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Study of the interfacial reaction between Ti3SiC2 particles and Al matrix
Journal of Alloys and Compounds 738 (2018) 1e9
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Review
Study of the interfacial reaction between Ti3SiC2 particles and Al matrix Jiangbo Zhang b, Wenyang Liu a, Yiming Jin b, Shangjiang Wu a, Taotao Hu b, Yong Li b, Xiangpeng Xiao b, * a b
School of Material Science and Engineering, Jiangxi University of Science and Technology, Ganzhou 341000, China Jiangxi Nonferrous Metal Processing Engineering Technology Research Center, Ganzhou 341000, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 September 2017 Received in revised form 9 December 2017 Accepted 12 December 2017 Available online 14 December 2017
MAX phases have a unique laminated ternary structure and excellent self-lubricating properties. Their tribological properties are similar to that of graphite in addition to having excellent high-temperature oxidation resistance. Because of these features, MAX is an ideal external reinforcement for metalbased self-lubricating composites. Here, Ti3SiC2 powder and pure aluminum (Al) powder were combined as the raw materials to prepare Ti3SiC2/Al-based composite without significant interface reaction between particles and matrix using spark plasma sintering. The degree of relative density of the composite was estimated to be as high as 96.6%. The features of the interfacial reaction between Ti3SiC2 and Al matrix were studied over varying times of isothermal treatment. The results showed that the holding temperature was the key factor in the initiation of the interfacial reaction. No interfacial reaction was observed when the reaction was conducted at 500 C/10 h. At a holding temperature of 600 C, the Si atoms in Ti3SiC2 entered the matrix to form Al(Si) solid solution. Ti then bound to both Al and C to form Al3Ti and TiC. Varying the holding time had a significant impact on the morphology of the reaction products. At temperatures below 700 C, prolongation of the holding time caused particles to gradually eVolve from irregular to lath-shaped. Ti3SiC2 completely decomposed at holding times of 10 h. The performance of the material also changed throughout the process of thermal treatment. As the holding time was prolonged, the density of the material continuously decreased, while the hardness first increased and then decreased. The coefficient of friction and degree of wear first decreased and then increased. The minimum coefficient of friction and wear were 0.23 and 0.01 102 g, respectively. © 2017 Elsevier B.V. All rights reserved.
Keywords: Aluminum matrix composite Frictional wear Interfacial reaction Ti3SiC2
Contents 1. 2. 3.
4.
Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 Experimental procedures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 Results and discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 3.1. Analysis of the composite after sintering densification . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 3.2. Analysis of the composite after isothermal treatment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 3.3. Performance analysis of the composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 3.3.1. Hardness and density . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 3.3.2. Coefficient of friction and wear rate . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8 Acknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9
* Corresponding author. E-mail address: [email protected] (X. Xiao). https://doi.org/10.1016/j.jallcom.2017.12.123 0925-8388/© 2017 Elsevier B.V. All rights reserved.
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1. Introduction The rapid advance of technology and engineering science creates a need for high performance materials. As the use of a single material can no longer meet these requirements, metal matrix composites are being increasingly used because of their excellent, comprehensive performance. Aluminum (Al) matrix composites have the advantage of high specific strength, specific modulus and specific stiffness, and therefore are increasingly utilized in the aviation, automobile and military industries [1e3]. Particles commonly added into the Al matrix are Al2O3 [4], TiC [5], SiC [6] and BN [7], which are hard ceramic particles able to enhance the hardness, and wear and deformation resistance of the Al matrix. However, the properties of these particles are greatly different from those of the Al matrix, which leads to poor plasticity and toughness of the particle-reinforced Al matrix composite, non-uniform distribution of the particles, low binding properties and low secondary process ability [8,9]. Recently, laminated ternary MAX, a novel phase that integrates metallic and ceramic features, has drawn widespread attention [10e12]. MAX refers to a ternary transition metal carbide or nitride, with the general formula Mnþ1AXn where M refers to a transition metal; A is a group IIIA or IVA element; X is C or N, n ¼ 1e3 [13,14]. Ti3SiC2 is representative of MAX compounds. To be specific, Ti3SiC2 displays electrical and thermal conductivity, a high degree of hardness, high elastic modulus and good ductility similar to metals. It also possesses the properties of ceramics, such as thermal stability, high temperature oxidation and heat resistance, creep properties, a low coefficient of friction and excellent selflubricating performance. Therefore, Ti3SiC2 can serve as the ideal reinforced phase for metal matrix compounds [15,16]. Pure Ti3SiC2 does not decompose below 1400 C [17]. The high temperature stability of Ti3SiC2 therefore inhibits an interfacial reaction and preparation of metal matrix composites reinforced by pure particles. However, many studies have shown that when added into a copper matrix, Ti3SiC2 particles will decompose at a much lower temperature, that is, a reduction from an initial 1400 Ce900 C [18,19]. This leads to a severe interfacial reaction and the formation of complex products, depending on the temperature and proportions of the components. In general, the addition of Ti3SiC2 into a copper matrix is a costly approach [20,21]. An Al matrix has the features of low density, high specific strength, high specific stiffness and corrosion resistance. When added to Ti3SiC2, Al matrix composites acquire excellent self-lubricating properties. However, the interfacial reaction between Ti3SiC2 and the Al matrix is the major factor that determines the performance of Al matrix composites. The affinity between Al and O2 is relatively high, and oxidation is easily achieved, which interferes with the study of interfacial reactions, and does not favor the preparation of composites with a high density. Dunia [22] et al. suggested that the dynamic vacuum of the SPS chamber creates a protective environment that is beneficial for oxidation-sensitive materials. In this study, we used SPS sintering preparation of high-density Ti3SiC2/Al composites without internal oxidation and obvious interfacial reaction, to investigate the interfacial reactions between Ti3SiC2 particles and the Al matrix by isothermal treatment, and evaluated the effects of the interfacial reaction on the mechanical properties of the composites. 2. Experimental procedures The raw materials were 300-mesh pure Al powder and 1000mesh Ti3SiC2 powder (see Table in Fig. 1 for particle size determination). After premixing, alcohol was added to a PMQW series Omnibearing planetary ball mill for wet ball milling for 9 h. The ratio of grinding media to material was 11:3 under argon
protection, at a milling rate of 200 r/min. After complete drying and cold molding, sintering was performed in a ST-5 SPS machine for 10 min at a temperature of 600 C and pressure of 40 MPa, followed by isothermal treatment in a muffle furnace at either 500 C, 600 C or 700 C for 1 h, 3 h, 5 h or 10 h. A small muffle furnace was used. The lining material of the furnace was silicon carbide and it operated with resistance heating. Material density was measured according to standard GB/T 1423-1996 “Method of measurement of density for precious metals and their alloys”. Hardness was measured using a MC010-HV-1000 Vickers hardness tester under a load of 0.2 N and holding time of 15 s. The mean of five measurements was taken as the final value. Frictional wear was determined on the M-2000 type wear tester. A block-on-ring test arrangement was used with a friction pair made from 45Cr steel with a hardness of 60 HRC and diameter of 20 mm. The test was performed for 600 s at room temperature. The experimental force was set to 50 N, and friction velocity at 200 r/ min. Each specimen was subjected to coarse grinding, fine grinding and polishing, then corroded with 0.5% HF for approximately 30s. The metallurgical structure was evaluated using a CMM-772 light microscope. Measurement of the phase of the composite was achieved using a PANalytical Axios Max X-ray fluorescence spectrometer. The microscopic structure and micro-area composition were analyzed using an FEI ML650F scanning electron microscope (SEM) and Bruker energy dispersive spectrometer, respectively. 3. Results and discussion 3.1. Analysis of the composite after sintering densification Fig. 2 shows the XRD patterns of the mixed Ti3SiC2/Al powder and Ti3SiC2-reinforced Al matrix composite, whose three diffraction peaks between 30 and 45 are located at 39.5 , 40.8 and 42.5 . The phases of the ball milled powder and the sintered composite correspond to the diffraction peaks of Al and Ti3SiC2, with no peaks from impurities. This indicated that wet ball milling and SPS technology can prepare pure Ti3SiC2-reinforced Al matrix composite. According to Table 1, the comprehensive performance of the sintered Ti3SiC2-reinforced Al matrix composite is very satisfactory, with compactness reaching up to 96.6% and the coefficient of friction as low as 0.28. Fig. 3 shows SEM images of mixed Al and Ti3SiC2 powders after ball milling and the sintered composite. In Fig. 3(a), the large gray particles are Al, with a diameter of 10e80 mm. The white particles of 1e5 mm are uniformly distributed around the large particles. Fig. 3(b) shows the sintered composite, in which fine round particles 1e8 mm in diameter are uniformly distributed and predominate at the grain boundaries. Comparison shows that after SPS, the diameter of the Ti3SiC2 particles greatly increased and were more likely to cluster at the triangular grain boundaries. This is possibly because the poor wet-ability between the ceramic particles and metal led to the failure of the matrix grains to engulf the reinforcement particles. As the Ti3SiC2 particles grew along the sintering neck, they were gradually pushed to the triangular grain boundaries and clustered together. Fig. 4 shows the surface scanning results of Ti3SiC2-reinforced Al matrix composite after SPS at 600 C. As can be seen in Fig. 4(a), the Al matrix takes up a greater proportion of the cross-section, with grain boundaries of elemental Al showing as black. The particles are generally located along the black grain boundaries, with a distinct interface observed between the reinforcement particles and the matrix. The distribution regions of the elements Ti and Si nearly overlap, both located in the black grain boundaries of the Al matrix grains. This further indicates that Ti3SiC2 does not decompose during SPS.
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Fig. 1. Particle size determination of Ti3SiC2.
Fig. 2. XRD patterns of the Al/Ti3SiC2 ball milled powder and sintered composite.
3.2. Analysis of the composite after isothermal treatment Fig. 5 shows the XRD patterns of the sintered Al matrix composites after isothermal treatment for 10 h at different temperatures (500 Ce700 C) and at 700 C for different holding times (1e0 h). In Fig. 5(a), no new phases were produced in the composite after long-term isothermal treatment under 500 C, with Al and Ti3SiC2 still being the major phases. This indicates no apparent interfacial reaction occurring at this temperature. After the
incubating at 600 C for 10 h, the strong diffraction peak of Al and weak diffraction peak of Ti3SiC2 appear, together with the diffraction peaks of new phases. These new phases were identified as Al3Ti and TiC, indicating the occurrence of the interfacial reaction. It is thus inferred that the interfacial reaction occurred at 500 Ce600 C. Thus, temperature is the decisive factor in the interfacial reaction occurring at this temperature. In other words, the interfacial reaction only occurs when temperature is adequately high (△G ¼ △G0-RTlnax, DG<0). Besides, the magnified XRD plot shown in Fig. 5(a) demonstrates that at a holding temperature that was increased from 500 to 700 C, the diffraction peak of Al was shifted towards the left. This shift can be attributed to the dissolution of elemental Si produced by the decomposition of Ti3SiC2 into the Al matrix, leading to the formation of solute atoms and lattice distortion. As shown in Fig. 5(b), after isothermal treatment at 700 C for 5 and 10 h, the diffraction peak of Al3Ti became very strong, indicating the presence of a large quantity in the matrix. Concurrently, the diffraction peaks of Ti3SiC2 at 40.8 and 42.5 disappeared. The diffraction peak of Al3Ti at 39.5 partially overlapped with the peak at 39.1. As the holding time was prolonged from 1 h to 10 h, the diffraction peak at this position shifted to a smaller diffraction angle. Considering that the other two diffraction peaks of Ti3SiC2 disappeared while the remaining diffraction peaks of Al3Ti strengthened, it can be inferred that Ti3SiC2 completely decomposed after long-term isothermal treatment at a high temperature. As indicated by the XRD plots, it can be inferred that the interfacial reaction between Ti3SiC2 and Al occurred above 600 C. The reaction formula can be simplified as follows:
Table 1 Performance of the sintered composite. Hardness (HV)
Electrical conductivity (%IACS)
Compactness
Coefficient of friction
38
12
96.6%
0.28
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Ti3 SiC2 þ Al/Al3 Ti þ TiC þ AlðSiÞ
Fig. 3. SEM images of the mixed Al/Ti3SiC2 powder after ball milling and sintered composite (a) Ball milled powder; (b) Sintered composite.
(1-1)
Gu et al. [23] prepared Ti3SiC2-reinforced Al matrix composite using hot pressed sintering and studied the characteristics of the interfacial reaction. They found that the main products of the reaction were Al3Ti, Al4SiC4 and Al4C3 at sintering temperatures of 600 Ce650 C and sintering times of 30e180 min. However, since the Ti3SiC2 powder contained a small quantity of TiC, it was difficult to determine whether the TiC was the product of the interfacial reaction. Our results indicated that there was no diffraction peak for TiC in the XRD plot at a temperature of 500 C. However, at 600 C, a weak diffraction peak corresponding to TiC appeared, which was strengthened when the temperature increased to 700 C. This indicates that TiC was the product of the interfacial reaction under this condition, and not due to an impurity arising from the Ti3SiC2 powder. Fig. 6 shows SEM images of the composites after isothermal treatment for 10 h at different temperatures. Most particles were distributed uniformly at the grain boundaries, with an irregular morphology and a dimension less than 5 mm, as indicated by number “1”, “2” in Fig. 6 (a). As temperature increased, the grains increased in size. It can be inferred that the Al element in the matrix diffused into the particles to replace Si, leading to intra-crystalline growth of the particles. Meanwhile, the element Ti migrated continuously, growing into a discrete phase, as shown at position 3 in Fig. 6(c). As the temperature increased, the grains in the composite grew, but insignificantly. This is probably due to a hindering effect on the particles along the grain boundaries. At 700 C, the second phase particles grew deep into one side of the grains, leading to clear, straight edges on this side. This is presumed to be due to phase changes. Fig. 7 shows area scans of the composites after isothermal treatment at 500 C for 10 h. These show that the particle size is approximately 1e15 mm, most particles being distributed uniformly around the grain boundaries. Al matrix occupied the largest area, the grain size varying dramatically, the largest being approximately
Fig. 4. Area scan of the sintered Al matrix composite Fig. 5 XRD plots of the composites after isothermal treatment for different time periods and at different temperatures (a) (500 Ce700 C)/10 h (b) 700 C/(1e10 h).
J. Zhang et al. / Journal of Alloys and Compounds 738 (2018) 1e9
Fig. 5. XRD plots of the composites after isothermal treatment for different time periods and at different temperatures (a) (500 Ce700 C)/10 h (b) 700 C/(1e10 h).
50 mm and the smallest about 10 mm. The distribution regions of the Ti and Si elements nearly overlapped, both being located along the boundaries of the Al matrix grains. Thus, it can be inferred that the Ti3SiC2 particles remain intact and do not decompose. This is consistent with the analysis of XRD plots. Fig. 8 shows a line scan of the particle-reinforced composite after isothermal treatment at 500 C for 10 h. According to the curves of the Ti and Si elements and combined with the results of XRD, it was determined that this particle was Ti3SiC2. There was a high content of Al outside the particles. The closer to the particles where an analysis was performed, the lower the content of Al element. There was a much lower content of Al element inside the particles than Ti and Si. As shown by the curve for Al, no interfacial reactions occurred at 500 C, with Al element not diffusing into the particles. The element Si accounted for a higher proportion inside the particles, indicating an enrichment of Si. This further proves that Ti3SiC2 remained intact and did not decompose or react with the matrix. Fig. 9 presents the SEM/EDS results of the composite after isothermal treatment at 700 C for 10 h and shows a uniform distribution of the second phase particles. Most were located along the grain boundaries, and a few inside the grains. Compared with the composite after isothermal treatment at 500 C, most particles had grown significantly and deeply within the grains. Furthermore, there were also some second phase particles that penetrated the entire grains. The difference in contrast between Al element and
5
matrix decreased, that is, although the contrast of the particles was less than that of the matrix, the interface between the matrix and the reinforcement particles was no longer clear. It could be observed that Al element was already inside the particles. The distribution of Ti element was generally consistent with the position, shape and size of the particles, indicating that it was the main constituent element in the particles. Combined with the XRD plot (Fig. 4(b)), the results suggested that Ti-containing phases in the alloy were mainly Al3Ti and TiC, and the particles were preliminarily determined to be Al3Ti and/or TiC. The distribution pattern of Si element was consistent with that of Ti element. Both elements were enriched near the particles. However, closer observation revealed that the enrichment area of the Si element was much smaller than that of Ti. It can be inferred that during the interfacial reaction, some of the Si element in the periphery of the second phase became free from the control of the particles and it diffused into the Al matrix to form the Al (Si) solid solution. In Lu et al. [24], a large quantity of Cu element diffused into the Ti3SiC2 particles during the reaction, whilst Si atoms in the Ti3SiC2 particles diffused in the opposite direction, leading to the decomposition of Ti3SiC2. Their findings were consistent with ours. Fig. 10 shows the line scan of the composite after isothermal treatment at 700 C for 10 h. According to the curves of the Ti and Si element and combined with the XRD plots, these particles should be Al3Ti. The Al element had a higher content outside the particle, and the content decreased along the boundary between the particle and the Al matrix. The Al content inside the particle was much lower than outside, but it was still higher than the contents of the Ti and Si elements inside the particle. The content of Si element remained low throughout the process, indicating that most elemental Si inside the Ti3SiC2 particle diffused into the Al matrix. This further indicates that Ti3SiC2 decomposed. To determine the main Ti-containing phases after the interfacial reaction, X-ray fluorescence spectrometry was performed on positions 1 and 2 in Fig. 10. The results are shown in Table 2. The proportion of Al to Ti elements in the middle of the lath-shaped particles was 66.3:21.66, close to the Al/Ti ratio of 3:1 in Al3Ti, confirming that the lath-shaped second phase was indeed Al3Ti. In addition, a small quantity of Si element was detected at position 1 in the matrix, indicating the presence of Al-Si solid solution. In addition, Wang et al. [25] prepared Al3Ti/Al matrix composite by an in-situ reaction and observed lath-shaped Al3Ti. Many studies have performed that focused on the interfacial reaction between MAX particles and the metal matrix. V. Dunia [26] and others prepared Ti3SiC2/Cu-based composites by SPS. When the content of Ti3SiC2 was high (18 vol%), the interface reaction between the particles and the matrix was initiated. Moreover, when the content was low (5 vol%), the interface reaction did not occur. The investigators believed that composites in which 18 vol% Ti3SiC2 was added, had a higher intrinsic resistivity and higher hardness, which prevented the formation of intimate contact between the composite particles, resulting in partial melting of the Cu matrix during the SPS sintering at 850 C. The preparation of 20e70 vol% Ti3SiC2/Cu composites was achieved by the hotpressing sintering method as described by Zhou et al. [27]. The sintering temperature was 900 Ce1070 C, the sintering time was 30 min, and the pressure was 40 Mpa. It was found that the interface reaction between particles and matrix occurred at temperatures higher than 900 C, forming the TiCx phase, Cu-Si compound, and the Cu-Si solid solution. Moreover, it was believed that interdiffusion of Si and Cu occurred during the high-temperature bonding and that Cu first diffused into the grain boundary of Ti3SiC2 and subsequently reacted with the Si in Ti3SiC2 to form a Cu(Si) solid solution. In our study, SPS sintering did not cause an obvious interfacial reaction between Ti3SiC2 and the Al matrix.
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Fig. 6. SEM images of the Al matrix composites after isothermal treatment at:(a) 500 C; (b) 600 C; (c) 700 C.
Fig. 7. Area scan of composites after isothermal treatment at 500 C for 10 h.
According to findings presented in a previous study [26], it was concluded that the addition of a lower amount of Ti3SiC2 and the choice of SPS process parameters were potentially the main reasons. However, the long-term heat treatment at a lower melting point temperature (600 C) led to the interfacial reaction. Based on these findings, we can conclude that interfacial reactions can occur
between metals and Ti3SiC2 at temperatures lower than the decomposition temperature of Ti3SiC2 and the melting point of metal matrix. In addition, temperature and time are the key factors that are involved in affecting the occurrence of interfacial reactions. Due to the low temperature and short duration of SPS, the densification of Ti3SiC2/metal-based composites can be achieved without
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obvious interfacial reaction. However, to avoid severe interfacial reactions during sintering, it is warranted to carefully select the content of the MAX phase and process parameters. Furthermore, the interfacial reaction between Ti3SiC2 and metal includes the diffusion of metal atoms into Ti3SiC2 particles and de-intercalation of the Si atom of Ti3SiC2. There is possibly no clear evidence that either of these two processes occurred first or simultaneously. It can therefore be concluded that initially the diffusion of metal atoms into the particles takes place through the fact that “interfacial reaction occurs at a temperature that is far below the Ti3SiC2 decomposition temperature”, however there is still lack of sufficient experimental evidence to proof this statement.
3.3. Performance analysis of the composites
Fig. 8. Line scan of the composite after isothermal treatment at 500 C for 10 h.
3.3.1. Hardness and density Fig. 11 shows the relationship between performance of the Ti3SiC2-reinforced Al matrix composites and the isothermal treatment parameters. When the holding temperature was either 500 C or 600 C, the composite hardness first increased and then decreased. When the holding temperature increased to 700 C, the composite hardness decreased continuously. The higher the holding temperature, the lower the density. Either an increase in holding temperature or prolongation of holding time could promote an interfacial reaction, further leading to a reduction in composite hardness and density. Appropriate interfacial reactions promote the binding of particles to the matrix, which is conducive to load transfer. This manifests as an increase in hardness on the macroscopic scale. As holding time is prolonged, the grains grow. Decomposition of Ti3SiC2 can cause a changes in the stable structure of the composites in addition to a reduction in hardness. The reduction in density may be attributed to the decomposition of Ti3SiC2 that results in the formation of new stable phase, Al3Ti. The density of Al3Ti is 3.36 g/cm3, far lower than that of Ti3SiC2 (4.53 g/ cm3).
Fig. 9. Area scan of the composite after isothermal treatment at 700 C for 10 h.
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Fig. 10. Line scan of the composite after insulation at 600 C for 10 h. Fig. 11. Influence of time on the interfacial reaction on the Vickers hardness of the composites (a) Vickers; (b) Density.
Table 2 Percentage of atoms present at positions 1 and 2 in Fig. 9 (a.t.%). No.
Al
Ti
Si
C
1 2
97.13 66.30
0.27 21.66
2.60 4.67
e 7.37
decomposed and no longer provided a self-lubricating effect. Additionally, the composite hardness decreased, leading to a rapid increase in wear. 4. Conclusions
3.3.2. Coefficient of friction and wear rate Fig. 12 shows the relationship between the coefficient of friction and degree of wear as a function of holding time. As shown in Fig. 12(a), the coefficient of friction first decreased and then increased as holding time was prolonged. The coefficient of friction reached a minimum value of 0.23 after isothermal treatment for 1 h. Thus, appropriate isothermal treatment is conducive to the diffusion and binding of atoms. Ti3SiC2 was still the main second phase, and formed a lubricating film on the friction surface, effectively reducing the coefficient of friction. As holding time was prolonged, the strength of the material decreased as Ti3SiC2 was consumed by the interfacial reaction, which produced Al3Ti. TiC itself does not possess self-lubricating properties, so the coefficient of friction increased. In Fig. 12(b), as holding time was prolonged, the degree of wear of the composite first decreased and then increased, the minimum wear being as low as 0.01 102 g. Thus, appropriate isothermal treatment enhanced the hardness of the composite. The lower the coefficient of friction, the stronger the wear resistance. As holding time was prolonged, Ti3SiC2
Ti3SiC2-reinforced Al matrix composites were prepared by SPS. The influence of holding time and temperature on the interfacial reaction between Ti3SiC2 and Al matrix was studied in the subsequent isothermal treatment process. (1) SPS was used to prepare Ti3SiC2/Al-based composites without significant interface reaction between particles and matrix. When the holding temperature was 500 C, no interfacial reaction occurred, but at a temperature above 600 C, the reaction took place. At a holding temperature of 700 C and a isothermal treatment time above 10 h, Ti3SiC2 completely decomposed. The main products of the interfacial reaction were Al3Ti, TiC and Al (Si) solid solution. (2) During high-temperature long-term isothermal treatment, reinforcement particles in the Ti3SiC2-reinforced Al matrix composite changed from irregular morphology to a lathshaped morphology. As the holding time was prolonged, the density decreased linearly, while the Vickers hardness
J. Zhang et al. / Journal of Alloys and Compounds 738 (2018) 1e9
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References
Fig. 12. Variation in coefficient of friction over time (a) Coefficient of friction; (b) Degree of Wear.
first increased and then decreased. Thus, the holding temperature had a significant impact on hardness, which decreased most rapidly under 700 C. (3) As holding time was prolonged, the coefficient of friction of the composites first decreased and then increased. The minimum coefficient of friction was 0.23. The degree of wear first decreased and then decreased, the minimum being 0.01 102 g.
Acknowledgments This work was supported by National Natural Science Foundation of China (No. 51461017 and 51561008) and the Jiangxi Science and Technology Natural Science Fund(No. 20171ACB21044 and 20161BAB206137).
[1] A.H. Zou, X.L. Zhou, Y.L. Que, et al., Computational and experimental investigation on thermal conductivity of SiC particle reinforced Al-Matrix composite containing pores, Mater. Sci. Forum 762 (2013) 769e776. [2] S. Jerome, B. Ravisankar, P.K. Mahato, et al., Synthesis and evaluation of mechanical and high temperature tribological properties of in-situ AleTiC composites, Tribol. Int. 43 (11) (2010) 2029e2036. [3] M. Rahimian, N. Parvin, N. Ehsani, Investigation of particle size and amount of alumina on microstructure and mechanical properties of Al matrix composite made by powder metallurgy, Mater. Sci. Eng. 527 (4e5) (2010) 1031e1038. [4] A. Canakci, T. Varol, C. Nazik, Effects of amount of methanol on characteristics of mechanically alloyed AleAl2O3 composite powders, Mater. Technol. 27 (4) (2014) 320e327. [5] S. Sobula, E. Olejnik, T. Tokarski, Wear resistance of TiC reinforced cast steel matrix composite, Arch. Foundry Eng. 17 (2017) 143e146. [6] M. Barmouz, P. Asadi, M.K.B. Givi, et al., Investigation of mechanical properties of Cu/SiC composite fabricated by FSP: effect of SiC particles' size and volume fraction, Mater. Sci. Eng. 528 (3) (2011) 1740e1749. [7] M. Hotta, T. Goto, Spark plasma sintering of TiN-Cubic BN composites, J. Ceram. Soc. Jpn. 118 (1374) (2010) 137e140. [8] P. Ravindran, K. Manisekar, R. Narayanasamy, et al., Tribological behaviour of powder metallurgy-processed aluminium hybrid composites with the addition of graphite solid lubricant, Ceram. Int. 39 (2) (2013) 1169e1182. [9] A. Baradeswaran, A.E. Perumal, Wear and mechanical characteristics of Al 7075/graphite composites, Compos. B 56 (1) (2014) 472e476. [10] M.W. Barsoum, M. Radovic, Elastic and mechanical properties of the MAX phases, Annu. Rev. Mater. Res. 41 (41) (2010) 195e227. [11] Yue Zhenming, Yang Limin, Gong Jianhong, et al., Experimental investigation on microstructure and mechanical properties of CBNeTi3SiC2 composites, Adv. Eng. Mater. 18 (9) (2016) 1568e1573. [12] Sun Zhimei, Progress in research and development on MAX phases:a family of layered ternary compounds, Int. Mater. Rev. 56 (3) (2013) 143e166. [13] J. Xiao, T. Yang, C. Wang, et al., Investigations on radiation tolerance of Mnþ1AXn phases: study of Ti3SiC2, Ti3AlC2, Cr2AlC, Cr2GeC, Ti2AlC, and Ti2AlN, J. Am. Ceram. Soc. 98 (4) (2015) 1323e1331. [14] S. Zhao, J. Xue, Y. Wang, et al., Ab initio study of irradiation tolerance for different Mnþ1AXn phases: Ti3SiC2 and Ti3AlC2, J. Appl. Phys. 115 (2) (2014) 201e343. [15] Zhou Yanchun, Sun Zhimei, Electronic structure and bonding properties in layered ternary carbide Ti3SiC2, J. Phys. Condens. Matter 12 (12) (2000) L457eL462. [16] Heng Junjun, T.L. Ngai, Hu Changxun, et al., Effect of Ti3SiC2 content on the property of a warm compacted Cu-Ti3SiC2 composite, Mater. Sci. Forum 628e629 (2009) 471e476. [17] Y.Y. Zhu, S.S. Li, L. Li, et al., Synthesis and decomposition of Ti3SiC2 under 15GPa at 1400 C, Key Eng. Mater. 602e603 (2014) 499e502. [18] S. Li, J. Xie, J. Zhao, et al., Mechanical properties and mechanism of damage tolerance for Ti3SiC2, Mater. Lett. 57 (1) (2002) 119e123. [19] T.L. Ngai, W. Zheng, Y. Li, Effect of sintering temperature on the preparation of CueTi3SiC2 metal matrix composite, Prog. Nat. Sci.: Mater. Int. 23 (1) (2013) 70e76. [20] T.A. Chernyshova, R.S. Mikheev, I.E. Kalashnikov, et al., Development and testing of Al-SiC and Al-TiC composite materials for application in friction units of oil-production equipment, Inorg. Mater. Appl. Res. 2 (3) (2011) 282e289. [21] Sumit, Chabri, Subhranshu, et al., Development and characterization of Al2O3 dispersed Al/Mg/Cu/Ti matrix composite, J. Mater. Sci. Technol. 29 (11) (2013) 1085e1090. [22] W.L. Gu, C.K. Yan, Y.C. Zhou, Reactions between Al and Ti3SiC2 in temperature range of 600e650 C, Scripta Mater. 49 (11) (2003) 1075e1080. [23] J.R. Lu, Y. Zhou, Y. Zheng, et al., Interface structure and wetting behaviour of Cu/Ti3SiC2 system, Adv. Appl. Ceram. 114 (1) (2015) 39e44. [24] Dudina, V. Dina, Bulina, et al., Ti3SiC2-Cu composites by mechanical milling and spark plasma sintering: possible microstructure formation scenarios, Met. Mater. Int. 19 (6) (2013) 1235e1241. [25] Y. Zhou, W. Gu, Chemical reaction and stability of Ti3SiC2 in Cu during hightemperature processing of Cu/Ti3SiC2 composites, Zeitschrift Fur Metallkunde 95 (1) (2004) 50e56. [26] D.V. Dudina, A.K. Mukherjee, Reactive spark plasma sintering: successes and challenges of nanomaterial synthesis, J. Nanomater. 2013 (2013) 12. [27] X. Wang, A. Jha, R. Brydson, In situ fabrication of Al3Ti particle reinforced aluminium alloy metal matrix composites, Mater. Sci. Eng. 364 (1) (2004) 339e345.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Review
Study of the interfacial reaction between Ti3SiC2 particles and Al matrix Jiangbo Zhang b, Wenyang Liu a, Yiming Jin b, Shangjiang Wu a, Taotao Hu b, Yong Li b, Xiangpeng Xiao b, * a b
School of Material Science and Engineering, Jiangxi University of Science and Technology, Ganzhou 341000, China Jiangxi Nonferrous Metal Processing Engineering Technology Research Center, Ganzhou 341000, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 September 2017 Received in revised form 9 December 2017 Accepted 12 December 2017 Available online 14 December 2017
MAX phases have a unique laminated ternary structure and excellent self-lubricating properties. Their tribological properties are similar to that of graphite in addition to having excellent high-temperature oxidation resistance. Because of these features, MAX is an ideal external reinforcement for metalbased self-lubricating composites. Here, Ti3SiC2 powder and pure aluminum (Al) powder were combined as the raw materials to prepare Ti3SiC2/Al-based composite without significant interface reaction between particles and matrix using spark plasma sintering. The degree of relative density of the composite was estimated to be as high as 96.6%. The features of the interfacial reaction between Ti3SiC2 and Al matrix were studied over varying times of isothermal treatment. The results showed that the holding temperature was the key factor in the initiation of the interfacial reaction. No interfacial reaction was observed when the reaction was conducted at 500 C/10 h. At a holding temperature of 600 C, the Si atoms in Ti3SiC2 entered the matrix to form Al(Si) solid solution. Ti then bound to both Al and C to form Al3Ti and TiC. Varying the holding time had a significant impact on the morphology of the reaction products. At temperatures below 700 C, prolongation of the holding time caused particles to gradually eVolve from irregular to lath-shaped. Ti3SiC2 completely decomposed at holding times of 10 h. The performance of the material also changed throughout the process of thermal treatment. As the holding time was prolonged, the density of the material continuously decreased, while the hardness first increased and then decreased. The coefficient of friction and degree of wear first decreased and then increased. The minimum coefficient of friction and wear were 0.23 and 0.01 102 g, respectively. © 2017 Elsevier B.V. All rights reserved.
Keywords: Aluminum matrix composite Frictional wear Interfacial reaction Ti3SiC2
Contents 1. 2. 3.
4.
Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 Experimental procedures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 Results and discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 3.1. Analysis of the composite after sintering densification . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 3.2. Analysis of the composite after isothermal treatment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 3.3. Performance analysis of the composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 3.3.1. Hardness and density . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 3.3.2. Coefficient of friction and wear rate . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8 Acknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9
* Corresponding author. E-mail address: [email protected] (X. Xiao). https://doi.org/10.1016/j.jallcom.2017.12.123 0925-8388/© 2017 Elsevier B.V. All rights reserved.
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1. Introduction The rapid advance of technology and engineering science creates a need for high performance materials. As the use of a single material can no longer meet these requirements, metal matrix composites are being increasingly used because of their excellent, comprehensive performance. Aluminum (Al) matrix composites have the advantage of high specific strength, specific modulus and specific stiffness, and therefore are increasingly utilized in the aviation, automobile and military industries [1e3]. Particles commonly added into the Al matrix are Al2O3 [4], TiC [5], SiC [6] and BN [7], which are hard ceramic particles able to enhance the hardness, and wear and deformation resistance of the Al matrix. However, the properties of these particles are greatly different from those of the Al matrix, which leads to poor plasticity and toughness of the particle-reinforced Al matrix composite, non-uniform distribution of the particles, low binding properties and low secondary process ability [8,9]. Recently, laminated ternary MAX, a novel phase that integrates metallic and ceramic features, has drawn widespread attention [10e12]. MAX refers to a ternary transition metal carbide or nitride, with the general formula Mnþ1AXn where M refers to a transition metal; A is a group IIIA or IVA element; X is C or N, n ¼ 1e3 [13,14]. Ti3SiC2 is representative of MAX compounds. To be specific, Ti3SiC2 displays electrical and thermal conductivity, a high degree of hardness, high elastic modulus and good ductility similar to metals. It also possesses the properties of ceramics, such as thermal stability, high temperature oxidation and heat resistance, creep properties, a low coefficient of friction and excellent selflubricating performance. Therefore, Ti3SiC2 can serve as the ideal reinforced phase for metal matrix compounds [15,16]. Pure Ti3SiC2 does not decompose below 1400 C [17]. The high temperature stability of Ti3SiC2 therefore inhibits an interfacial reaction and preparation of metal matrix composites reinforced by pure particles. However, many studies have shown that when added into a copper matrix, Ti3SiC2 particles will decompose at a much lower temperature, that is, a reduction from an initial 1400 Ce900 C [18,19]. This leads to a severe interfacial reaction and the formation of complex products, depending on the temperature and proportions of the components. In general, the addition of Ti3SiC2 into a copper matrix is a costly approach [20,21]. An Al matrix has the features of low density, high specific strength, high specific stiffness and corrosion resistance. When added to Ti3SiC2, Al matrix composites acquire excellent self-lubricating properties. However, the interfacial reaction between Ti3SiC2 and the Al matrix is the major factor that determines the performance of Al matrix composites. The affinity between Al and O2 is relatively high, and oxidation is easily achieved, which interferes with the study of interfacial reactions, and does not favor the preparation of composites with a high density. Dunia [22] et al. suggested that the dynamic vacuum of the SPS chamber creates a protective environment that is beneficial for oxidation-sensitive materials. In this study, we used SPS sintering preparation of high-density Ti3SiC2/Al composites without internal oxidation and obvious interfacial reaction, to investigate the interfacial reactions between Ti3SiC2 particles and the Al matrix by isothermal treatment, and evaluated the effects of the interfacial reaction on the mechanical properties of the composites. 2. Experimental procedures The raw materials were 300-mesh pure Al powder and 1000mesh Ti3SiC2 powder (see Table in Fig. 1 for particle size determination). After premixing, alcohol was added to a PMQW series Omnibearing planetary ball mill for wet ball milling for 9 h. The ratio of grinding media to material was 11:3 under argon
protection, at a milling rate of 200 r/min. After complete drying and cold molding, sintering was performed in a ST-5 SPS machine for 10 min at a temperature of 600 C and pressure of 40 MPa, followed by isothermal treatment in a muffle furnace at either 500 C, 600 C or 700 C for 1 h, 3 h, 5 h or 10 h. A small muffle furnace was used. The lining material of the furnace was silicon carbide and it operated with resistance heating. Material density was measured according to standard GB/T 1423-1996 “Method of measurement of density for precious metals and their alloys”. Hardness was measured using a MC010-HV-1000 Vickers hardness tester under a load of 0.2 N and holding time of 15 s. The mean of five measurements was taken as the final value. Frictional wear was determined on the M-2000 type wear tester. A block-on-ring test arrangement was used with a friction pair made from 45Cr steel with a hardness of 60 HRC and diameter of 20 mm. The test was performed for 600 s at room temperature. The experimental force was set to 50 N, and friction velocity at 200 r/ min. Each specimen was subjected to coarse grinding, fine grinding and polishing, then corroded with 0.5% HF for approximately 30s. The metallurgical structure was evaluated using a CMM-772 light microscope. Measurement of the phase of the composite was achieved using a PANalytical Axios Max X-ray fluorescence spectrometer. The microscopic structure and micro-area composition were analyzed using an FEI ML650F scanning electron microscope (SEM) and Bruker energy dispersive spectrometer, respectively. 3. Results and discussion 3.1. Analysis of the composite after sintering densification Fig. 2 shows the XRD patterns of the mixed Ti3SiC2/Al powder and Ti3SiC2-reinforced Al matrix composite, whose three diffraction peaks between 30 and 45 are located at 39.5 , 40.8 and 42.5 . The phases of the ball milled powder and the sintered composite correspond to the diffraction peaks of Al and Ti3SiC2, with no peaks from impurities. This indicated that wet ball milling and SPS technology can prepare pure Ti3SiC2-reinforced Al matrix composite. According to Table 1, the comprehensive performance of the sintered Ti3SiC2-reinforced Al matrix composite is very satisfactory, with compactness reaching up to 96.6% and the coefficient of friction as low as 0.28. Fig. 3 shows SEM images of mixed Al and Ti3SiC2 powders after ball milling and the sintered composite. In Fig. 3(a), the large gray particles are Al, with a diameter of 10e80 mm. The white particles of 1e5 mm are uniformly distributed around the large particles. Fig. 3(b) shows the sintered composite, in which fine round particles 1e8 mm in diameter are uniformly distributed and predominate at the grain boundaries. Comparison shows that after SPS, the diameter of the Ti3SiC2 particles greatly increased and were more likely to cluster at the triangular grain boundaries. This is possibly because the poor wet-ability between the ceramic particles and metal led to the failure of the matrix grains to engulf the reinforcement particles. As the Ti3SiC2 particles grew along the sintering neck, they were gradually pushed to the triangular grain boundaries and clustered together. Fig. 4 shows the surface scanning results of Ti3SiC2-reinforced Al matrix composite after SPS at 600 C. As can be seen in Fig. 4(a), the Al matrix takes up a greater proportion of the cross-section, with grain boundaries of elemental Al showing as black. The particles are generally located along the black grain boundaries, with a distinct interface observed between the reinforcement particles and the matrix. The distribution regions of the elements Ti and Si nearly overlap, both located in the black grain boundaries of the Al matrix grains. This further indicates that Ti3SiC2 does not decompose during SPS.
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3
Fig. 1. Particle size determination of Ti3SiC2.
Fig. 2. XRD patterns of the Al/Ti3SiC2 ball milled powder and sintered composite.
3.2. Analysis of the composite after isothermal treatment Fig. 5 shows the XRD patterns of the sintered Al matrix composites after isothermal treatment for 10 h at different temperatures (500 Ce700 C) and at 700 C for different holding times (1e0 h). In Fig. 5(a), no new phases were produced in the composite after long-term isothermal treatment under 500 C, with Al and Ti3SiC2 still being the major phases. This indicates no apparent interfacial reaction occurring at this temperature. After the
incubating at 600 C for 10 h, the strong diffraction peak of Al and weak diffraction peak of Ti3SiC2 appear, together with the diffraction peaks of new phases. These new phases were identified as Al3Ti and TiC, indicating the occurrence of the interfacial reaction. It is thus inferred that the interfacial reaction occurred at 500 Ce600 C. Thus, temperature is the decisive factor in the interfacial reaction occurring at this temperature. In other words, the interfacial reaction only occurs when temperature is adequately high (△G ¼ △G0-RTlnax, DG<0). Besides, the magnified XRD plot shown in Fig. 5(a) demonstrates that at a holding temperature that was increased from 500 to 700 C, the diffraction peak of Al was shifted towards the left. This shift can be attributed to the dissolution of elemental Si produced by the decomposition of Ti3SiC2 into the Al matrix, leading to the formation of solute atoms and lattice distortion. As shown in Fig. 5(b), after isothermal treatment at 700 C for 5 and 10 h, the diffraction peak of Al3Ti became very strong, indicating the presence of a large quantity in the matrix. Concurrently, the diffraction peaks of Ti3SiC2 at 40.8 and 42.5 disappeared. The diffraction peak of Al3Ti at 39.5 partially overlapped with the peak at 39.1. As the holding time was prolonged from 1 h to 10 h, the diffraction peak at this position shifted to a smaller diffraction angle. Considering that the other two diffraction peaks of Ti3SiC2 disappeared while the remaining diffraction peaks of Al3Ti strengthened, it can be inferred that Ti3SiC2 completely decomposed after long-term isothermal treatment at a high temperature. As indicated by the XRD plots, it can be inferred that the interfacial reaction between Ti3SiC2 and Al occurred above 600 C. The reaction formula can be simplified as follows:
Table 1 Performance of the sintered composite. Hardness (HV)
Electrical conductivity (%IACS)
Compactness
Coefficient of friction
38
12
96.6%
0.28
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Ti3 SiC2 þ Al/Al3 Ti þ TiC þ AlðSiÞ
Fig. 3. SEM images of the mixed Al/Ti3SiC2 powder after ball milling and sintered composite (a) Ball milled powder; (b) Sintered composite.
(1-1)
Gu et al. [23] prepared Ti3SiC2-reinforced Al matrix composite using hot pressed sintering and studied the characteristics of the interfacial reaction. They found that the main products of the reaction were Al3Ti, Al4SiC4 and Al4C3 at sintering temperatures of 600 Ce650 C and sintering times of 30e180 min. However, since the Ti3SiC2 powder contained a small quantity of TiC, it was difficult to determine whether the TiC was the product of the interfacial reaction. Our results indicated that there was no diffraction peak for TiC in the XRD plot at a temperature of 500 C. However, at 600 C, a weak diffraction peak corresponding to TiC appeared, which was strengthened when the temperature increased to 700 C. This indicates that TiC was the product of the interfacial reaction under this condition, and not due to an impurity arising from the Ti3SiC2 powder. Fig. 6 shows SEM images of the composites after isothermal treatment for 10 h at different temperatures. Most particles were distributed uniformly at the grain boundaries, with an irregular morphology and a dimension less than 5 mm, as indicated by number “1”, “2” in Fig. 6 (a). As temperature increased, the grains increased in size. It can be inferred that the Al element in the matrix diffused into the particles to replace Si, leading to intra-crystalline growth of the particles. Meanwhile, the element Ti migrated continuously, growing into a discrete phase, as shown at position 3 in Fig. 6(c). As the temperature increased, the grains in the composite grew, but insignificantly. This is probably due to a hindering effect on the particles along the grain boundaries. At 700 C, the second phase particles grew deep into one side of the grains, leading to clear, straight edges on this side. This is presumed to be due to phase changes. Fig. 7 shows area scans of the composites after isothermal treatment at 500 C for 10 h. These show that the particle size is approximately 1e15 mm, most particles being distributed uniformly around the grain boundaries. Al matrix occupied the largest area, the grain size varying dramatically, the largest being approximately
Fig. 4. Area scan of the sintered Al matrix composite Fig. 5 XRD plots of the composites after isothermal treatment for different time periods and at different temperatures (a) (500 Ce700 C)/10 h (b) 700 C/(1e10 h).
J. Zhang et al. / Journal of Alloys and Compounds 738 (2018) 1e9
Fig. 5. XRD plots of the composites after isothermal treatment for different time periods and at different temperatures (a) (500 Ce700 C)/10 h (b) 700 C/(1e10 h).
50 mm and the smallest about 10 mm. The distribution regions of the Ti and Si elements nearly overlapped, both being located along the boundaries of the Al matrix grains. Thus, it can be inferred that the Ti3SiC2 particles remain intact and do not decompose. This is consistent with the analysis of XRD plots. Fig. 8 shows a line scan of the particle-reinforced composite after isothermal treatment at 500 C for 10 h. According to the curves of the Ti and Si elements and combined with the results of XRD, it was determined that this particle was Ti3SiC2. There was a high content of Al outside the particles. The closer to the particles where an analysis was performed, the lower the content of Al element. There was a much lower content of Al element inside the particles than Ti and Si. As shown by the curve for Al, no interfacial reactions occurred at 500 C, with Al element not diffusing into the particles. The element Si accounted for a higher proportion inside the particles, indicating an enrichment of Si. This further proves that Ti3SiC2 remained intact and did not decompose or react with the matrix. Fig. 9 presents the SEM/EDS results of the composite after isothermal treatment at 700 C for 10 h and shows a uniform distribution of the second phase particles. Most were located along the grain boundaries, and a few inside the grains. Compared with the composite after isothermal treatment at 500 C, most particles had grown significantly and deeply within the grains. Furthermore, there were also some second phase particles that penetrated the entire grains. The difference in contrast between Al element and
5
matrix decreased, that is, although the contrast of the particles was less than that of the matrix, the interface between the matrix and the reinforcement particles was no longer clear. It could be observed that Al element was already inside the particles. The distribution of Ti element was generally consistent with the position, shape and size of the particles, indicating that it was the main constituent element in the particles. Combined with the XRD plot (Fig. 4(b)), the results suggested that Ti-containing phases in the alloy were mainly Al3Ti and TiC, and the particles were preliminarily determined to be Al3Ti and/or TiC. The distribution pattern of Si element was consistent with that of Ti element. Both elements were enriched near the particles. However, closer observation revealed that the enrichment area of the Si element was much smaller than that of Ti. It can be inferred that during the interfacial reaction, some of the Si element in the periphery of the second phase became free from the control of the particles and it diffused into the Al matrix to form the Al (Si) solid solution. In Lu et al. [24], a large quantity of Cu element diffused into the Ti3SiC2 particles during the reaction, whilst Si atoms in the Ti3SiC2 particles diffused in the opposite direction, leading to the decomposition of Ti3SiC2. Their findings were consistent with ours. Fig. 10 shows the line scan of the composite after isothermal treatment at 700 C for 10 h. According to the curves of the Ti and Si element and combined with the XRD plots, these particles should be Al3Ti. The Al element had a higher content outside the particle, and the content decreased along the boundary between the particle and the Al matrix. The Al content inside the particle was much lower than outside, but it was still higher than the contents of the Ti and Si elements inside the particle. The content of Si element remained low throughout the process, indicating that most elemental Si inside the Ti3SiC2 particle diffused into the Al matrix. This further indicates that Ti3SiC2 decomposed. To determine the main Ti-containing phases after the interfacial reaction, X-ray fluorescence spectrometry was performed on positions 1 and 2 in Fig. 10. The results are shown in Table 2. The proportion of Al to Ti elements in the middle of the lath-shaped particles was 66.3:21.66, close to the Al/Ti ratio of 3:1 in Al3Ti, confirming that the lath-shaped second phase was indeed Al3Ti. In addition, a small quantity of Si element was detected at position 1 in the matrix, indicating the presence of Al-Si solid solution. In addition, Wang et al. [25] prepared Al3Ti/Al matrix composite by an in-situ reaction and observed lath-shaped Al3Ti. Many studies have performed that focused on the interfacial reaction between MAX particles and the metal matrix. V. Dunia [26] and others prepared Ti3SiC2/Cu-based composites by SPS. When the content of Ti3SiC2 was high (18 vol%), the interface reaction between the particles and the matrix was initiated. Moreover, when the content was low (5 vol%), the interface reaction did not occur. The investigators believed that composites in which 18 vol% Ti3SiC2 was added, had a higher intrinsic resistivity and higher hardness, which prevented the formation of intimate contact between the composite particles, resulting in partial melting of the Cu matrix during the SPS sintering at 850 C. The preparation of 20e70 vol% Ti3SiC2/Cu composites was achieved by the hotpressing sintering method as described by Zhou et al. [27]. The sintering temperature was 900 Ce1070 C, the sintering time was 30 min, and the pressure was 40 Mpa. It was found that the interface reaction between particles and matrix occurred at temperatures higher than 900 C, forming the TiCx phase, Cu-Si compound, and the Cu-Si solid solution. Moreover, it was believed that interdiffusion of Si and Cu occurred during the high-temperature bonding and that Cu first diffused into the grain boundary of Ti3SiC2 and subsequently reacted with the Si in Ti3SiC2 to form a Cu(Si) solid solution. In our study, SPS sintering did not cause an obvious interfacial reaction between Ti3SiC2 and the Al matrix.
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Fig. 6. SEM images of the Al matrix composites after isothermal treatment at:(a) 500 C; (b) 600 C; (c) 700 C.
Fig. 7. Area scan of composites after isothermal treatment at 500 C for 10 h.
According to findings presented in a previous study [26], it was concluded that the addition of a lower amount of Ti3SiC2 and the choice of SPS process parameters were potentially the main reasons. However, the long-term heat treatment at a lower melting point temperature (600 C) led to the interfacial reaction. Based on these findings, we can conclude that interfacial reactions can occur
between metals and Ti3SiC2 at temperatures lower than the decomposition temperature of Ti3SiC2 and the melting point of metal matrix. In addition, temperature and time are the key factors that are involved in affecting the occurrence of interfacial reactions. Due to the low temperature and short duration of SPS, the densification of Ti3SiC2/metal-based composites can be achieved without
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obvious interfacial reaction. However, to avoid severe interfacial reactions during sintering, it is warranted to carefully select the content of the MAX phase and process parameters. Furthermore, the interfacial reaction between Ti3SiC2 and metal includes the diffusion of metal atoms into Ti3SiC2 particles and de-intercalation of the Si atom of Ti3SiC2. There is possibly no clear evidence that either of these two processes occurred first or simultaneously. It can therefore be concluded that initially the diffusion of metal atoms into the particles takes place through the fact that “interfacial reaction occurs at a temperature that is far below the Ti3SiC2 decomposition temperature”, however there is still lack of sufficient experimental evidence to proof this statement.
3.3. Performance analysis of the composites
Fig. 8. Line scan of the composite after isothermal treatment at 500 C for 10 h.
3.3.1. Hardness and density Fig. 11 shows the relationship between performance of the Ti3SiC2-reinforced Al matrix composites and the isothermal treatment parameters. When the holding temperature was either 500 C or 600 C, the composite hardness first increased and then decreased. When the holding temperature increased to 700 C, the composite hardness decreased continuously. The higher the holding temperature, the lower the density. Either an increase in holding temperature or prolongation of holding time could promote an interfacial reaction, further leading to a reduction in composite hardness and density. Appropriate interfacial reactions promote the binding of particles to the matrix, which is conducive to load transfer. This manifests as an increase in hardness on the macroscopic scale. As holding time is prolonged, the grains grow. Decomposition of Ti3SiC2 can cause a changes in the stable structure of the composites in addition to a reduction in hardness. The reduction in density may be attributed to the decomposition of Ti3SiC2 that results in the formation of new stable phase, Al3Ti. The density of Al3Ti is 3.36 g/cm3, far lower than that of Ti3SiC2 (4.53 g/ cm3).
Fig. 9. Area scan of the composite after isothermal treatment at 700 C for 10 h.
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Fig. 10. Line scan of the composite after insulation at 600 C for 10 h. Fig. 11. Influence of time on the interfacial reaction on the Vickers hardness of the composites (a) Vickers; (b) Density.
Table 2 Percentage of atoms present at positions 1 and 2 in Fig. 9 (a.t.%). No.
Al
Ti
Si
C
1 2
97.13 66.30
0.27 21.66
2.60 4.67
e 7.37
decomposed and no longer provided a self-lubricating effect. Additionally, the composite hardness decreased, leading to a rapid increase in wear. 4. Conclusions
3.3.2. Coefficient of friction and wear rate Fig. 12 shows the relationship between the coefficient of friction and degree of wear as a function of holding time. As shown in Fig. 12(a), the coefficient of friction first decreased and then increased as holding time was prolonged. The coefficient of friction reached a minimum value of 0.23 after isothermal treatment for 1 h. Thus, appropriate isothermal treatment is conducive to the diffusion and binding of atoms. Ti3SiC2 was still the main second phase, and formed a lubricating film on the friction surface, effectively reducing the coefficient of friction. As holding time was prolonged, the strength of the material decreased as Ti3SiC2 was consumed by the interfacial reaction, which produced Al3Ti. TiC itself does not possess self-lubricating properties, so the coefficient of friction increased. In Fig. 12(b), as holding time was prolonged, the degree of wear of the composite first decreased and then increased, the minimum wear being as low as 0.01 102 g. Thus, appropriate isothermal treatment enhanced the hardness of the composite. The lower the coefficient of friction, the stronger the wear resistance. As holding time was prolonged, Ti3SiC2
Ti3SiC2-reinforced Al matrix composites were prepared by SPS. The influence of holding time and temperature on the interfacial reaction between Ti3SiC2 and Al matrix was studied in the subsequent isothermal treatment process. (1) SPS was used to prepare Ti3SiC2/Al-based composites without significant interface reaction between particles and matrix. When the holding temperature was 500 C, no interfacial reaction occurred, but at a temperature above 600 C, the reaction took place. At a holding temperature of 700 C and a isothermal treatment time above 10 h, Ti3SiC2 completely decomposed. The main products of the interfacial reaction were Al3Ti, TiC and Al (Si) solid solution. (2) During high-temperature long-term isothermal treatment, reinforcement particles in the Ti3SiC2-reinforced Al matrix composite changed from irregular morphology to a lathshaped morphology. As the holding time was prolonged, the density decreased linearly, while the Vickers hardness
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References
Fig. 12. Variation in coefficient of friction over time (a) Coefficient of friction; (b) Degree of Wear.
first increased and then decreased. Thus, the holding temperature had a significant impact on hardness, which decreased most rapidly under 700 C. (3) As holding time was prolonged, the coefficient of friction of the composites first decreased and then increased. The minimum coefficient of friction was 0.23. The degree of wear first decreased and then decreased, the minimum being 0.01 102 g.
Acknowledgments This work was supported by National Natural Science Foundation of China (No. 51461017 and 51561008) and the Jiangxi Science and Technology Natural Science Fund(No. 20171ACB21044 and 20161BAB206137).
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