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Evolution of as-cast and heat-treated microstructure of a commercial bearing alloy
Materials Science and Engineering A 398 (2005) 332–343
Evolution of as-cast and heat-treated microstructure of a commercial bearing alloy Pritha Choudhury, Karabi Das, Siddhartha Das ∗ Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721302, India Received 24 November 2004; received in revised form 17 March 2005; accepted 29 March 2005
Abstract The present study involves the microstructural changes of a zinc–aluminium alloy (ZA27) by nickel, a high melting point element, and the commercial master alloys, Al–5Ti–1B and Al–5Sr. The evolution of as-cast microstructure and the effect of heat treatment on the as-cast microstructure and hardness have been presented in detail. The addition of Ti, B and Sr as master alloys results in significant refinement in the dendrite structure of the ZA27 alloy. The addition of Ni to ZA27-based alloy results in the formation of Ni-aluminides as rod as well as blocky-shaped particles. High amount of Ni (2 wt.%) refines the dendritic structure significantly. © 2005 Elsevier B.V. All rights reserved. Keywords: ZA27; Dendrite structure; Bearing allloy
1. Introduction Zinc–aluminium (ZA) alloys, with a unique combination of properties, are competitive alternative materials to most aluminium casting alloys, bearing bronze, cast iron, as well as to plastics and steel fabrications [1–5]. ZA alloys are a group of high-performance, high-aluminium zinc alloys with excellent bearing properties [3]. Some of the attractive properties of the Zn–Al casting alloys are high sliding wear resistance (usefulness as bearings and bushings) [3], good machinability, excellent corrosion resistance in a variety of environments [6] and high damping capacity (suitability in defence applications) [7]. Addition of some alloying elements is known to have significant effect on the microstructure of these alloys [8,9]. High melting point element like Ni is reported to improve the sliding wear properties of the Zn–Al casting alloys [2,10]. However, literature pertaining to the systematic evolution of microstructure in such alloys is lacking. In this paper, it is aimed to present systematically the evolution of microstructure in a high-aluminium zinc-based alloy (ZA27) to which Ti, B, and Sr have been added in the form of commercial ∗
Corresponding author. E-mail address: [email protected] (S. Das).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.03.098
master alloys, Al–5Ti–1B and Al–5Sr. Such master alloys are generally used for the grain refinement and/or modification of some aluminium alloys. However, the use of grain refiners for the zinc–aluminium based alloys is not a common practice. The effect of Ni on the microstructure of ZA27 is also studied.
2. Experimental procedure All alloys based on ZA27 composition were prepared by conventional melting and casting route. They are identified as ZA1, ZTiB1, ZSr1, ZNi1, ZNi3 and ZNi4 for the sake of brevity and their compositions are shown in Table 1. Please note that ZA1 is the base ZA27 alloy. The alloys ZNi1, ZNi3 and ZNi4 contain 0.3, 0.9 and 2.0 wt.% Ni, respectively. For microstructural characterization, samples were metallographically polished according to standard practices and etched suitably using diluted nitric acid (5 vol.%) in water. The metallographically polished samples were subjected to image analysis using the Leica QWin image analysis software attached with a Leica DMRX optical microscope. Scanning electron microscopy was carried out in a JEOL scanning electron microscope (SEM) attached with an energy dispersive X-ray spectroscopic (EDS) facility. For transmission electron
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Table 1 Chemical composition of the alloys Alloys
Zn (wt.%)
Al (wt.%)
Cu (wt.%)
Mg (wt.%)
Ti (wt.%)
B (wt.%)
Sr (wt.%)
Ni (wt.%)
ZA1 ZTiB1 ZSr1 ZNi1 ZNi3 ZNi4 Al–5Ti–1B Al–5Sr
69.95 69.85 71.64 71.89 70.32 69.20 – –
28.0 27.72 26.0 26.2 26.0 26.8 94.0 95.0
2.0 2.25 2.20 1.56 1.58 0.80 – –
0.05 0.05 0.05 0.05 0.05 0.05 – –
– 0.1 – – – – 5 –
– 0.02 – – – – 1.00 –
– – 0.1 – – – – 5.0
– – – 0.3 0.9 2.0 – –
microscopy, thin foils with a thickness less than 100 m were prepared by dry grinding on silicon carbide papers. Discs of 3 mm diameter were punched from the thin foils. Final thinning of the discs was performed by electropolishing in a Struers Tenupol-5 double-jet electropolisher using 8% perchloric acid in 92% methanol as electrolyte at a voltage of 11.5 V and a temperature of −30 to −20 ◦ C. The thin foils were observed in Philips CM12 transmission electron microscope at an operating voltage of 120 kV. Bulk hardness of all the alloys were measured using a Brinell hardness tester, with a 10 mm diameter steel-ball indenter. The load applied was 500 kg for 60 s.
the (Al) solid solution undergoes a eutectoid reaction to form a mixture of ␣Al + (Zn). Accordingly, the base alloy, ZA1 comprises of primary ␣ dendrites surrounded by the eutectoid ␣Al + (Zn) as shown in Fig. 1. 3.1.2. Alloy ZTiB1 The optical and transmission electron (TEM) micrographs of ZTiB1 are shown in Fig. 2. The sequence of formation of
3. Results and discussion 3.1. Microstructure 3.1.1. Alloy ZA1 The optical micrograph of ZA1 is shown in Fig. 1. Zn–Al binary system is principally a eutectic system, consisting of a eutectoid reaction and a miscibility gap in the solid state having the critical point at 351.5 ◦ C. The ZA1 alloy fall in the hypereutectoid range of the Zn–Al phase diagram [11]. The phase to form first upon solidification is the primary ␣ (the solid-solution of zinc in aluminium). In the process, the excess zinc is rejected to the surrounding liquid. Subsequently,
Fig. 1. Optical micrograph of as-cast ZA1.
Fig. 2. (a) Optical and (b) TEM micrographs of as-cast ZTiB1.
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Fig. 3. (a) Optical and (b)–(d) TEM micrographs of as-cast ZSr1.
the Al-rich and Zn-rich phases are presumably same as in ZA1, already described. However, it is observed from the optical micrograph that the dendrites are rosette-shaped and much smaller in size than those in the unmodified alloy, i.e., ZA1 alloy. The lamellar structure of the eutectoid constituents is observed in the TEM micrographs (Fig. 2(b)). The average width of the lamellae is calculated to be 0.02 m. The lamellae perpendicular to the surface have a cellular appearance (Fig. 2(b)). Large black regions are titanium aluminide precipitates formed in the molten state. When the master alloy, Al–5Ti–1B is added to the molten alloy, the existing Ti-aluminides have enough time for growth during solidification and subsequent cooling, resulting in large particles (0.06–0.08 m). 3.1.3. ZSr1 The optical and TEM photographs of as-cast ZSr1 are shown in Fig. 3. It is observed from the optical micrographs that the addition of Sr in the form of Al–5Sr master alloy increases the curvature of the dendrites.
Fig. 3(b) shows a typical lamellar morphology of the eutectoid constituents. The average width of the eutectoid lamellae in this alloy is calculated to be 0.08 m. The fine precipitates (average size 0.03 m) observed in Fig. 3(c) are formed in the solid state and hence did not have enough time for growth. These precipitates are most likely those of AlCu or Al4 Cu9 (details given in a later section). Anti phase domain boundaries (APB) have been observed in this alloy (as marked by an arrow in Fig. 3(d)). During the fast cooling of the alloy, a number of small ordered regions have been formed rapidly. These ordered regions are separated from each other by the APBs, which are formed in the solid state. The ordered regions are most likely those of the intermetallic Al4 Sr, which was pre-existing in the master alloy and has been detected in as-cast ZSr1 by X-ray diffraction (described in a later section). 3.1.4. ZNi1, ZNi3 and ZNi4 Figs. 4–6 show the optical and SEM micrographs of the alloy ZNi1, ZNi3 and ZNi4. The sequence of formation of the Zn-rich and Al-rich phases are presumably same as in ZA1.
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Fig. 4. (a) Optical and (b) SEM micrographs of as-cast ZNi1: arrow indicates nickel-aluminide (Al3 Ni) particle.
Fig. 5. (a) Optical and (b) SEM micrographs of as-cast ZNi3: arrow indicates nickel-aluminide (Al3 Ni) particle.
Fig. 6. (a) Optical and (b) SEM micrographs of as-cast ZNi4: arrow indicates nickel-aluminide (Al3 Ni) particle.
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Fig. 7. Transmission electron micrograph (TEM) of as-cast ZNi3.
The particles of Ni-aluminides are observed in the optical as well as SEM micrographs. With increasing Ni, the morphology of the particles changes from blocky to rod-shape. The particles are blocky, mostly rod with some blocky and rodshaped in alloy ZNi1, ZNi3 and ZNi4, respectively. The Ni that has been added in powder form remains in the solid state during melting of the alloy (∼680 ◦ C). Hence, the formation of the Ni-aluminides occurs in the solid state. It is clearly evident from optical micrographs (Fig. 6(a)) that a significant refinement of dendrites, is achieved in ZNi4. The lamellar structure of the eutectoid constituents are clearly visible in Fig. 7. The fine precipitates (0.01–0.04 m) seen inside the eutectoid lamellae did not have enough time for growth and hence, are very small in size. The shape of these particles is nearly spherical which also suggests that they are in the initial stages of formation. These fine precipitates are likely to be those of AlCu and Al4 Cu9 and they formed in the solid state. The bigger black regions represent the precipitates formed in the liquid state. These precipitates formed on pre-existing particles and having longer time for growth they became larger in size (0.08–0.11 m in length). These particles are most likely those of the Ni-aluminides which started growing on the Ni particles added in powder form during melting. 3.2. Phase identification 3.2.1. ZA1 The major phases observed in ZA1 are shown in the X-ray diffractogram of Fig. 8. The alloys are cast in metal mold and hence subjected to rapid cooling. According to the low temperature (200 and 20 ◦ C) isothermal sections of Zn–Al–Cu [12,13], AlCu and CuZn4 are the phases existing in equilibrium with Al- and Zn-rich phases at respective temperatures. Al4 Cu9 is a phase found in the Cu-rich side of the isothermal section of Zn–Al–Cu ternary system at 20 ◦ C and also at the higher temperature of 350 ◦ C [14]. Since elemental Cu is added to the Zn–Al melt, Al4 Cu9 phase forms at
Fig. 8. X-ray diffractogram of as-cast ZA1 (Co target).
the solid Cu-melt interface. Owing to the rapid cooling and solidification, there is insufficient time for the complete decomposition of Al4 Cu9 into AlCu and hence some amount of Al4 Cu9 still remains in the as-cast state as observed in the XRD pattern. Other phases AlCu and CuZn4 are observed in the XRD pattern. The AlCu phase is not supposed to exist in the 20 ◦ C isothermal section. However, it is observed in the as-cast microstructure owing to a rapid cooling associated with casting in a metal mold. Mg present in minor quantities is added in order to strengthen and increase corrosion resistance of the alloy. Mg is present as the intermetallics AlMg and Mg2 Zn11 (Fig. 8) and the remaining amount goes into solid solution with the Al-rich phase. From the isothermal section of the Al–Mg–Zn system at 25 ◦ C [15], it is observed that Mg2 Zn11 may be present in ZA1. However, owing to rapid cooling, AlMg of the Mg-rich end does not get enough time for decomposition, and, hence, observed in the as-cast alloy (Fig. 8). However, one should note that peaks for AlMg and Mg2 Zn11 overlap with those of Zn. 3.2.2. ZTiB1 The major phases in the Al–5Ti–1B master alloy are shown in the X-ray diffractogram in Fig. 9. In the Al–Ti binary system [11], Al – 5 wt.% Ti is a hypoperitectic alloy from the Ti-end. It is observed from the isothermal section at 650 ◦ C of Al–Ti–B ternary system [16] that the alloy Al–5Ti–1B consists of Al3 Ti, TiB2 and the Al-rich solid solution. The phases Al3 Ti and (Al) are observed in Fig. 9. Thus, Al3 Ti is the major Ti-aluminide phase as obtained from the binary Al–Ti and the ternary Al–Ti–B phase diagrams. In Al–5Ti–1B, the volume percent of TiB2 is empirically calculated to be ∼0.714, which is far too less for the occurrence of any peak in the X-ray diffractogram. However, one small peak of very low intensity corresponding to that of TiB2 (Fig. 9) coincides with Al (the major phase), and therefore may not be considered.
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Fig. 9. X-ray diffractograms of master alloys, Al–5Ti–1B and Al–5Sr (Cu target).
ZTiB1 contains the phases present in the base alloy, ZA1, as well as additional phases containing Ti (Fig. 10). The phases Al3 Ti and AlTi are already present in the master alloy as stated earlier. These intermetallic phases having very high melting points remain undissolved in the Zn-based alloy during melting which was carried out at ∼680 ◦ C. The amount of Ti added is only 0.1 wt.% and hence it results in a very small volume fraction of the Ti-containing intermetallic phases. Generally peaks of Ti-containing intermetallics coincide with any other phase like Zn or Al. The phase Al24 Ti8 is an ordered superstructure of Al3 Ti. The other phases remain the same as that in the unmodified ZA27 alloy.
Fig. 10. X-ray diffractograms of as-cast ZTiB1 and ZSr1 (Co target).
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3.2.3. ZSr1 It is observed from the Al–Sr binary phase diagram [10] that the master alloy Al–5Sr (5 wt.% Sr) is a hypereutectic alloy of Al and Al4 Sr from the Al end. Al has negligible solid solubility for Sr. In addition to Al4 Sr, the master alloy also contains negligible amount of AlSr as shown in Fig. 9. The composition of AlSr closely matches that of Al7 Sr8 [11]. The intermetallic phases, Al4 Sr and AlSr (high melting points) of the master alloy remain undissolved during melting of the alloy ZSr1. Since the amount of Sr added is too small (0.1 wt.%), the volume fraction of the resultant Sr-containing intermetallic phases is also too less for the appearance of any peak in the X-ray diffractogram (Fig. 10). Hence, the peaks of Al4 Sr and AlSr having significant intensity are the ones coinciding with the peaks of the major phase Zn. The other phases present are those also found in the base alloy (ZA1). The APBs observed in the TEM of this alloy (Fig. 3(d)) formed in ordered Al4 Sr phase. 3.2.4. ZNi1, ZNi3 and ZNi4 The X-ray diffractograms of the alloys containing Ni are shown in Fig. 11. It is seen from this figure that the phases present in addition to those of ZA1 are those of the Nicontaining intermetallics, like, Al3 Ni, AlNi, Ni3 Al, Ni3 Zn22 and Ni5 Zn21 . According to the Al–Ni binary phase diagram [11], in the ZA27 alloys containing Ni, Al3 Ni presumably is the Ni-aluminide phase existing in equilibrium with the other phases. During the rapid cooling of the alloys, there is not enough time for diffusion at the Ni–Al interfaces. This could have resulted in the formation of the other Ni-aluminide phases already mentioned. The amount of Ni3 Al, NiAl or Al3 Ni2 formed in any of these alloys is too less for the appearance of any peak of reasonable intensity in the X-ray diffractogram. Hence, only those peaks of Ni3 Al, NiAl or Al3 Ni2 coinciding with any of the major phases like Zn or
Fig. 11. X-ray diffractograms of as-cast ZNi1, ZNi3 and ZNi4 (Co target).
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Table 2 Hardness of the as-cast alloys Alloys
Bulk hardness (BHN/500/60)
ZA1 ZSr1 ZTiB1 ZNi1 ZNi3 ZNi4
119 158 148 124 130 138
AlCu are observed in the X-ray diffractograms of Fig. 11. The maximum solubility of Ni in Zn is negligible, as seen from the Ni–Zn binary phase diagram [17]. This results in the formation of minor amounts of the Ni–Zn intermetallics, Ni3 Zn22 and Ni5 Zn21 . Similar to those of the minor Ni-aluminides, only those peaks of Ni–Zn phases which coincide with other major phases in the alloy are observed in Fig. 11. 3.3. Hardness The bulk hardness of the as-cast alloys are given in Table 2. The addition of Al–5Ti–1B and Al–5Sr results in significant improvement in hardness. The refinement of the microstructure results in the significant increase in hardness of ZTiB1 and ZSr1. The addition of Ni also increases hardness to a certain extent. The addition of Ni results in the formation of hard Ni-aluminide particles contributing towards the increased hardness of the alloys ZNi1, ZNi3 and ZNi4. ZA27based alloys experience a spinodal reaction. However, no reference could be cited regarding the effect of the alloying elements on the spinodal reaction of this system. It is not possible to conclude the effect of alloying elements on spinodal reaction and ultimately the hardness at this point. 3.4. Effect of heat treatment on microstructure and hardness The as-cast sample was homogenized at 365 ◦ C for 1 h followed by quenching in water at room temperature. This treatment removes the segregation effects and the inhomogenities present in the as-cast alloy. The temperature of homogenization, 365 ◦ C, is selected on the basis of the binary Al–Zn phase diagram. In the binary system, a single-phase supersaturated solid-solution of Al exists at 365 ◦ C. The quenched samples were subsequently aged in an oven at 185 ◦ C for 15, 30, 45 and 65 min. 3.4.1. Alloy ZA1 The X-ray diffractograms of ZA1 in the heat-treated condition are shown in Fig. 12(a). The dissolution of Zn in Al results in the disappearance or drastic decrease in intensity of most of its peaks in the solutionized state (Fig. 12(b)). Since cooling is rapid, there is not enough time for the diffusion of Cu from Al4 Cu9 . As a result the peaks of Al4 Cu9 are still present in the as-quenched state. During heating of the quenched sample from room temperature, initially Zn
Fig. 12. X-ray diffractograms of (a) heat-treated ZA1 and (b) as-cast and then solutionized ZA1.
diffuses out of the supersaturated (Al) solid solution. This results in the formation of the broad exothermic peak between 50 and 100 ◦ C in the DSC curves (Fig. 13). The peaks in the upward direction indicate exothermic reaction, whereas those in the downward direction indicate endothermic reaction. At the same time, Cu atoms also diffuse out of the supersaturated (Al) phase. These Cu atoms together with the pre-existing free Cu clusters, form a diffusion couple between the Cu-rich end and Al + Zn rich phases. It is observed from the isothermal sections of Zn–Al–Cu at 200 ◦ C [12] and 20 ◦ C [13] that the phase to form first is CuZn4 , which together with the precipitation of Zn-rich phase contributes towards the broad peak formation in the temperature range of 50–100 ◦ C. With further progress in heating and hence diffusion, the phase field of AlCu is crossed resulting in the formation of this phase. The formation of AlCu results in the exothermic peak in the
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Fig. 13. Non-isothermal DSC curves of ZA1. Table 3 Activation energy for precipitation reaction (A) of the alloys Alloys
Activation energy (kJ/mol K)
ZA1 ZTiB1 ZSr1 ZNi1 ZNi3 ZNi4
35.24 27.69 41.53 22.56 35.24 35.24
temperature range of 150–200 ◦ C (marked A in Fig. 13). The sharp endothermic peak at about 280 ◦ C signifies the eutectoid reaction [18]. During aging, Cu diffuses out of Al4 Cu9 resulting in the formation of AlCu and subsequent decrease in amount and disappearance of Al4 Cu9 phase, as observed in Fig. 12(a). The activation energy of the precipitation reaction “A” in ZA1 is calculated by the ASTM E698 method which is based on the equation, E = −2.19R[(d log10 Q)/(d(1/T))]. The term within the parenthesis is the slope of the plot of log(Q) versus 1/T, where Q is the heating rate, T the peak temperature in absolute scale and R is the universal gas constant. The method is based on the assumption that the order of the
Fig. 14. Effect of heat treatment on hardness of ZA1, ZTiB1 and ZSr1.
Fig. 15. X-ray diffractograms of heat-treated ZTiB1 (Cu target).
Fig. 16. Non-isothermal DSC curves of ZTiB1.
Fig. 17. X-ray diffractograms of heat-treated ZSr1 (Cu target).
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reaction is 1. The activation energies so obtained for each alloy is shown in Table 3. Fig. 14 shows the effect of heat treatment on the hardness of ZA1. Hardness value at 0 min denotes that of the as-quenched sample. Beyond 30 min of aging the hardness drops sharply owing to grain growth.
Fig. 18. Non-isothermal DSC curves of ZSr1.
3.4.2. ZTiB1 Fig. 15 shows the X-ray diffractograms of the heat-treated ZTiB1. The DSC curves for this alloy are shown in Fig. 16. Except for the Ti-aluminides, the formation of the other phases remains same as that of ZA1. The exothermic peak corresponding to the precipitation of Zn and CuZn4 is not pronounced in this figure. As in the alloy ZA1, the precipitation of AlCu takes place in the temperature range of 150–200 ◦ C (the exothermic peak marked A in Fig. 16). The activation energy for the reaction “A” is shown in Table 3.
Fig. 19. X-ray diffractograms of heat-treated (a) ZNi1 (Co target), (b) ZNi3 (Cu target) and (c) ZNi4 (Cu target).
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Fig. 20. Non-isothermal DSC thermograms of (a) ZNi1, (b) ZNi3 and (c) ZNi4.
The effect of heat treatment on the hardness of ZTiB1 is shown in Fig. 14. The hardness is retained for a longer time than ZA1 owing to the presence of Al3 Ti precipitates. It may be pointed out that no significant rise in hardness is observed owing to simultaneous precipitation and grain growth at such a high aging temperature. Aging at a low temperature, say 100 ◦ C, would have shown a peak hardness [19]. 3.4.3. ZSr1 The X-ray diffractograms and DSC curves of ZSr1 are shown in Figs. 17 and 18, respectively. The signal corresponding to the precipitation of Zn and CuZn4 is quite weak (Fig. 18). The presence of Sr appears to prevent the diffusion of Cu from Al4 Cu9 during the rapid cooling of this alloy. As a result, Al4 Cu9 is retained in the as-quenched state and AlCu is not formed. Even during aging, the diffusion of Cu from Al4 Cu9 is quite less resulting in the retainment of most of Al4 Cu9 and the formation of a minor amount of AlCu, unlike that of the ZA1 alloy (Fig. 17). The significant increase in activation energy for reaction “A” (Table 3) is also suggestive of the reduction in the rate of the precipitation of AlCu. The effect of heat treatment on the hardness of ZSr1 is shown in Fig. 14. Sr is not as effective as Ti in retaining the hardness at such a high aging temperature. However, peak hardness is achievable at a low aging temperature like 100 ◦ C [19].
3.4.4. ZNi1, ZNi3 and ZNi4 Fig. 19(a)–(c) shows the X-ray diffractograms of ZNi1, ZNi3 and ZNi4, respectively. The DSC curves of alloy ZNi1, ZNi3 and ZNi4 are shown in Fig. 20(a)–(c), respectively. The broad exothermic peak below 100 ◦ C (as observed in the DSC of ZA1) is not observed in the DSC curves (Fig. 20). The signal for this peak is very weak in all the alloys. In ZNi1 and ZNi3, with a further decrease in amount of Cu, the amount of CuZn4 precipitated becomes negligibly small (Fig. 19(a) and (b)); it also contributes towards the weakening of its corresponding signal in the DSC curve (Fig. 20(a) and (b)). The X-ray diffractogram in Fig. 19(c) shows that the amount of CuZn4 precipitated during aging of ZNi4 is too less for the occurrence of any distinguishable peak. This is due to the very small amount of (0.8 wt.%) of Cu present in this alloy. Results in Table 3 show that the addition of Ni does not affect the activation energy for the precipitation reaction AlCu (peak “A”). The sharp endothermic peak at about 280 ◦ C signifies the eutectoid reaction [18]. The effect of heat treatment on the hardness of the ZA27 alloys containing Ni is shown in Fig. 21. Solutionization at 365 ◦ C followed by quenching of ZA27-based alloys containing Ni of varying amounts results in a significant rise in hardness. The solubility of Ni in Zn and Al is negligible even at elevated temperatures. Therefore, this increase in hardness owing to quenching cannot be attributed to the solid-solution
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Table 4 Summarized results on the effect of heat treatment on phases and hardness of the alloys Alloy
As cast
Solutionized (from 365 ◦ C)
Aging (185 ◦ C) 15 min
30 min
45 min
60 min
Phases present under different conditions
ZA1 (unmodified ZA27)
119
119
119
119
109
100
ZTiB1 (ZA27 + 0.1 wt.% Ti)
148
150
150
150
147
90
ZSr1 (ZA27 + 0.1 wt.% Sr)
158
158
158
100
96
91
ZNi1 (ZA27 + 0.3 wt.% Ni)
124
143
119
119
104
100
ZNi3 (ZA27 + 0.9 wt.% Ni)
130
158
158
143
136
100
ZNi4 (ZA27 + 2.0 wt.% Ni)
138
172
170
168
153
147
a b
Zn, AlCu, CuZn4 , AlMga , Al3 Mg2 a , Al4 Cu9 b , Mg2 Zn11 a in as-cast state; Al4 Cu9 b , AlCu, Znb in solutionized state; Zn, AlMga , Al3 Mg2 a , CuZn4 , AlCu, Al4 Cu9 b , Mg2 Zn11 a in aged condition Zn, Al, AlCub , CuZn4 , AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Ti, Ti8 Al24 a , AlTib , TiB2 a in as-cast condition; Znb , AlCu, Al4 Cu9 b , Al3 Ti, AlTib , in solutionized state; Zn, Al, AlCu, CuZn4 , AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Ti, AlTib , TiB2 a in aged condition Zn, Al, CuZn4 , AlMga , Al3 Mg2 a , Mg2 Zn11 a AlSrb , Al4 Srb present in as-cast condition; Znb , AlSr, Al4 Cu9 b in solutionized state; Zn, CuZn4 , AlMga , Al3 Mg2 a , Al4 Cu9 b , AlCub , AlSr, Al4 Srb , Mg2 Zn11 a in aged condition Zn, Al, AlCu, Al4 Cu9 a , AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Nib , Ni3 Ala , AlNia , Al3 Ni2 a , Ni3 Zn22 a , Ni5 Zn21 a in as-cast condition; Znb , AlCu, Al3 Nib , Al4 Cu9 b in solutionized state; Zn, AlCu, Al4 Cu9 b , Alb , AlMga , Al3 Mg2 a , Mg2 Zn11 a , CuZn4 b , Al3 Nib , Ni3 Ala , AlNia , Al3 Ni2 a , Ni3 Zn22 a , Ni5 Zn21 a in aged condition Zn, Al, AlCu, Al4 Cu9 a , AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Nib , NiAlb , Al3 Ni2 b , Ni3 Ala , Ni3 Zn22 a , Ni5 Zn21 a in as-cast condition; AlCu, Znb , Al4 Cu9 b , Al3 Nib in solutionized state; Zn, AlCu, Alb , CuZn4 a , Mg2 Zn11 a , AlMga , Al3 Mg2 a , Al3 Nia , Ni3 Ala , Al3 Ni2 a , Ni5 Zn21 a , Ni3 Zn22 a in aged condition Zn, Al, AlMga , Mg2 Zn11 a , Al3 Mg2 a , Al3 Nib , NiAlb , Al3 Ni2 b , Ni3 Ala , Ni3 Zn22 a , Ni5 Zn21 a present in as-cast condition; Znb , AlCu, Al4 Cu9 b , Al3 Nib in solutionized state; Zn, AlCu, AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Nib , Ni3 Ala , Al3 Ni2 a , NiAla in aged condition
Very small or negligible quantity. Minor quantity.
hardening. It is not possible at this moment to comment on the effect of Ni on the spinodal decomposition of such alloys. It is quite interesting to note that such alloys retain their hardness for a significant period during aging even at a high temperature. This retention of hardness improves with the increasing amount of Ni in the alloy. This behavior may be attributed to the presence of fine Al3 Ni precipitates preventing grain growth. Results discussed so far regarding the effect of heat treatment on the development of microstructure and hardness is summarized in Table 4.
4. Conclusions
Fig. 21. Effect of heat treatment on hardness of ZNi1, ZNi3 and ZNi4.
The as-cast alloys consist of primary Al-rich dendrites surrounded by the eutectoid constituents. The lamellar morphology of the eutectoid constituents have been clearly observed in the transmission electron micrographs of all the alloys studied. The addition of commercial master alloys has resulted in significant changes in the dendrite morphology. The addition of Ti and B results in rosette-shaped den-
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drites whereas that of Sr results in the formation of dendrites with reasonably sharp tips. The addition of Ni results in the formation of Ni–Al and Ni–Zn intermetallic phases randomly distributed in the matrix. In the presence of high amount of Ni, the dendrites undergo significant refinement. Owing to the microstructural refinement, as-cast alloys containing Ti or Sr undergo remarkable increase in hardness. The formation of hard Ni-containing intermetallic phases results in improvement in hardness to a certain extent. The effect of heat treatment is quite similar in all the alloys. Solutionizing and quenching treatment results in the dissolution of Zn in Al. During aging of the quenched samples from room temperature, Zn diffusion form the supersaturated (Al) solid-solution and precipitation of CuZn4 occurs at 50–100 ◦ C. This is followed by the precipitation of AlCu in the temperature range of 150–200 ◦ C. No significant improvement in hardness is observed in any of the alloys due to aging. This is owing to the simultaneous occurrence of precipitation and grain growth at such a high aging temperature (185 ◦ C). Ti is more effective than Sr in retaining the hardness at such a high aging temperature. The alloys containing Ni retain their hardness for a significant period during aging. Such a behavior may be attributed to the presence of fine Al3 Ni precipitates preventing grain growth.
Acknowledgement Financial support received from the Defence Research and Development Organization, Department of Defence Re-
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search and Development, Government of India to carry out this research is gratefully acknowledged.
References [1] J.K. Edward Jr., Adv. Mater. Process. 132 (1987) 51–57. [2] B.K. Prasad, A.K. Patwardhan, A.H. Yegneswaran, Metall. Mater. Trans. A 27 (1996) 3513–3523. [3] T. Calayag, D. Ferrea, Soc. Automotive Eng. 90 (1982) 40–44. [4] B.K. Prasad, A.K. Patwardhan, A.H. Yegneswaran, Scripta Metall. Mater. 37 (1997) 323–328. [5] Y. Chen, M. Tu, Mater. Sci. Technol. 14 (1998) 473. [6] K. Lal, R.K. Mahanti, C.S. Sivaramakrishnan, K.B. Jain, Mater. Trans. JIM 32 (1991) 621–625. [7] I.G. Ritchie, Z.L. Pan, F.E. Goodwin, Met. Mater. Trans. A 22 (1991) 607–616. [8] A.A. Abdel-Hamid, Z. Metallkd. 83 (1992) 314–318. [9] Z.M. El-Baradie, J. Mater. Process. Technol. 84 (1998) 73–78. [10] B.K. Prasad, A.K. Patwardhan, A.H. Yegneswaran, Wear 199 (1996) 142–151. [11] T.B. Massalski, J.L. Murray, L.H. Bennett, H. Baker (Eds.), Binary Alloy Phase Diagrams, vol. 1, 2nd ed., ASM, Metals Park, OH, 1990, p. 240. [12] D. Hanson, M.L.V. Gaylor, J. Inst. Met. London 34 (1925) 125–170. [13] E. Janecke, Kurzgefasstes Handbuch Aller Legierungen, Carl-Winter Press, Heidelberg, 1949, p. 569. [14] H.H. Arndt, K. Moeller, Z. Metallkd. 51 (1960) 596–600. [15] W. Koester, W. Dullenkopf, Z. Metallkd. 28 (1936) 363–367. [16] I. Maxwell, A. Hellawell, Metall. Trans. 3 (1972) 1487–1493. [17] T.B. Massalski, J.L. Murray, L.H. Bennett, H. Baker, Binary Alloy Phase Diagrams, vol. 3, 2nd ed., ASM, Metals Park, OH, 1990, p. 2888. [18] R. Kovacheva, R. Dobrev, S. Zadgorski, A. Lilova, Mater. Charact. 31 (1993) 217–224. [19] T. Savaskan, S. Murphy, Mater. Sci. Technol. 6 (1990) 695–703.
Evolution of as-cast and heat-treated microstructure of a commercial bearing alloy Pritha Choudhury, Karabi Das, Siddhartha Das ∗ Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721302, India Received 24 November 2004; received in revised form 17 March 2005; accepted 29 March 2005
Abstract The present study involves the microstructural changes of a zinc–aluminium alloy (ZA27) by nickel, a high melting point element, and the commercial master alloys, Al–5Ti–1B and Al–5Sr. The evolution of as-cast microstructure and the effect of heat treatment on the as-cast microstructure and hardness have been presented in detail. The addition of Ti, B and Sr as master alloys results in significant refinement in the dendrite structure of the ZA27 alloy. The addition of Ni to ZA27-based alloy results in the formation of Ni-aluminides as rod as well as blocky-shaped particles. High amount of Ni (2 wt.%) refines the dendritic structure significantly. © 2005 Elsevier B.V. All rights reserved. Keywords: ZA27; Dendrite structure; Bearing allloy
1. Introduction Zinc–aluminium (ZA) alloys, with a unique combination of properties, are competitive alternative materials to most aluminium casting alloys, bearing bronze, cast iron, as well as to plastics and steel fabrications [1–5]. ZA alloys are a group of high-performance, high-aluminium zinc alloys with excellent bearing properties [3]. Some of the attractive properties of the Zn–Al casting alloys are high sliding wear resistance (usefulness as bearings and bushings) [3], good machinability, excellent corrosion resistance in a variety of environments [6] and high damping capacity (suitability in defence applications) [7]. Addition of some alloying elements is known to have significant effect on the microstructure of these alloys [8,9]. High melting point element like Ni is reported to improve the sliding wear properties of the Zn–Al casting alloys [2,10]. However, literature pertaining to the systematic evolution of microstructure in such alloys is lacking. In this paper, it is aimed to present systematically the evolution of microstructure in a high-aluminium zinc-based alloy (ZA27) to which Ti, B, and Sr have been added in the form of commercial ∗
Corresponding author. E-mail address: [email protected] (S. Das).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.03.098
master alloys, Al–5Ti–1B and Al–5Sr. Such master alloys are generally used for the grain refinement and/or modification of some aluminium alloys. However, the use of grain refiners for the zinc–aluminium based alloys is not a common practice. The effect of Ni on the microstructure of ZA27 is also studied.
2. Experimental procedure All alloys based on ZA27 composition were prepared by conventional melting and casting route. They are identified as ZA1, ZTiB1, ZSr1, ZNi1, ZNi3 and ZNi4 for the sake of brevity and their compositions are shown in Table 1. Please note that ZA1 is the base ZA27 alloy. The alloys ZNi1, ZNi3 and ZNi4 contain 0.3, 0.9 and 2.0 wt.% Ni, respectively. For microstructural characterization, samples were metallographically polished according to standard practices and etched suitably using diluted nitric acid (5 vol.%) in water. The metallographically polished samples were subjected to image analysis using the Leica QWin image analysis software attached with a Leica DMRX optical microscope. Scanning electron microscopy was carried out in a JEOL scanning electron microscope (SEM) attached with an energy dispersive X-ray spectroscopic (EDS) facility. For transmission electron
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Table 1 Chemical composition of the alloys Alloys
Zn (wt.%)
Al (wt.%)
Cu (wt.%)
Mg (wt.%)
Ti (wt.%)
B (wt.%)
Sr (wt.%)
Ni (wt.%)
ZA1 ZTiB1 ZSr1 ZNi1 ZNi3 ZNi4 Al–5Ti–1B Al–5Sr
69.95 69.85 71.64 71.89 70.32 69.20 – –
28.0 27.72 26.0 26.2 26.0 26.8 94.0 95.0
2.0 2.25 2.20 1.56 1.58 0.80 – –
0.05 0.05 0.05 0.05 0.05 0.05 – –
– 0.1 – – – – 5 –
– 0.02 – – – – 1.00 –
– – 0.1 – – – – 5.0
– – – 0.3 0.9 2.0 – –
microscopy, thin foils with a thickness less than 100 m were prepared by dry grinding on silicon carbide papers. Discs of 3 mm diameter were punched from the thin foils. Final thinning of the discs was performed by electropolishing in a Struers Tenupol-5 double-jet electropolisher using 8% perchloric acid in 92% methanol as electrolyte at a voltage of 11.5 V and a temperature of −30 to −20 ◦ C. The thin foils were observed in Philips CM12 transmission electron microscope at an operating voltage of 120 kV. Bulk hardness of all the alloys were measured using a Brinell hardness tester, with a 10 mm diameter steel-ball indenter. The load applied was 500 kg for 60 s.
the (Al) solid solution undergoes a eutectoid reaction to form a mixture of ␣Al + (Zn). Accordingly, the base alloy, ZA1 comprises of primary ␣ dendrites surrounded by the eutectoid ␣Al + (Zn) as shown in Fig. 1. 3.1.2. Alloy ZTiB1 The optical and transmission electron (TEM) micrographs of ZTiB1 are shown in Fig. 2. The sequence of formation of
3. Results and discussion 3.1. Microstructure 3.1.1. Alloy ZA1 The optical micrograph of ZA1 is shown in Fig. 1. Zn–Al binary system is principally a eutectic system, consisting of a eutectoid reaction and a miscibility gap in the solid state having the critical point at 351.5 ◦ C. The ZA1 alloy fall in the hypereutectoid range of the Zn–Al phase diagram [11]. The phase to form first upon solidification is the primary ␣ (the solid-solution of zinc in aluminium). In the process, the excess zinc is rejected to the surrounding liquid. Subsequently,
Fig. 1. Optical micrograph of as-cast ZA1.
Fig. 2. (a) Optical and (b) TEM micrographs of as-cast ZTiB1.
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Fig. 3. (a) Optical and (b)–(d) TEM micrographs of as-cast ZSr1.
the Al-rich and Zn-rich phases are presumably same as in ZA1, already described. However, it is observed from the optical micrograph that the dendrites are rosette-shaped and much smaller in size than those in the unmodified alloy, i.e., ZA1 alloy. The lamellar structure of the eutectoid constituents is observed in the TEM micrographs (Fig. 2(b)). The average width of the lamellae is calculated to be 0.02 m. The lamellae perpendicular to the surface have a cellular appearance (Fig. 2(b)). Large black regions are titanium aluminide precipitates formed in the molten state. When the master alloy, Al–5Ti–1B is added to the molten alloy, the existing Ti-aluminides have enough time for growth during solidification and subsequent cooling, resulting in large particles (0.06–0.08 m). 3.1.3. ZSr1 The optical and TEM photographs of as-cast ZSr1 are shown in Fig. 3. It is observed from the optical micrographs that the addition of Sr in the form of Al–5Sr master alloy increases the curvature of the dendrites.
Fig. 3(b) shows a typical lamellar morphology of the eutectoid constituents. The average width of the eutectoid lamellae in this alloy is calculated to be 0.08 m. The fine precipitates (average size 0.03 m) observed in Fig. 3(c) are formed in the solid state and hence did not have enough time for growth. These precipitates are most likely those of AlCu or Al4 Cu9 (details given in a later section). Anti phase domain boundaries (APB) have been observed in this alloy (as marked by an arrow in Fig. 3(d)). During the fast cooling of the alloy, a number of small ordered regions have been formed rapidly. These ordered regions are separated from each other by the APBs, which are formed in the solid state. The ordered regions are most likely those of the intermetallic Al4 Sr, which was pre-existing in the master alloy and has been detected in as-cast ZSr1 by X-ray diffraction (described in a later section). 3.1.4. ZNi1, ZNi3 and ZNi4 Figs. 4–6 show the optical and SEM micrographs of the alloy ZNi1, ZNi3 and ZNi4. The sequence of formation of the Zn-rich and Al-rich phases are presumably same as in ZA1.
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Fig. 4. (a) Optical and (b) SEM micrographs of as-cast ZNi1: arrow indicates nickel-aluminide (Al3 Ni) particle.
Fig. 5. (a) Optical and (b) SEM micrographs of as-cast ZNi3: arrow indicates nickel-aluminide (Al3 Ni) particle.
Fig. 6. (a) Optical and (b) SEM micrographs of as-cast ZNi4: arrow indicates nickel-aluminide (Al3 Ni) particle.
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Fig. 7. Transmission electron micrograph (TEM) of as-cast ZNi3.
The particles of Ni-aluminides are observed in the optical as well as SEM micrographs. With increasing Ni, the morphology of the particles changes from blocky to rod-shape. The particles are blocky, mostly rod with some blocky and rodshaped in alloy ZNi1, ZNi3 and ZNi4, respectively. The Ni that has been added in powder form remains in the solid state during melting of the alloy (∼680 ◦ C). Hence, the formation of the Ni-aluminides occurs in the solid state. It is clearly evident from optical micrographs (Fig. 6(a)) that a significant refinement of dendrites, is achieved in ZNi4. The lamellar structure of the eutectoid constituents are clearly visible in Fig. 7. The fine precipitates (0.01–0.04 m) seen inside the eutectoid lamellae did not have enough time for growth and hence, are very small in size. The shape of these particles is nearly spherical which also suggests that they are in the initial stages of formation. These fine precipitates are likely to be those of AlCu and Al4 Cu9 and they formed in the solid state. The bigger black regions represent the precipitates formed in the liquid state. These precipitates formed on pre-existing particles and having longer time for growth they became larger in size (0.08–0.11 m in length). These particles are most likely those of the Ni-aluminides which started growing on the Ni particles added in powder form during melting. 3.2. Phase identification 3.2.1. ZA1 The major phases observed in ZA1 are shown in the X-ray diffractogram of Fig. 8. The alloys are cast in metal mold and hence subjected to rapid cooling. According to the low temperature (200 and 20 ◦ C) isothermal sections of Zn–Al–Cu [12,13], AlCu and CuZn4 are the phases existing in equilibrium with Al- and Zn-rich phases at respective temperatures. Al4 Cu9 is a phase found in the Cu-rich side of the isothermal section of Zn–Al–Cu ternary system at 20 ◦ C and also at the higher temperature of 350 ◦ C [14]. Since elemental Cu is added to the Zn–Al melt, Al4 Cu9 phase forms at
Fig. 8. X-ray diffractogram of as-cast ZA1 (Co target).
the solid Cu-melt interface. Owing to the rapid cooling and solidification, there is insufficient time for the complete decomposition of Al4 Cu9 into AlCu and hence some amount of Al4 Cu9 still remains in the as-cast state as observed in the XRD pattern. Other phases AlCu and CuZn4 are observed in the XRD pattern. The AlCu phase is not supposed to exist in the 20 ◦ C isothermal section. However, it is observed in the as-cast microstructure owing to a rapid cooling associated with casting in a metal mold. Mg present in minor quantities is added in order to strengthen and increase corrosion resistance of the alloy. Mg is present as the intermetallics AlMg and Mg2 Zn11 (Fig. 8) and the remaining amount goes into solid solution with the Al-rich phase. From the isothermal section of the Al–Mg–Zn system at 25 ◦ C [15], it is observed that Mg2 Zn11 may be present in ZA1. However, owing to rapid cooling, AlMg of the Mg-rich end does not get enough time for decomposition, and, hence, observed in the as-cast alloy (Fig. 8). However, one should note that peaks for AlMg and Mg2 Zn11 overlap with those of Zn. 3.2.2. ZTiB1 The major phases in the Al–5Ti–1B master alloy are shown in the X-ray diffractogram in Fig. 9. In the Al–Ti binary system [11], Al – 5 wt.% Ti is a hypoperitectic alloy from the Ti-end. It is observed from the isothermal section at 650 ◦ C of Al–Ti–B ternary system [16] that the alloy Al–5Ti–1B consists of Al3 Ti, TiB2 and the Al-rich solid solution. The phases Al3 Ti and (Al) are observed in Fig. 9. Thus, Al3 Ti is the major Ti-aluminide phase as obtained from the binary Al–Ti and the ternary Al–Ti–B phase diagrams. In Al–5Ti–1B, the volume percent of TiB2 is empirically calculated to be ∼0.714, which is far too less for the occurrence of any peak in the X-ray diffractogram. However, one small peak of very low intensity corresponding to that of TiB2 (Fig. 9) coincides with Al (the major phase), and therefore may not be considered.
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Fig. 9. X-ray diffractograms of master alloys, Al–5Ti–1B and Al–5Sr (Cu target).
ZTiB1 contains the phases present in the base alloy, ZA1, as well as additional phases containing Ti (Fig. 10). The phases Al3 Ti and AlTi are already present in the master alloy as stated earlier. These intermetallic phases having very high melting points remain undissolved in the Zn-based alloy during melting which was carried out at ∼680 ◦ C. The amount of Ti added is only 0.1 wt.% and hence it results in a very small volume fraction of the Ti-containing intermetallic phases. Generally peaks of Ti-containing intermetallics coincide with any other phase like Zn or Al. The phase Al24 Ti8 is an ordered superstructure of Al3 Ti. The other phases remain the same as that in the unmodified ZA27 alloy.
Fig. 10. X-ray diffractograms of as-cast ZTiB1 and ZSr1 (Co target).
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3.2.3. ZSr1 It is observed from the Al–Sr binary phase diagram [10] that the master alloy Al–5Sr (5 wt.% Sr) is a hypereutectic alloy of Al and Al4 Sr from the Al end. Al has negligible solid solubility for Sr. In addition to Al4 Sr, the master alloy also contains negligible amount of AlSr as shown in Fig. 9. The composition of AlSr closely matches that of Al7 Sr8 [11]. The intermetallic phases, Al4 Sr and AlSr (high melting points) of the master alloy remain undissolved during melting of the alloy ZSr1. Since the amount of Sr added is too small (0.1 wt.%), the volume fraction of the resultant Sr-containing intermetallic phases is also too less for the appearance of any peak in the X-ray diffractogram (Fig. 10). Hence, the peaks of Al4 Sr and AlSr having significant intensity are the ones coinciding with the peaks of the major phase Zn. The other phases present are those also found in the base alloy (ZA1). The APBs observed in the TEM of this alloy (Fig. 3(d)) formed in ordered Al4 Sr phase. 3.2.4. ZNi1, ZNi3 and ZNi4 The X-ray diffractograms of the alloys containing Ni are shown in Fig. 11. It is seen from this figure that the phases present in addition to those of ZA1 are those of the Nicontaining intermetallics, like, Al3 Ni, AlNi, Ni3 Al, Ni3 Zn22 and Ni5 Zn21 . According to the Al–Ni binary phase diagram [11], in the ZA27 alloys containing Ni, Al3 Ni presumably is the Ni-aluminide phase existing in equilibrium with the other phases. During the rapid cooling of the alloys, there is not enough time for diffusion at the Ni–Al interfaces. This could have resulted in the formation of the other Ni-aluminide phases already mentioned. The amount of Ni3 Al, NiAl or Al3 Ni2 formed in any of these alloys is too less for the appearance of any peak of reasonable intensity in the X-ray diffractogram. Hence, only those peaks of Ni3 Al, NiAl or Al3 Ni2 coinciding with any of the major phases like Zn or
Fig. 11. X-ray diffractograms of as-cast ZNi1, ZNi3 and ZNi4 (Co target).
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Table 2 Hardness of the as-cast alloys Alloys
Bulk hardness (BHN/500/60)
ZA1 ZSr1 ZTiB1 ZNi1 ZNi3 ZNi4
119 158 148 124 130 138
AlCu are observed in the X-ray diffractograms of Fig. 11. The maximum solubility of Ni in Zn is negligible, as seen from the Ni–Zn binary phase diagram [17]. This results in the formation of minor amounts of the Ni–Zn intermetallics, Ni3 Zn22 and Ni5 Zn21 . Similar to those of the minor Ni-aluminides, only those peaks of Ni–Zn phases which coincide with other major phases in the alloy are observed in Fig. 11. 3.3. Hardness The bulk hardness of the as-cast alloys are given in Table 2. The addition of Al–5Ti–1B and Al–5Sr results in significant improvement in hardness. The refinement of the microstructure results in the significant increase in hardness of ZTiB1 and ZSr1. The addition of Ni also increases hardness to a certain extent. The addition of Ni results in the formation of hard Ni-aluminide particles contributing towards the increased hardness of the alloys ZNi1, ZNi3 and ZNi4. ZA27based alloys experience a spinodal reaction. However, no reference could be cited regarding the effect of the alloying elements on the spinodal reaction of this system. It is not possible to conclude the effect of alloying elements on spinodal reaction and ultimately the hardness at this point. 3.4. Effect of heat treatment on microstructure and hardness The as-cast sample was homogenized at 365 ◦ C for 1 h followed by quenching in water at room temperature. This treatment removes the segregation effects and the inhomogenities present in the as-cast alloy. The temperature of homogenization, 365 ◦ C, is selected on the basis of the binary Al–Zn phase diagram. In the binary system, a single-phase supersaturated solid-solution of Al exists at 365 ◦ C. The quenched samples were subsequently aged in an oven at 185 ◦ C for 15, 30, 45 and 65 min. 3.4.1. Alloy ZA1 The X-ray diffractograms of ZA1 in the heat-treated condition are shown in Fig. 12(a). The dissolution of Zn in Al results in the disappearance or drastic decrease in intensity of most of its peaks in the solutionized state (Fig. 12(b)). Since cooling is rapid, there is not enough time for the diffusion of Cu from Al4 Cu9 . As a result the peaks of Al4 Cu9 are still present in the as-quenched state. During heating of the quenched sample from room temperature, initially Zn
Fig. 12. X-ray diffractograms of (a) heat-treated ZA1 and (b) as-cast and then solutionized ZA1.
diffuses out of the supersaturated (Al) solid solution. This results in the formation of the broad exothermic peak between 50 and 100 ◦ C in the DSC curves (Fig. 13). The peaks in the upward direction indicate exothermic reaction, whereas those in the downward direction indicate endothermic reaction. At the same time, Cu atoms also diffuse out of the supersaturated (Al) phase. These Cu atoms together with the pre-existing free Cu clusters, form a diffusion couple between the Cu-rich end and Al + Zn rich phases. It is observed from the isothermal sections of Zn–Al–Cu at 200 ◦ C [12] and 20 ◦ C [13] that the phase to form first is CuZn4 , which together with the precipitation of Zn-rich phase contributes towards the broad peak formation in the temperature range of 50–100 ◦ C. With further progress in heating and hence diffusion, the phase field of AlCu is crossed resulting in the formation of this phase. The formation of AlCu results in the exothermic peak in the
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Fig. 13. Non-isothermal DSC curves of ZA1. Table 3 Activation energy for precipitation reaction (A) of the alloys Alloys
Activation energy (kJ/mol K)
ZA1 ZTiB1 ZSr1 ZNi1 ZNi3 ZNi4
35.24 27.69 41.53 22.56 35.24 35.24
temperature range of 150–200 ◦ C (marked A in Fig. 13). The sharp endothermic peak at about 280 ◦ C signifies the eutectoid reaction [18]. During aging, Cu diffuses out of Al4 Cu9 resulting in the formation of AlCu and subsequent decrease in amount and disappearance of Al4 Cu9 phase, as observed in Fig. 12(a). The activation energy of the precipitation reaction “A” in ZA1 is calculated by the ASTM E698 method which is based on the equation, E = −2.19R[(d log10 Q)/(d(1/T))]. The term within the parenthesis is the slope of the plot of log(Q) versus 1/T, where Q is the heating rate, T the peak temperature in absolute scale and R is the universal gas constant. The method is based on the assumption that the order of the
Fig. 14. Effect of heat treatment on hardness of ZA1, ZTiB1 and ZSr1.
Fig. 15. X-ray diffractograms of heat-treated ZTiB1 (Cu target).
Fig. 16. Non-isothermal DSC curves of ZTiB1.
Fig. 17. X-ray diffractograms of heat-treated ZSr1 (Cu target).
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reaction is 1. The activation energies so obtained for each alloy is shown in Table 3. Fig. 14 shows the effect of heat treatment on the hardness of ZA1. Hardness value at 0 min denotes that of the as-quenched sample. Beyond 30 min of aging the hardness drops sharply owing to grain growth.
Fig. 18. Non-isothermal DSC curves of ZSr1.
3.4.2. ZTiB1 Fig. 15 shows the X-ray diffractograms of the heat-treated ZTiB1. The DSC curves for this alloy are shown in Fig. 16. Except for the Ti-aluminides, the formation of the other phases remains same as that of ZA1. The exothermic peak corresponding to the precipitation of Zn and CuZn4 is not pronounced in this figure. As in the alloy ZA1, the precipitation of AlCu takes place in the temperature range of 150–200 ◦ C (the exothermic peak marked A in Fig. 16). The activation energy for the reaction “A” is shown in Table 3.
Fig. 19. X-ray diffractograms of heat-treated (a) ZNi1 (Co target), (b) ZNi3 (Cu target) and (c) ZNi4 (Cu target).
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Fig. 20. Non-isothermal DSC thermograms of (a) ZNi1, (b) ZNi3 and (c) ZNi4.
The effect of heat treatment on the hardness of ZTiB1 is shown in Fig. 14. The hardness is retained for a longer time than ZA1 owing to the presence of Al3 Ti precipitates. It may be pointed out that no significant rise in hardness is observed owing to simultaneous precipitation and grain growth at such a high aging temperature. Aging at a low temperature, say 100 ◦ C, would have shown a peak hardness [19]. 3.4.3. ZSr1 The X-ray diffractograms and DSC curves of ZSr1 are shown in Figs. 17 and 18, respectively. The signal corresponding to the precipitation of Zn and CuZn4 is quite weak (Fig. 18). The presence of Sr appears to prevent the diffusion of Cu from Al4 Cu9 during the rapid cooling of this alloy. As a result, Al4 Cu9 is retained in the as-quenched state and AlCu is not formed. Even during aging, the diffusion of Cu from Al4 Cu9 is quite less resulting in the retainment of most of Al4 Cu9 and the formation of a minor amount of AlCu, unlike that of the ZA1 alloy (Fig. 17). The significant increase in activation energy for reaction “A” (Table 3) is also suggestive of the reduction in the rate of the precipitation of AlCu. The effect of heat treatment on the hardness of ZSr1 is shown in Fig. 14. Sr is not as effective as Ti in retaining the hardness at such a high aging temperature. However, peak hardness is achievable at a low aging temperature like 100 ◦ C [19].
3.4.4. ZNi1, ZNi3 and ZNi4 Fig. 19(a)–(c) shows the X-ray diffractograms of ZNi1, ZNi3 and ZNi4, respectively. The DSC curves of alloy ZNi1, ZNi3 and ZNi4 are shown in Fig. 20(a)–(c), respectively. The broad exothermic peak below 100 ◦ C (as observed in the DSC of ZA1) is not observed in the DSC curves (Fig. 20). The signal for this peak is very weak in all the alloys. In ZNi1 and ZNi3, with a further decrease in amount of Cu, the amount of CuZn4 precipitated becomes negligibly small (Fig. 19(a) and (b)); it also contributes towards the weakening of its corresponding signal in the DSC curve (Fig. 20(a) and (b)). The X-ray diffractogram in Fig. 19(c) shows that the amount of CuZn4 precipitated during aging of ZNi4 is too less for the occurrence of any distinguishable peak. This is due to the very small amount of (0.8 wt.%) of Cu present in this alloy. Results in Table 3 show that the addition of Ni does not affect the activation energy for the precipitation reaction AlCu (peak “A”). The sharp endothermic peak at about 280 ◦ C signifies the eutectoid reaction [18]. The effect of heat treatment on the hardness of the ZA27 alloys containing Ni is shown in Fig. 21. Solutionization at 365 ◦ C followed by quenching of ZA27-based alloys containing Ni of varying amounts results in a significant rise in hardness. The solubility of Ni in Zn and Al is negligible even at elevated temperatures. Therefore, this increase in hardness owing to quenching cannot be attributed to the solid-solution
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Table 4 Summarized results on the effect of heat treatment on phases and hardness of the alloys Alloy
As cast
Solutionized (from 365 ◦ C)
Aging (185 ◦ C) 15 min
30 min
45 min
60 min
Phases present under different conditions
ZA1 (unmodified ZA27)
119
119
119
119
109
100
ZTiB1 (ZA27 + 0.1 wt.% Ti)
148
150
150
150
147
90
ZSr1 (ZA27 + 0.1 wt.% Sr)
158
158
158
100
96
91
ZNi1 (ZA27 + 0.3 wt.% Ni)
124
143
119
119
104
100
ZNi3 (ZA27 + 0.9 wt.% Ni)
130
158
158
143
136
100
ZNi4 (ZA27 + 2.0 wt.% Ni)
138
172
170
168
153
147
a b
Zn, AlCu, CuZn4 , AlMga , Al3 Mg2 a , Al4 Cu9 b , Mg2 Zn11 a in as-cast state; Al4 Cu9 b , AlCu, Znb in solutionized state; Zn, AlMga , Al3 Mg2 a , CuZn4 , AlCu, Al4 Cu9 b , Mg2 Zn11 a in aged condition Zn, Al, AlCub , CuZn4 , AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Ti, Ti8 Al24 a , AlTib , TiB2 a in as-cast condition; Znb , AlCu, Al4 Cu9 b , Al3 Ti, AlTib , in solutionized state; Zn, Al, AlCu, CuZn4 , AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Ti, AlTib , TiB2 a in aged condition Zn, Al, CuZn4 , AlMga , Al3 Mg2 a , Mg2 Zn11 a AlSrb , Al4 Srb present in as-cast condition; Znb , AlSr, Al4 Cu9 b in solutionized state; Zn, CuZn4 , AlMga , Al3 Mg2 a , Al4 Cu9 b , AlCub , AlSr, Al4 Srb , Mg2 Zn11 a in aged condition Zn, Al, AlCu, Al4 Cu9 a , AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Nib , Ni3 Ala , AlNia , Al3 Ni2 a , Ni3 Zn22 a , Ni5 Zn21 a in as-cast condition; Znb , AlCu, Al3 Nib , Al4 Cu9 b in solutionized state; Zn, AlCu, Al4 Cu9 b , Alb , AlMga , Al3 Mg2 a , Mg2 Zn11 a , CuZn4 b , Al3 Nib , Ni3 Ala , AlNia , Al3 Ni2 a , Ni3 Zn22 a , Ni5 Zn21 a in aged condition Zn, Al, AlCu, Al4 Cu9 a , AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Nib , NiAlb , Al3 Ni2 b , Ni3 Ala , Ni3 Zn22 a , Ni5 Zn21 a in as-cast condition; AlCu, Znb , Al4 Cu9 b , Al3 Nib in solutionized state; Zn, AlCu, Alb , CuZn4 a , Mg2 Zn11 a , AlMga , Al3 Mg2 a , Al3 Nia , Ni3 Ala , Al3 Ni2 a , Ni5 Zn21 a , Ni3 Zn22 a in aged condition Zn, Al, AlMga , Mg2 Zn11 a , Al3 Mg2 a , Al3 Nib , NiAlb , Al3 Ni2 b , Ni3 Ala , Ni3 Zn22 a , Ni5 Zn21 a present in as-cast condition; Znb , AlCu, Al4 Cu9 b , Al3 Nib in solutionized state; Zn, AlCu, AlMga , Al3 Mg2 a , Mg2 Zn11 a , Al3 Nib , Ni3 Ala , Al3 Ni2 a , NiAla in aged condition
Very small or negligible quantity. Minor quantity.
hardening. It is not possible at this moment to comment on the effect of Ni on the spinodal decomposition of such alloys. It is quite interesting to note that such alloys retain their hardness for a significant period during aging even at a high temperature. This retention of hardness improves with the increasing amount of Ni in the alloy. This behavior may be attributed to the presence of fine Al3 Ni precipitates preventing grain growth. Results discussed so far regarding the effect of heat treatment on the development of microstructure and hardness is summarized in Table 4.
4. Conclusions
Fig. 21. Effect of heat treatment on hardness of ZNi1, ZNi3 and ZNi4.
The as-cast alloys consist of primary Al-rich dendrites surrounded by the eutectoid constituents. The lamellar morphology of the eutectoid constituents have been clearly observed in the transmission electron micrographs of all the alloys studied. The addition of commercial master alloys has resulted in significant changes in the dendrite morphology. The addition of Ti and B results in rosette-shaped den-
P. Choudhury et al. / Materials Science and Engineering A 398 (2005) 332–343
drites whereas that of Sr results in the formation of dendrites with reasonably sharp tips. The addition of Ni results in the formation of Ni–Al and Ni–Zn intermetallic phases randomly distributed in the matrix. In the presence of high amount of Ni, the dendrites undergo significant refinement. Owing to the microstructural refinement, as-cast alloys containing Ti or Sr undergo remarkable increase in hardness. The formation of hard Ni-containing intermetallic phases results in improvement in hardness to a certain extent. The effect of heat treatment is quite similar in all the alloys. Solutionizing and quenching treatment results in the dissolution of Zn in Al. During aging of the quenched samples from room temperature, Zn diffusion form the supersaturated (Al) solid-solution and precipitation of CuZn4 occurs at 50–100 ◦ C. This is followed by the precipitation of AlCu in the temperature range of 150–200 ◦ C. No significant improvement in hardness is observed in any of the alloys due to aging. This is owing to the simultaneous occurrence of precipitation and grain growth at such a high aging temperature (185 ◦ C). Ti is more effective than Sr in retaining the hardness at such a high aging temperature. The alloys containing Ni retain their hardness for a significant period during aging. Such a behavior may be attributed to the presence of fine Al3 Ni precipitates preventing grain growth.
Acknowledgement Financial support received from the Defence Research and Development Organization, Department of Defence Re-
343
search and Development, Government of India to carry out this research is gratefully acknowledged.
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