prediction for yield strength of MMC was proposed: Fig. 5. TEM micrographs of dislocation states and precipitates near interface in the matrix of specimen 9.

maximum for five cycles. After TRSC treatment of SiCw /Al, the dislocation density nearby the interface is reduced obviously, and it gave high σ MYS [4]. The maximum of σ MYS was obtained after T6 treatment (quenching + aging), in which dislocation is less and linear [8]. From these studies, the dependence between σ TRS and dislocation density and state exists. From all tension test results in this study (Table 4), it is obvious that the σ MYS is very close to the first movable dislocation driving stress σ 0 . So those factors which affect σ 0 are important in determination of σ MYS . As the stress to move the first movable dislocation is concerned, critical resolution shear stress (CRSS) of a certain slip system is definite. Related to this CRSS, the total stress for driving the first movable dislocation could be regarded as a definite value σ th . It is plausible to suggest the following formulation: σth = σTRS + σex

(4)

where σ TRS is the thermal residual stress in the matrix nearby interface and σ ex is the applied external stress. σ TRS is determined mainly by fabrication process and followed heat treat. σ ex is close to σ 0 . Moreover, the moveable dislocations have to overcome the resistances from the parallel dislocations in the same sliding planes, from short fibers (or reinforcements, precipitates), and from the dislocations in the different sliding planes, such as forest dislocation. Loading transfer (LT) from matrix to fiber [11] could be neglected in the stage of microyield. However, LT, dislocation multiplication and interaction with short fiber must be taken into consideration for the stage of macro-yield. As shown in Fig. 2a and b, the dislocations is very small and linear, they can be considered to be movable. This means that the residual stress σ TRS is low. With fiber fraction volume increasing, the dislocation density increase greatly, and network of dislocation forms nearby interface (as Fig. 2c). Dislocation density and state could be regarded as the indicator of thermal stress. Obviously, σ TRS in Fig. 2c is much more than that in Fig. 2a. The number of movable dislocation in Fig. 2c is less than that in Fig. 2a. Dislocation tangling enhances resistance to a degree in Fig. 2c. However, this role cannot overcome the

 σcy = σmy

     1 l l Vw 2 + + (1 − Vw ) = σmy 1 + Vw 2 d 2d (5)

where σ cy and σ my are the yield strength of MMC and matrix alloy, respectively. l/d and Vw are the aspect ratio and volume fraction of reinforcing fibers, respectively. d is the diameter of the reinforcing fibers. The above formula shows that, the yield strength of the MMC increases with the increase of aspect ratio of the fibers. The σ 0.2 of Al2 O3 –SiO2 (sf)/Al–Si MMCs from specimens 5–7 in Table 4 follows this prediction. Comparing to the σ 0.2 from specimen 2, σ 0.2 from specimens 5–7 looks somewhat low. It may be connected with micro-crack in fibers in Fig. 3. Because of breakage of fiber, the actual fiber length is reduced indeed. Dislocation density from specimen 5 in Fig. 3 is higher than that from specimen 2 in Fig. 2. It means that high σ TRS exists in the matrix near interface in specimen 5. According to Eq. (4), σ MYS from specimen 5 should to be lower than that from specimen 2. Probably because of relaxation of σ TRS induced by fiber breakage, values of σ MYS from these both samples are almost the same. With fiber length becoming long, local thermal residual stress concentration will intensify. Micro-Al2 O3 particles were used to reduce σ TRS in the matrix near interface [9]. That is to say, σ MYS will decreases with fiber’s length, as shown in Table 4. Comparing the dislocation and state in Fig. 4 to those in Fig. 5, it is shown that rearrangement of dislocation occurs during the heat treat of quenching followed by aging. In the matrix around interface. Dislocation density becomes much lower in Fig. 5, it means that a large amount of thermal residual stress is relaxed. Although densified dislocation structure is observed in the middle part of matrix between two parallel short fibers, linear-like dislocations are also found in matrix very close to interface. Thus it is of reason to say that σ 0 from these two samples is almost the same value. In that the σ MYS from specimen 15 is higher than that from specimen 9 as-cast. Close to interface, some precipitates are found, these fine and uniform distribution pay contribution to improvement of σ MYS and σ 02 [9]. As barriers to dislocation movement and source of dislocation multiplication, these precipitates from aging treatment are significant to increase macro-yield resistance.

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550 ◦ C/1 h WQ + 170 ◦ C/6 h (T6) AC, σ 0.2 and σ MYS reach 200 and 58 MPa, respectively.

5. Conclusions In this study, by continuous loading using Instron-5569 tester with a special extensometer with accuracy of 10−7 , effects of many factors on micro-yield strength (σ MYS ) and macro-yield strength (σ 0.2 ) of Al2 O3 –SiO2 (sf)/Al–Si MMCs were investigated. These factors incorporate parameters of short fiber, such as volume fraction, aspect ratio (length), and distribution mode of fiber in matrix, as well as heat treatments. Dislocation configurations were observed by TEM. The roles of thermal residual stress (σ TRS ) and dislocation configuration are analyzed. The following conclusion could be made: (1) The σ MYS of Al2 O3 –SiO2 MMCs is mainly controlled by the thermal residual stress (σ TRS ) and the applied external stress to drive the first movable dislocation (σ 0 ). In these MMCs, the smaller the σ TRS is, the higher the σ MYS is. (2) Effects of parameters of short fiber on σ MYS of Al2 O3 –SiO2 MMCs could be contributed to their effects on thermal residual stress. In the range of Vf 10–40%, σ MYS increases with Vf for the composites as annealed, while macro-yield stress (σ 0.2 ) deceases with Vf . (3) The aspect ratio (Ra ) of short fiber and fiber’s distribution mode also play an important role in yield behavior of Al2 O3 –SiO2 MMCs. σ MYS from coupons as annealed decreases from 58 to 30 MPa with increasement of fiber length from 60 to 125 ␮m, while σ 0.2 increases from 105 to 125 MPa. Comparing to random distribution of fiber, σ MYS reduces, and σ 0.2 rises for composites with aligned fibers. (4) Suitable heat treatment of quenching followed aging is beneficial to improve σ 0.2 and σ MYS of Al2 O3 –SiO2 (sf)/Al–Si composites simultaneously. For the heat treatment of

Acknowledgements The authors are grateful for the financial support of the National Natural Science Foundation of the People’s Republic of China (Grant No. 19972021) and Natural Science Foundation from Guangdong Province in PR China (Grant No. 980563). References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16]

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