- Email: [email protected]
Effect of prior cold work on age hardening of Cu–3Ti–1Cr alloy
Materials Characterization 57 (2006) 348 – 357
Effect of prior cold work on age hardening of Cu–3Ti–1Cr alloy R. Markandeya a , S. Nagarjuna b,⁎, D.S. Sarma c a
c
Department of Metallurgical Engineering, College of Engineering, Jawaharlal Nehru Technological University, Kukatpally, Hyderabad-500 072, India b Defence Metallurgical Research Laboratory, Kanchanbagh, Hyderabad-500 058, India Department of Metallurgical Engineering, Institute of Technology, Banaras Hindu University, Varanasi-221 005, India Received 1 August 2005; accepted 28 February 2006
Abstract The influence of 50%, 75% and 90% cold work on the age hardening behavior of Cu–3Ti–1Cr alloy has been investigated by hardness and tensile tests, and light optical and transmission electron microscopy. Hardness increased from 118 Hv in the solution-treated condition to 373 Hv after 90% cold work and peak aging. Cold deformation reduced the peak aging time and temperature of the alloy. The yield strength and ultimate tensile strength reached a maximum of 1090 and 1110 MPa, respectively, following 90% deformation and peak aging. The microstructure of the deformed alloy exhibited elongated grains and deformation twins. The maximum strength on peak aging was obtained due to precipitation of the ordered, metastable and coherent β′-Cu4Ti phase, in addition to high dislocation density and deformation twins. Over-aging resulted in decreases in hardness and strength due to the formation of incoherent and equilibrium β-Cu3Ti phase in the form of a cellular structure. However, the morphology of the discontinuous precipitation changed to a globular form on high deformation. The mechanical properties of Cu–3Ti–1Cr alloy are superior to those of Cu–2.7Ti, Cu–3Ti–1Cd and the commercial Cu–0.5Be–2.5Co alloys in the cold-worked and peak-aged condition. © 2006 Elsevier Inc. All rights reserved. Keywords: Cu–3Ti–1Cr alloy; Precipitation hardening; Prior cold work; Tensile properties; Discontinuous precipitation
1. Introduction World wide-research on the binary Cu–Ti alloys has indicated that they have to serve as substitutes for the expensive and toxic Cu–Be–Co alloys. Extensive research has been carried out by many researchers on the mechanisms of spinodal decomposition and precipitation strengthening in Cu–Ti alloys [1–7]. Nagarjuna et al. have studied the structure–property correlations of Cu–Ti alloys in various conditions, viz., solution ⁎ Corresponding author. Tel.: +91 40 24586409; fax: +91 40 24342123. E-mail address: [email protected] (S. Nagarjuna). 1044-5803/$ - see front matter © 2006 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2006.02.017
treatment, solution treatment + aging and solution treatment + cold work + aging [8,9]. It was reported that compositional modulations would occur during solution treatment itself in high Ti alloys (N4.0 wt.%) and age hardening takes place by the formation of the metastable β′-Cu4Ti phase in both undeformed and deformed alloys. The tensile properties of these alloys are comparable to those of the Cu–Be–Co alloys [8,9]. Our earlier studies [10,11] on the addition of 1 wt.% Cd to a Cu–Ti alloy revealed that the tensile properties of the alloy increased considerably in solution-treated condition due to substitutional solid solution strengthening. The tensile properties of the ternary Cu–Ti–Cd alloys were also comparable to those of a Cu–Be–Co
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
alloy in the undeformed condition. The strengthening mechanism and microstructure of the Cu–Ti–Cd alloys were similar to those of the binary Cu–Ti alloys. Prior cold work increased the tensile properties considerably, but no microstructural changes were reported in the Cu– Ti–Cd alloys during peak aging. The strengthening mechanisms of cold-worked Cu–Ti–Cd alloys [12,13] were also similar to those in undeformed condition. Our recent studies on Cu–Ti alloys with 1 wt.% Cr addition revealed that the tensile properties of the Cu– Ti–Cr alloys increased considerably during peak aging after solution treatment [14]. Age hardening in these alloys was similar to that in the binary Cu–Ti alloys. Further, we recently investigated the effects of prior cold work on the age hardening of a Cu–4Ti–1Cr alloy [15], wherein it has been reported that the properties increased significantly with cold deformation followed by aging. Moreover, the tensile properties are comparable to those of the Cu–Ti–Cd and Cu–2Be–0.6Co alloys. However, little work has been reported on the effects of prior cold work on the age hardening of Cu–Ti–Cr alloys with a low Ti content (b4.0 wt.%). Therefore, the present investigation was undertaken with the aim of filling this gap and to generate necessary data on a Cu– 3Ti–1Cr alloy. This paper presents results on the effects of prior cold work on a Cu–3Ti–1Cr alloy and focuses on the strength properties, microstructure and strengthening mechanism. 2. Experimental procedure A Cu–Ti–Cr alloy was prepared with the aimed composition of 3 wt.% Ti, 1 wt.% Cr and Cu as the balance. A 30 kg melt of the alloy was made by melting the appropriate proportions of pure copper, titanium and chromium in the graphite crucible of Stokes vacuum induction melting (VIM) furnace and then casting into a graphite mould. The ingot was homogenized at 850 °C for 24 h and analyzed for its composition. The analysis report confirmed the composition to be 2.97 wt.% Ti, 0.99 wt.% Cr and balance Cu. The ingot was subsequently hot forged and rolled after soaking at 850 °C to obtain 10 mm thick flats and 12 mm diameter rods. After solutionizing at 860 °C for 2 h, the samples were given different amounts of deformation viz., 50%, 75% and 90%, at room temperature. Vickers hardness at 10 kg load was measured after aging the deformed samples at 400, 450 and 500 °C to establish peak aging time and temperature. A total of 6 indents were taken for each sample at frequent time intervals after aging at the above temperatures and the average value of hardness recorded.
349
Flat tensile specimens as per ASTM E8M standard were prepared from the solution treated + cold deformed + aged alloy, then tested at room temperature and at a strain rate of 10− 3 s− 1, using an Instron universal tensile testing machine. Metallographic specimens were prepared by mechanical polishing then etching in a solution of 10 g K2Cr2O7, 5 ml H2SO4, a few drops of HCl and 95 ml distilled water. The microstructures thus revealed were examined in a light optical microscope (LOM). Thin slices were cut from the cold deformed and aged bulk samples using a low-speed Isomet cutting machine and then mechanically polished to obtain 40–50 μm thick slices. Discs of 3 mm diameter were punched out of these slices and electropolished in a twin-jet electropolisher in a solution of 30 vol.% HNO3 and 70 vol.% methanol at − 45 °C and 20 V. These samples were studied using a JEOL 200CX transmission electron microscope (TEM) operating at 160 kV. 3. Results 3.1. Hardness Fig. 1 shows the effect of prior cold work on the hardness of the Cu–3Ti–1Cr alloy during aging at 400, 450 and 500 °C. These figures also include the variation in hardness with aging time of the undeformed alloy [14] at the above aging temperatures. Cold work prior to aging increased the hardness in solution-treated (ST + CW) as well as peak-aged (ST + CW + PA) conditions. However, peak aging time and temperature decreased with increasing amount of deformation. The peak aging time decreased from 2 h for 50% deformation to 1 h for 90% deformation when aged at 450 °C compared to 24 h for undeformed alloy at the same temperature. The maximum hardness for the Cu–3Ti–1Cr alloy rose from 118 Hv in the undeformed (solution-treated) condition to 373 Hv with 90% deformation and peak aging at 400 °C (Fig. 1a). Similarly, the maximum hardness obtained for 90% deformed alloy after aging at 450 °C was 340 Hv (Fig. 1b) and at 500 °C, 312 Hv (Fig. 1c). The variation of hardness with the amount of prior cold work before and after aging is shown in Fig. 2, which reveals that hardness increased considerably with the amount of prior cold work. Peak hardness was found to be nearly independent of aging temperature for the alloy with deformations of 50% and 75%, but with 90% deformation the peak hardness was highest after aging at 400 °C and thereafter dropped off considerably as the aging temperature increased. Peak hardness after aging at 500 °C is observed to be almost independent of the amount of prior deformation.
350
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
Fig. 2. Effect of prior cold work on hardness of solution-treated and peak-aged Cu–3Ti–1Cr alloy.
3.2. Tensile properties Fig. 3 compares the variation of YS and UTS with the amount of cold work following peak aging at 400 °C for the Cu–3Ti–1Cr alloy of this investigation with the data obtained on a Cu–4Ti–1Cr alloy [15]. For the latter alloy, both the YS and the UTS increased continuously up to 90% deformation. By contrast, in the case of the Cu–3Ti–1Cr alloy, the YS and UTS were nearly constant up to 50% deformation, but then increased slowly up to 75% deformation and reached maxima at 90% deformation. Further, the YS and UTS were nearly the same for 90% deformed and peak-aged Cu–3Ti–1Cr alloy.
Fig. 1. Effect of prior cold work on hardness during aging of Cu–3Ti– 1Cr alloy. (a) 400 °C; (b) 450 °C; (c) 500 °C. Fig. 3. Effect of prior cold work on strength properties of peak-aged Cu–Ti–Cr alloys.
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
351
The mechanical properties of the cold deformed Cu– 3Ti–1Cr alloy and peak aged at 400 °C are compared with those of the binary Cu–2.7Ti, Cu–3Ti–1Cd and Cu– 0.5Be–2.5Co alloys in Table 1. The YS of the Cu–3Ti–1Cr alloy increased from 757 MPa with 50% deformation to 1090 MPa with 90% deformation while the UTS increased from 890 MPa with 50% deformation to 1110 MPa on 90% deformation compared to 721 MPa YS and 854 MPa UTS in the undeformed and peak-aged condition. The ductility (% elongation) of the Cu–3Ti–1Cr alloy decreased from 17 for 50% cold work to 2.5 for 90% cold work compared to 21 in peak-aged condition for undeformed alloy. The YS and UTS are higher for the Cu–3Ti–1Cr alloy in 90% coldwork and aged condition than those of Cu–2.7Ti [8] and Cu–3Ti–1Cd [13] alloys in a similar condition. Further, the mechanical properties of the Cu–3Ti–1Cr alloy are superior to those of the commercial Cu–0.5Be–2.5Co alloy [16] in the cold-worked and peak-aged condition. 3.3. Light optical microscopy The microstructure of the undeformed Cu–3Ti–1Cr alloy [14] as seen in the LOM is briefly described here for comparison. In the solution-treated condition it displayed an equiaxed single-phase structure with fine grains. On peak aging at 450 °C for 24 h, the alloy exhibited a similar microstructure. However, it revealed onset of discontinuous precipitation at the grain boundaries of the matrix when over-aged, either at 450 °C for 88 h or 500 °C for 24 h. Fig. 4(a), (b) and (c) shows, respectively, the microstructures of the solution-treated and 50% cold-worked Cu–3Ti–1Cr alloy in peak-aged (400 °C/2 h), over-aged Table 1 Comparison of mechanical properties of solution-treated (ST) + coldworked (CW) + peak-aged (PA) Cu–3Ti–1Cr alloy with those reported for Cu–2.7Ti [9], Cu–3Ti–1Cd [13] and Cu–0.5Be–2Co [16] alloys Property
ST
Cold-worked Cu– Cu– 3Ti–1Cr, peak aged at 2.7Ti 400 °C (time) [9]
Cu– 3Ti– 1Cd [13]
Cu– 0.5Be– 2.5Co [16]
0% 50% 90% CW CW CW (24 h) (2 h) (1 h)
ST + 90% CW + PA
ST + 90% CW + PA
ST + full hard + aged (TH04 temper)
950 1000 355
922 1035 375
YS, MPa 215 721 UTS, MPa 464 854 Hardness, 118 277 Hv Elongation, %
43
21
757 890 300
17
1090 1110 373
2.5
3.5
690–825 760–895 220–270 (95–102 HRB) 6.0 10–20
Fig. 4. Light optical micrographs of solution-treated and 50% coldworked Cu–3Ti–1Cr alloy. (a) Peak aged at 400 °C/2 h; (b) over-aged at 450 °C/8 h; (c) over-aged at 500 °C/24 h.
1 (450 °C for 8 h) and over-aged 2 (500 °C for 24 h) conditions. The equiaxed grains were disturbed by the cold deformation and the alloy exhibited a single-phase structure with severe strain markings in the peak-aged condition (Fig. 4(a)). The deformation structure recovered slightly on over-aging at 450 °C for 8 h and the occasional occurrence of discontinuous precipitation was observed at the grain boundaries of the matrix (Fig. 4(b)). The discontinuous precipitation grew to a large extent on over-aging at 500 °C for 24 h (Fig. 4(c)). On 90% cold work, the grain elongation was severe, as revealed in Fig. 5. When peak aged at 400 °C for 1 h, a deformed grain structure along with the onset of discontinuous precipitation was seen (Fig. 5(a)). On over-
352
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
aging at 450 °C for 8 h, elongated grains and a significant amount of discontinuous precipitation were observed (Fig. 5(b)), while a substantial amount of the matrix is covered by discontinuous precipitation on over-aging at 500 °C for 24 h (Fig. 5(c)). 3.4. Transmission electron microscopy TEM studies were carried out on 50% and 90% deformed Cu–3Ti–1Cr alloy to understand the structural changes that occurred during peak aging and over-aging after high deformations. Fig. 6, a bright field (BF) image of the Cu–3Ti–1Cr alloy in the peak-aged condition (400 °C for 2 h) after
Fig. 6. TEM BF image of solution-treated and 50% cold-worked Cu– 3Ti–1Cr alloy peak aged at 400 °C for 2 h, showing fine precipitate of β′ Cu4Ti, twins and dislocation cells.
Fig. 5. Light optical micrographs of solution-treated and 90% coldworked Cu–3Ti–1Cr alloy. (a) Peak aged at 400 °C/1 h; (b) over-aged at 450 °C/8 h; (c) over-aged at 500 °C/24 h.
50% prior cold work, reveals dislocation cells and deformation twins. In addition, an ordered, metastable and coherent precipitate, β′-Cu4Ti, was seen. The selected area diffraction (SAD) pattern and its schematic diagram to identify the precipitate are not shown here because the same was reported in our recent work on undeformed Cu–Ti–Cr alloys [14], wherein the metastable and coherent phase was identified as β′-Cu4Ti precipitate. The β′ phase is reported to have a body-centered tetragonal (bct) crystal structure [5]. Deformation twins are more clearly revealed in the BF and dark field (DF) images of the 50% deformed alloy over-aged at 450 °C for 8 h (Fig. 7(a) and (b), respectively). The SAD pattern and its schematic diagram (shown in Fig. 7(c) and (d), respectively) identify and confirm the presence of twins. Fig. 7(a) also reveals the formation of some discontinuous precipitation surrounded by a highly strained matrix. Fig. 8 shows the BF images of the alloy with 50% cold deformation and over-aged at 500 °C for 24 h. Fig. 8(a) shows the coarse β-Cu3Ti equilibrium precipitate formed on over-aging while Fig. 8(b) shows the highly strained matrix and some fine β precipitation. The SAD pattern and its schematic to identify the precipitate β-Cu3Ti are not included here as they are the same as depicted in Fig. 11. Fig. 9, a BF image of 90% deformed Cu–3Ti–1Cr alloy peak aged at 400 °C, reveals increased twin density and dislocation cell structure. The ordered, metastable and coherent β′-Cu4Ti precipitate was also observed here. The deformation twins in the matrix of the 90% deformed Cu–3Ti–1Cr alloy over-aged at 450 °C for 8 h are clearly visible in the BF and DF images of Fig. 10(a) and (b). The SAD in Fig. 10(c) and its schematic in Fig. 10(d) confirm the presence of twins in the alloy. The dual phases of matrix and coarse β-Cu3Ti precipitate revealed in the BF and DF images of Fig. 11(a) and (b), respectively, are from
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
353
Fig. 7. TEM images of solution-treated and 50% cold-worked Cu–3Ti–1Cr alloy over-aged at 450 °C for 8 h. (a) BF image; (b) DF image; (c) SAD; (d) schematic of SAD of deformation twins.
this alloy after over-aging at 500 °C for 24 h. Interphase boundary dislocations can also be seen in the α-matrix (Fig. 11(a)). The SAD of Fig. 11(c) and its schematic (Fig. 11(d)) confirm the presence of equilibrium precipitate, β. 4. Discussion The present investigation has revealed that the hardness and strength of solution-treated Cu–3Ti–1Cr alloy
Fig. 8. TEM BF images of solution-treated and 50% cold-worked Cu– 3Ti–1Cr alloy over-aged at 500 °C for 24 h, showing (a) matrix and coarse β precipitate and (b) matrix with very fine β precipitate.
Fig. 9. TEM BF image of solution-treated and 90% cold-worked Cu– 3Ti–1Cr alloy peak aged at 400 °C for 1 h, showing fine precipitate of β′ Cu4Ti, deformation twins and dislocation cells.
354
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
Fig. 10. TEM images of solution-treated and 90% cold-worked Cu–3Ti–1Cr alloy over-aged at 450 °C for 8 h. (a) BF image; (b) DF image; (c) SAD; (d) schematic of SAD.
are significantly improved by cold work and subsequent peak aging at 400 °C. 4.1. Solution treatment and cold work (ST + CW) Cold rolling to a 90% reduction in thickness increased the hardness of the solution-treated Cu–3Ti–1Cr alloy (118 Hv) by 115%, to 254 Hv. By comparison, the same processing of a Cu–4Ti–1Cr alloy resulted in only a 40% increase, but in that case the initial solution-treated hardness was 224 Hv [14]. The higher initial hardness of the Cu–4Ti–1Cr alloy is attributed to compositional modulations due to the higher titanium content [5] and the presence of β′-Cu4Ti precipitate, whereas the solutiontreated Cu–3Ti–1Cr alloy is characterized by the presence of a single-phase (homogeneous) structure. Thus the higher percentage increase in hardness due to the application of cold working seen in the Cu–3Ti–1Cr alloy is attributed to the greater effect of the dislocation networks formed during the deformation process. 4.2. Solution treatment, cold work and peak aging (ST + CW + PA) The hardness of the Cu–3Ti–1Cr alloy increased considerably on cold working followed by aging,
reaching a maximum hardness of 373 Hv after 90% deformation and aging for 1 h at 400 °C. The highest hardness values for all the deformations were obtained when the alloy was aged at 400 °C whereas the undeformed alloy exhibited its peak hardness after aging at 450 °C [14]. This indicates that cold deformation reduced the peak aging temperature, which is in agreement with the findings on the binary Cu–Ti alloys [9]. Because of the presence of Cr in a super-saturated solid solution and to the extensive deformation, the aging kinetics were accelerated, giving rise to maximum strength values at a lower aging temperature (400 °C) and promoting the precipitation of a very fine metastable β′ phase in increased volume fraction. The peak aging times were generally in the 1–2 h range for the deformed alloy, considerably less than the 24 h required for the undeformed alloy, an effect similar to that observed with deformed binary Cu–Ti alloys. Similar observations were made by Dutkiewicz [6] on Cu–2.4 wt.% Ti and Cu–4.29 wt.% Ti alloys, Saji and Hornbogen [7] on a Cu–4 wt.% Ti alloy and Nagarjuna et al. [9] on Cu– 1.5 wt.% Ti and Cu–4.5 wt.% Ti alloys. The early precipitation of the β′ phase due to the presence of high strains and of Cr in super-saturated solid solution also reduced the peak aging times considerably compared with those of the undeformed alloys.
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
355
Fig. 11. TEM images of solution-treated and 90% cold-worked Cu–3Ti–1Cr alloy over-aged at 500 °C for 24 h. (a) BF image; (b) DF image; (c) SAD; (d) schematic of SAD.
The YS and UTS are increased significantly on peak aging the deformed alloy at 400 °C. The precipitation of a very fine dispersion of β′-Cu4Ti precipitate resulted in a 400% increase in the YS of the alloy after 90% deformation and peak aging when compared with that of the alloy in solution-treated condition. At the same time, the corresponding maximum increment in the UTS amounted to only 140%. The percentage increases in the YS and UTS of the present alloy are much higher than those recorded with cold-worked and peak-aged Cu–3Ti– 1Cd alloy (244% and 114%, respectively) [13], an effect attributed to the different solid solution strengthening levels of Cr and Cd in the Cu matrix. The precipitation hardening, in addition to the strain hardening, plays a major role in strengthening the Cu– 3Ti–1Cr alloy. A coherent, ordered and metastable precipitate, the β′-Cu4Ti phase (bct structure) is formed on either a TH04 temper or aging after solution treatment. On aging, the YS and UTS of the solution-treated and undeformed alloy were increased considerably, from 215 to 721 MPa and from 464 to 854 MPa, respectively [14]. The aging of the cold-worked alloy further increased the tensile strength properties over those of the peak-aged undeformed alloy, i.e., 90% deformation and peak aging raised the YS to 1090 MPa and the UTS to 1110 MPa. This indicates that cold working has a considerable effect
in improving the tensile strength properties of Cu–Ti–Cr alloy. The precipitation of the fine, coherent, metastable β′Cu4Ti phase plays a predominant role in improving the tensile strength properties in the cold-worked and peakaged alloys. The mechanical twins and dislocation cells formed during the cold working process also participate in improving the properties significantly. On 90% deformation, the grains are severely deformed and the dislocation density is increased extensively. Because of the sub-grain structure of the alloy, the YS and UTS are increased by an order of 30–50% over those of the peakaged undeformed alloys. This reveals the important role that cold working has in improving the tensile strength properties. The twinning mode of deformation observed in the Cu–3Ti–1Cr alloy is, similar to that noted in earlier reports on the Cu–Ti alloys [8,9,17,18]. 4.3. Solution treatment, cold work and over-aging (ST + CW + OA) Over-aging of the alloy is reflected in decreases in hardness. Aging of 90% cold-worked Cu–3Ti–1Cr alloy at 500 °C for 24 h resulted in a drastic fall in hardness, reaching to a level below that in the solution-treated condition. This was not observed in the undeformed Cu–
356
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
Ti–Cr alloys [14]. This indicates that the matrix in this material (90% deformation, aging for 24 h at 500 °C) has undergone considerable recovery. This can be observed in the TEM images (Fig. 11). These also show the formation of globular (discontinuous) precipitation in the matrix. It was reported earlier that aging at 500 °C of solutiontreated and cold-worked Cu–1.81Be–0.28Co alloy produced faster rates of recovery and sub-grain structure in the matrix of that material [19]. When the deformed Cu–3Ti–1Cr alloy becomes overaged, the formation of incoherent equilibrium precipitates of β-Cu3Ti in the matrix as well as a lamellar-type of discontinuous precipitation at the grain boundaries is noticed. This was observed when the 50% and 90% deformed Cu–3Ti–1Cr alloys were over-aged at 450 °C. When over-aged at 500 °C for 24 h, the lamellar morphology of the discontinuous precipitation was converted to a globular morphology of coarse β surrounded by the α matrix. This means that long-time aging at a high temperature (500 °C) leads to the complete release of the large strains induced by severe deformations, at the same time promoting the dissociation of the lamellar-type discontinuous precipitation into a dual-phase structure of globular coarse precipitate β surrounded by the α matrix. 5. Conclusions The following conclusions are drawn from the present investigations on the age hardening behavior of coldworked Cu–3Ti–1Cr alloy: 1. Substantial hardening occurred during the aging of the deformed alloy at 400 °C. The maximum hardness of 373 Hv; YS of 1090 MPa and UTS of 1110 MPa were obtained on material cold worked 90% then aged for 2 h at 400 °C. 2. The ductility is decreased significantly in the peakaged alloy due to prior cold work, reaching 2.5% elongation on 90% CW + PA at 400 °C. 3. Deformation accelerated the aging, the peak aging time and temperature decreasing significantly to 1 h and 400 °C, respectively, for 90% deformation. 4. For all the deformations, the predominant strengthening mechanism of age hardening is the precipitation of an ordered, metastable and coherent β′-Cu4Ti phase in the matrix in addition to deformation twins and high dislocation density. 5. Over-aging in these alloys is associated with the formation of the equilibrium phase, β-Cu3Ti. In addition, over-aging gave rise to a lamellar form of discontinuous precipitation at lower deformations and globulartype discontinuous precipitation at 90% deformation.
6. The mechanical properties of the Cu–3Ti–1Cr alloy were found to be superior to those of Cu–2.7Ti, Cu– 3Ti–1Cd and the commercial Cu–0.5Be–2.5Co alloys in the cold-worked and peak-aged condition. Acknowledgements The authors greatly acknowledge the Head, Department of Metallurgical Engineering, IT, BHU, Varanasi, for providing laboratory facilities. The authors are thankful to the Research and Development Organisation of Ministry of Defence, Govt. of India, New Delhi for financial support. The authors thank Director, Defence Metallurgical Research Laboratory, Hyderabad, for permission to publish this work. One of the authors (RM) is thankful to Jawaharlal Nehru Technological University, Hyderabad for granting leave under QIP. References [1] Saarivirta MJ, Canon HS. Copper–titanium alloys have high strength. Metal Progr 1959;76:81–4. [2] Cornnie JA, Datta A, Soffa WA. An electron microscopy study of precipitation in Cu–Ti sideband alloys. Metall Trans 1973;4:727–33. [3] Michels HT, Cadoff IB, Levine E. Precipitation hardening in Cu– 3.6wt.% Ti. Metall Trans 1972;3:667–74. [4] Laughlin DE, Cahn JW. Spinodal decomposition in age hardening Cu–Ti alloys. Acta Metall 1975;23:329–39. [5] Datta A, Soffa WA. The structure and properties of age hardened Cu– Ti alloys. Acta Metall 1976;24:987–1001. [6] Dutkiewicz J. Spinodal decomposition, ordering, and discontinuous precipitation in deformed and aged copper–titanium alloys. Met Technol 1978;5:333–40. [7] Saji S, Hornbogen E. Combined recrystallisation and precipitation reactions in a Cu–4wt.%Ti alloy. Z Metall 1978;69:741–6. [8] Nagarjuna S, Balasubramanian K, Sarma DS. Effects of cold work on precipitation hardening of Cu–4.5 mass% Ti alloy. Mater Trans JIM 1995;36(8):1058–66. [9] Nagarjuna S, Balasubramanian K, Sarma DS. Effect of prior cold work on mechanical properties, electrical conductivity and microstructure of aged Cu–Ti alloys. J Mater Sci 1999;34:2929–42. [10] Markandeya R, Nagarjuna S, Sarma DS. Precipitation hardening of Cu–4Ti–1Cd alloy. J Mater Sci 2004;39:1579–87. [11] Markandeya, R., Nagarjuna, S, Sarma, DS. Precipitation hardening of Cu–3Ti–1Cd alloy. Submitted to J Mater Engng and Performance. [12] Markandeya R, Nagarjuna S, Sarma DS. Effect of prior cold work on precipitation hardening of Cu–4Ti–1Cd alloy. J Mater Sci 2006;41:1165. [13] Markandeya R, Nagarjuna S, Sarma DS. Characterization of prior cold worked and age hardened Cu–3Ti–1Cd alloy. Mater Charact 2005;54(5):360–9. [14] Markandeya R, Nagarjuna S, Sarma DS. Precipitation hardening of Cu–Ti–Cr alloys. Mater Sci Eng A 2004;371:291–305. [15] Markandeya R, Nagarjuna S, Sarma DS. Effect of prior cold work on age hardening of Cu–4Ti–1Cr alloy. Mater Sci Eng A 2005;404: 305–13. [16] Properties of wrought copper and copper alloys – C17500 97Cu–0.5Be–2.5Co. In: Metals Hand Book 10th edition, Vol. 2,
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357 Properties and selection: Non-ferrous alloys and pure metals. Materials Park, OH: ASM International, 1990. pp. 288–289. [17] Nagarjuna S, Srinivas M, Balasubramanian K, Sarma DS. On the deformation characteristics of solution treated Cu–Ti alloys. Scr Metall Mater 1995;33:1455–60.
357
[18] Saito K. Deformation twinning in an age-hardened Cu–4 wt.% Ti alloy. J Phys Soc Jpn 1969;27(5):1234–45. [19] Bonfield W, Edwards BC. Precipitation hardening in Cu–1.81% Be–0.28 wt.% Co, Part 3. The effect of prior deformation on the ageing sequence. J Mater Sci 1974;9:415–22.
Effect of prior cold work on age hardening of Cu–3Ti–1Cr alloy R. Markandeya a , S. Nagarjuna b,⁎, D.S. Sarma c a
c
Department of Metallurgical Engineering, College of Engineering, Jawaharlal Nehru Technological University, Kukatpally, Hyderabad-500 072, India b Defence Metallurgical Research Laboratory, Kanchanbagh, Hyderabad-500 058, India Department of Metallurgical Engineering, Institute of Technology, Banaras Hindu University, Varanasi-221 005, India Received 1 August 2005; accepted 28 February 2006
Abstract The influence of 50%, 75% and 90% cold work on the age hardening behavior of Cu–3Ti–1Cr alloy has been investigated by hardness and tensile tests, and light optical and transmission electron microscopy. Hardness increased from 118 Hv in the solution-treated condition to 373 Hv after 90% cold work and peak aging. Cold deformation reduced the peak aging time and temperature of the alloy. The yield strength and ultimate tensile strength reached a maximum of 1090 and 1110 MPa, respectively, following 90% deformation and peak aging. The microstructure of the deformed alloy exhibited elongated grains and deformation twins. The maximum strength on peak aging was obtained due to precipitation of the ordered, metastable and coherent β′-Cu4Ti phase, in addition to high dislocation density and deformation twins. Over-aging resulted in decreases in hardness and strength due to the formation of incoherent and equilibrium β-Cu3Ti phase in the form of a cellular structure. However, the morphology of the discontinuous precipitation changed to a globular form on high deformation. The mechanical properties of Cu–3Ti–1Cr alloy are superior to those of Cu–2.7Ti, Cu–3Ti–1Cd and the commercial Cu–0.5Be–2.5Co alloys in the cold-worked and peak-aged condition. © 2006 Elsevier Inc. All rights reserved. Keywords: Cu–3Ti–1Cr alloy; Precipitation hardening; Prior cold work; Tensile properties; Discontinuous precipitation
1. Introduction World wide-research on the binary Cu–Ti alloys has indicated that they have to serve as substitutes for the expensive and toxic Cu–Be–Co alloys. Extensive research has been carried out by many researchers on the mechanisms of spinodal decomposition and precipitation strengthening in Cu–Ti alloys [1–7]. Nagarjuna et al. have studied the structure–property correlations of Cu–Ti alloys in various conditions, viz., solution ⁎ Corresponding author. Tel.: +91 40 24586409; fax: +91 40 24342123. E-mail address: [email protected] (S. Nagarjuna). 1044-5803/$ - see front matter © 2006 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2006.02.017
treatment, solution treatment + aging and solution treatment + cold work + aging [8,9]. It was reported that compositional modulations would occur during solution treatment itself in high Ti alloys (N4.0 wt.%) and age hardening takes place by the formation of the metastable β′-Cu4Ti phase in both undeformed and deformed alloys. The tensile properties of these alloys are comparable to those of the Cu–Be–Co alloys [8,9]. Our earlier studies [10,11] on the addition of 1 wt.% Cd to a Cu–Ti alloy revealed that the tensile properties of the alloy increased considerably in solution-treated condition due to substitutional solid solution strengthening. The tensile properties of the ternary Cu–Ti–Cd alloys were also comparable to those of a Cu–Be–Co
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
alloy in the undeformed condition. The strengthening mechanism and microstructure of the Cu–Ti–Cd alloys were similar to those of the binary Cu–Ti alloys. Prior cold work increased the tensile properties considerably, but no microstructural changes were reported in the Cu– Ti–Cd alloys during peak aging. The strengthening mechanisms of cold-worked Cu–Ti–Cd alloys [12,13] were also similar to those in undeformed condition. Our recent studies on Cu–Ti alloys with 1 wt.% Cr addition revealed that the tensile properties of the Cu– Ti–Cr alloys increased considerably during peak aging after solution treatment [14]. Age hardening in these alloys was similar to that in the binary Cu–Ti alloys. Further, we recently investigated the effects of prior cold work on the age hardening of a Cu–4Ti–1Cr alloy [15], wherein it has been reported that the properties increased significantly with cold deformation followed by aging. Moreover, the tensile properties are comparable to those of the Cu–Ti–Cd and Cu–2Be–0.6Co alloys. However, little work has been reported on the effects of prior cold work on the age hardening of Cu–Ti–Cr alloys with a low Ti content (b4.0 wt.%). Therefore, the present investigation was undertaken with the aim of filling this gap and to generate necessary data on a Cu– 3Ti–1Cr alloy. This paper presents results on the effects of prior cold work on a Cu–3Ti–1Cr alloy and focuses on the strength properties, microstructure and strengthening mechanism. 2. Experimental procedure A Cu–Ti–Cr alloy was prepared with the aimed composition of 3 wt.% Ti, 1 wt.% Cr and Cu as the balance. A 30 kg melt of the alloy was made by melting the appropriate proportions of pure copper, titanium and chromium in the graphite crucible of Stokes vacuum induction melting (VIM) furnace and then casting into a graphite mould. The ingot was homogenized at 850 °C for 24 h and analyzed for its composition. The analysis report confirmed the composition to be 2.97 wt.% Ti, 0.99 wt.% Cr and balance Cu. The ingot was subsequently hot forged and rolled after soaking at 850 °C to obtain 10 mm thick flats and 12 mm diameter rods. After solutionizing at 860 °C for 2 h, the samples were given different amounts of deformation viz., 50%, 75% and 90%, at room temperature. Vickers hardness at 10 kg load was measured after aging the deformed samples at 400, 450 and 500 °C to establish peak aging time and temperature. A total of 6 indents were taken for each sample at frequent time intervals after aging at the above temperatures and the average value of hardness recorded.
349
Flat tensile specimens as per ASTM E8M standard were prepared from the solution treated + cold deformed + aged alloy, then tested at room temperature and at a strain rate of 10− 3 s− 1, using an Instron universal tensile testing machine. Metallographic specimens were prepared by mechanical polishing then etching in a solution of 10 g K2Cr2O7, 5 ml H2SO4, a few drops of HCl and 95 ml distilled water. The microstructures thus revealed were examined in a light optical microscope (LOM). Thin slices were cut from the cold deformed and aged bulk samples using a low-speed Isomet cutting machine and then mechanically polished to obtain 40–50 μm thick slices. Discs of 3 mm diameter were punched out of these slices and electropolished in a twin-jet electropolisher in a solution of 30 vol.% HNO3 and 70 vol.% methanol at − 45 °C and 20 V. These samples were studied using a JEOL 200CX transmission electron microscope (TEM) operating at 160 kV. 3. Results 3.1. Hardness Fig. 1 shows the effect of prior cold work on the hardness of the Cu–3Ti–1Cr alloy during aging at 400, 450 and 500 °C. These figures also include the variation in hardness with aging time of the undeformed alloy [14] at the above aging temperatures. Cold work prior to aging increased the hardness in solution-treated (ST + CW) as well as peak-aged (ST + CW + PA) conditions. However, peak aging time and temperature decreased with increasing amount of deformation. The peak aging time decreased from 2 h for 50% deformation to 1 h for 90% deformation when aged at 450 °C compared to 24 h for undeformed alloy at the same temperature. The maximum hardness for the Cu–3Ti–1Cr alloy rose from 118 Hv in the undeformed (solution-treated) condition to 373 Hv with 90% deformation and peak aging at 400 °C (Fig. 1a). Similarly, the maximum hardness obtained for 90% deformed alloy after aging at 450 °C was 340 Hv (Fig. 1b) and at 500 °C, 312 Hv (Fig. 1c). The variation of hardness with the amount of prior cold work before and after aging is shown in Fig. 2, which reveals that hardness increased considerably with the amount of prior cold work. Peak hardness was found to be nearly independent of aging temperature for the alloy with deformations of 50% and 75%, but with 90% deformation the peak hardness was highest after aging at 400 °C and thereafter dropped off considerably as the aging temperature increased. Peak hardness after aging at 500 °C is observed to be almost independent of the amount of prior deformation.
350
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
Fig. 2. Effect of prior cold work on hardness of solution-treated and peak-aged Cu–3Ti–1Cr alloy.
3.2. Tensile properties Fig. 3 compares the variation of YS and UTS with the amount of cold work following peak aging at 400 °C for the Cu–3Ti–1Cr alloy of this investigation with the data obtained on a Cu–4Ti–1Cr alloy [15]. For the latter alloy, both the YS and the UTS increased continuously up to 90% deformation. By contrast, in the case of the Cu–3Ti–1Cr alloy, the YS and UTS were nearly constant up to 50% deformation, but then increased slowly up to 75% deformation and reached maxima at 90% deformation. Further, the YS and UTS were nearly the same for 90% deformed and peak-aged Cu–3Ti–1Cr alloy.
Fig. 1. Effect of prior cold work on hardness during aging of Cu–3Ti– 1Cr alloy. (a) 400 °C; (b) 450 °C; (c) 500 °C. Fig. 3. Effect of prior cold work on strength properties of peak-aged Cu–Ti–Cr alloys.
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
351
The mechanical properties of the cold deformed Cu– 3Ti–1Cr alloy and peak aged at 400 °C are compared with those of the binary Cu–2.7Ti, Cu–3Ti–1Cd and Cu– 0.5Be–2.5Co alloys in Table 1. The YS of the Cu–3Ti–1Cr alloy increased from 757 MPa with 50% deformation to 1090 MPa with 90% deformation while the UTS increased from 890 MPa with 50% deformation to 1110 MPa on 90% deformation compared to 721 MPa YS and 854 MPa UTS in the undeformed and peak-aged condition. The ductility (% elongation) of the Cu–3Ti–1Cr alloy decreased from 17 for 50% cold work to 2.5 for 90% cold work compared to 21 in peak-aged condition for undeformed alloy. The YS and UTS are higher for the Cu–3Ti–1Cr alloy in 90% coldwork and aged condition than those of Cu–2.7Ti [8] and Cu–3Ti–1Cd [13] alloys in a similar condition. Further, the mechanical properties of the Cu–3Ti–1Cr alloy are superior to those of the commercial Cu–0.5Be–2.5Co alloy [16] in the cold-worked and peak-aged condition. 3.3. Light optical microscopy The microstructure of the undeformed Cu–3Ti–1Cr alloy [14] as seen in the LOM is briefly described here for comparison. In the solution-treated condition it displayed an equiaxed single-phase structure with fine grains. On peak aging at 450 °C for 24 h, the alloy exhibited a similar microstructure. However, it revealed onset of discontinuous precipitation at the grain boundaries of the matrix when over-aged, either at 450 °C for 88 h or 500 °C for 24 h. Fig. 4(a), (b) and (c) shows, respectively, the microstructures of the solution-treated and 50% cold-worked Cu–3Ti–1Cr alloy in peak-aged (400 °C/2 h), over-aged Table 1 Comparison of mechanical properties of solution-treated (ST) + coldworked (CW) + peak-aged (PA) Cu–3Ti–1Cr alloy with those reported for Cu–2.7Ti [9], Cu–3Ti–1Cd [13] and Cu–0.5Be–2Co [16] alloys Property
ST
Cold-worked Cu– Cu– 3Ti–1Cr, peak aged at 2.7Ti 400 °C (time) [9]
Cu– 3Ti– 1Cd [13]
Cu– 0.5Be– 2.5Co [16]
0% 50% 90% CW CW CW (24 h) (2 h) (1 h)
ST + 90% CW + PA
ST + 90% CW + PA
ST + full hard + aged (TH04 temper)
950 1000 355
922 1035 375
YS, MPa 215 721 UTS, MPa 464 854 Hardness, 118 277 Hv Elongation, %
43
21
757 890 300
17
1090 1110 373
2.5
3.5
690–825 760–895 220–270 (95–102 HRB) 6.0 10–20
Fig. 4. Light optical micrographs of solution-treated and 50% coldworked Cu–3Ti–1Cr alloy. (a) Peak aged at 400 °C/2 h; (b) over-aged at 450 °C/8 h; (c) over-aged at 500 °C/24 h.
1 (450 °C for 8 h) and over-aged 2 (500 °C for 24 h) conditions. The equiaxed grains were disturbed by the cold deformation and the alloy exhibited a single-phase structure with severe strain markings in the peak-aged condition (Fig. 4(a)). The deformation structure recovered slightly on over-aging at 450 °C for 8 h and the occasional occurrence of discontinuous precipitation was observed at the grain boundaries of the matrix (Fig. 4(b)). The discontinuous precipitation grew to a large extent on over-aging at 500 °C for 24 h (Fig. 4(c)). On 90% cold work, the grain elongation was severe, as revealed in Fig. 5. When peak aged at 400 °C for 1 h, a deformed grain structure along with the onset of discontinuous precipitation was seen (Fig. 5(a)). On over-
352
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
aging at 450 °C for 8 h, elongated grains and a significant amount of discontinuous precipitation were observed (Fig. 5(b)), while a substantial amount of the matrix is covered by discontinuous precipitation on over-aging at 500 °C for 24 h (Fig. 5(c)). 3.4. Transmission electron microscopy TEM studies were carried out on 50% and 90% deformed Cu–3Ti–1Cr alloy to understand the structural changes that occurred during peak aging and over-aging after high deformations. Fig. 6, a bright field (BF) image of the Cu–3Ti–1Cr alloy in the peak-aged condition (400 °C for 2 h) after
Fig. 6. TEM BF image of solution-treated and 50% cold-worked Cu– 3Ti–1Cr alloy peak aged at 400 °C for 2 h, showing fine precipitate of β′ Cu4Ti, twins and dislocation cells.
Fig. 5. Light optical micrographs of solution-treated and 90% coldworked Cu–3Ti–1Cr alloy. (a) Peak aged at 400 °C/1 h; (b) over-aged at 450 °C/8 h; (c) over-aged at 500 °C/24 h.
50% prior cold work, reveals dislocation cells and deformation twins. In addition, an ordered, metastable and coherent precipitate, β′-Cu4Ti, was seen. The selected area diffraction (SAD) pattern and its schematic diagram to identify the precipitate are not shown here because the same was reported in our recent work on undeformed Cu–Ti–Cr alloys [14], wherein the metastable and coherent phase was identified as β′-Cu4Ti precipitate. The β′ phase is reported to have a body-centered tetragonal (bct) crystal structure [5]. Deformation twins are more clearly revealed in the BF and dark field (DF) images of the 50% deformed alloy over-aged at 450 °C for 8 h (Fig. 7(a) and (b), respectively). The SAD pattern and its schematic diagram (shown in Fig. 7(c) and (d), respectively) identify and confirm the presence of twins. Fig. 7(a) also reveals the formation of some discontinuous precipitation surrounded by a highly strained matrix. Fig. 8 shows the BF images of the alloy with 50% cold deformation and over-aged at 500 °C for 24 h. Fig. 8(a) shows the coarse β-Cu3Ti equilibrium precipitate formed on over-aging while Fig. 8(b) shows the highly strained matrix and some fine β precipitation. The SAD pattern and its schematic to identify the precipitate β-Cu3Ti are not included here as they are the same as depicted in Fig. 11. Fig. 9, a BF image of 90% deformed Cu–3Ti–1Cr alloy peak aged at 400 °C, reveals increased twin density and dislocation cell structure. The ordered, metastable and coherent β′-Cu4Ti precipitate was also observed here. The deformation twins in the matrix of the 90% deformed Cu–3Ti–1Cr alloy over-aged at 450 °C for 8 h are clearly visible in the BF and DF images of Fig. 10(a) and (b). The SAD in Fig. 10(c) and its schematic in Fig. 10(d) confirm the presence of twins in the alloy. The dual phases of matrix and coarse β-Cu3Ti precipitate revealed in the BF and DF images of Fig. 11(a) and (b), respectively, are from
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
353
Fig. 7. TEM images of solution-treated and 50% cold-worked Cu–3Ti–1Cr alloy over-aged at 450 °C for 8 h. (a) BF image; (b) DF image; (c) SAD; (d) schematic of SAD of deformation twins.
this alloy after over-aging at 500 °C for 24 h. Interphase boundary dislocations can also be seen in the α-matrix (Fig. 11(a)). The SAD of Fig. 11(c) and its schematic (Fig. 11(d)) confirm the presence of equilibrium precipitate, β. 4. Discussion The present investigation has revealed that the hardness and strength of solution-treated Cu–3Ti–1Cr alloy
Fig. 8. TEM BF images of solution-treated and 50% cold-worked Cu– 3Ti–1Cr alloy over-aged at 500 °C for 24 h, showing (a) matrix and coarse β precipitate and (b) matrix with very fine β precipitate.
Fig. 9. TEM BF image of solution-treated and 90% cold-worked Cu– 3Ti–1Cr alloy peak aged at 400 °C for 1 h, showing fine precipitate of β′ Cu4Ti, deformation twins and dislocation cells.
354
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
Fig. 10. TEM images of solution-treated and 90% cold-worked Cu–3Ti–1Cr alloy over-aged at 450 °C for 8 h. (a) BF image; (b) DF image; (c) SAD; (d) schematic of SAD.
are significantly improved by cold work and subsequent peak aging at 400 °C. 4.1. Solution treatment and cold work (ST + CW) Cold rolling to a 90% reduction in thickness increased the hardness of the solution-treated Cu–3Ti–1Cr alloy (118 Hv) by 115%, to 254 Hv. By comparison, the same processing of a Cu–4Ti–1Cr alloy resulted in only a 40% increase, but in that case the initial solution-treated hardness was 224 Hv [14]. The higher initial hardness of the Cu–4Ti–1Cr alloy is attributed to compositional modulations due to the higher titanium content [5] and the presence of β′-Cu4Ti precipitate, whereas the solutiontreated Cu–3Ti–1Cr alloy is characterized by the presence of a single-phase (homogeneous) structure. Thus the higher percentage increase in hardness due to the application of cold working seen in the Cu–3Ti–1Cr alloy is attributed to the greater effect of the dislocation networks formed during the deformation process. 4.2. Solution treatment, cold work and peak aging (ST + CW + PA) The hardness of the Cu–3Ti–1Cr alloy increased considerably on cold working followed by aging,
reaching a maximum hardness of 373 Hv after 90% deformation and aging for 1 h at 400 °C. The highest hardness values for all the deformations were obtained when the alloy was aged at 400 °C whereas the undeformed alloy exhibited its peak hardness after aging at 450 °C [14]. This indicates that cold deformation reduced the peak aging temperature, which is in agreement with the findings on the binary Cu–Ti alloys [9]. Because of the presence of Cr in a super-saturated solid solution and to the extensive deformation, the aging kinetics were accelerated, giving rise to maximum strength values at a lower aging temperature (400 °C) and promoting the precipitation of a very fine metastable β′ phase in increased volume fraction. The peak aging times were generally in the 1–2 h range for the deformed alloy, considerably less than the 24 h required for the undeformed alloy, an effect similar to that observed with deformed binary Cu–Ti alloys. Similar observations were made by Dutkiewicz [6] on Cu–2.4 wt.% Ti and Cu–4.29 wt.% Ti alloys, Saji and Hornbogen [7] on a Cu–4 wt.% Ti alloy and Nagarjuna et al. [9] on Cu– 1.5 wt.% Ti and Cu–4.5 wt.% Ti alloys. The early precipitation of the β′ phase due to the presence of high strains and of Cr in super-saturated solid solution also reduced the peak aging times considerably compared with those of the undeformed alloys.
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
355
Fig. 11. TEM images of solution-treated and 90% cold-worked Cu–3Ti–1Cr alloy over-aged at 500 °C for 24 h. (a) BF image; (b) DF image; (c) SAD; (d) schematic of SAD.
The YS and UTS are increased significantly on peak aging the deformed alloy at 400 °C. The precipitation of a very fine dispersion of β′-Cu4Ti precipitate resulted in a 400% increase in the YS of the alloy after 90% deformation and peak aging when compared with that of the alloy in solution-treated condition. At the same time, the corresponding maximum increment in the UTS amounted to only 140%. The percentage increases in the YS and UTS of the present alloy are much higher than those recorded with cold-worked and peak-aged Cu–3Ti– 1Cd alloy (244% and 114%, respectively) [13], an effect attributed to the different solid solution strengthening levels of Cr and Cd in the Cu matrix. The precipitation hardening, in addition to the strain hardening, plays a major role in strengthening the Cu– 3Ti–1Cr alloy. A coherent, ordered and metastable precipitate, the β′-Cu4Ti phase (bct structure) is formed on either a TH04 temper or aging after solution treatment. On aging, the YS and UTS of the solution-treated and undeformed alloy were increased considerably, from 215 to 721 MPa and from 464 to 854 MPa, respectively [14]. The aging of the cold-worked alloy further increased the tensile strength properties over those of the peak-aged undeformed alloy, i.e., 90% deformation and peak aging raised the YS to 1090 MPa and the UTS to 1110 MPa. This indicates that cold working has a considerable effect
in improving the tensile strength properties of Cu–Ti–Cr alloy. The precipitation of the fine, coherent, metastable β′Cu4Ti phase plays a predominant role in improving the tensile strength properties in the cold-worked and peakaged alloys. The mechanical twins and dislocation cells formed during the cold working process also participate in improving the properties significantly. On 90% deformation, the grains are severely deformed and the dislocation density is increased extensively. Because of the sub-grain structure of the alloy, the YS and UTS are increased by an order of 30–50% over those of the peakaged undeformed alloys. This reveals the important role that cold working has in improving the tensile strength properties. The twinning mode of deformation observed in the Cu–3Ti–1Cr alloy is, similar to that noted in earlier reports on the Cu–Ti alloys [8,9,17,18]. 4.3. Solution treatment, cold work and over-aging (ST + CW + OA) Over-aging of the alloy is reflected in decreases in hardness. Aging of 90% cold-worked Cu–3Ti–1Cr alloy at 500 °C for 24 h resulted in a drastic fall in hardness, reaching to a level below that in the solution-treated condition. This was not observed in the undeformed Cu–
356
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357
Ti–Cr alloys [14]. This indicates that the matrix in this material (90% deformation, aging for 24 h at 500 °C) has undergone considerable recovery. This can be observed in the TEM images (Fig. 11). These also show the formation of globular (discontinuous) precipitation in the matrix. It was reported earlier that aging at 500 °C of solutiontreated and cold-worked Cu–1.81Be–0.28Co alloy produced faster rates of recovery and sub-grain structure in the matrix of that material [19]. When the deformed Cu–3Ti–1Cr alloy becomes overaged, the formation of incoherent equilibrium precipitates of β-Cu3Ti in the matrix as well as a lamellar-type of discontinuous precipitation at the grain boundaries is noticed. This was observed when the 50% and 90% deformed Cu–3Ti–1Cr alloys were over-aged at 450 °C. When over-aged at 500 °C for 24 h, the lamellar morphology of the discontinuous precipitation was converted to a globular morphology of coarse β surrounded by the α matrix. This means that long-time aging at a high temperature (500 °C) leads to the complete release of the large strains induced by severe deformations, at the same time promoting the dissociation of the lamellar-type discontinuous precipitation into a dual-phase structure of globular coarse precipitate β surrounded by the α matrix. 5. Conclusions The following conclusions are drawn from the present investigations on the age hardening behavior of coldworked Cu–3Ti–1Cr alloy: 1. Substantial hardening occurred during the aging of the deformed alloy at 400 °C. The maximum hardness of 373 Hv; YS of 1090 MPa and UTS of 1110 MPa were obtained on material cold worked 90% then aged for 2 h at 400 °C. 2. The ductility is decreased significantly in the peakaged alloy due to prior cold work, reaching 2.5% elongation on 90% CW + PA at 400 °C. 3. Deformation accelerated the aging, the peak aging time and temperature decreasing significantly to 1 h and 400 °C, respectively, for 90% deformation. 4. For all the deformations, the predominant strengthening mechanism of age hardening is the precipitation of an ordered, metastable and coherent β′-Cu4Ti phase in the matrix in addition to deformation twins and high dislocation density. 5. Over-aging in these alloys is associated with the formation of the equilibrium phase, β-Cu3Ti. In addition, over-aging gave rise to a lamellar form of discontinuous precipitation at lower deformations and globulartype discontinuous precipitation at 90% deformation.
6. The mechanical properties of the Cu–3Ti–1Cr alloy were found to be superior to those of Cu–2.7Ti, Cu– 3Ti–1Cd and the commercial Cu–0.5Be–2.5Co alloys in the cold-worked and peak-aged condition. Acknowledgements The authors greatly acknowledge the Head, Department of Metallurgical Engineering, IT, BHU, Varanasi, for providing laboratory facilities. The authors are thankful to the Research and Development Organisation of Ministry of Defence, Govt. of India, New Delhi for financial support. The authors thank Director, Defence Metallurgical Research Laboratory, Hyderabad, for permission to publish this work. One of the authors (RM) is thankful to Jawaharlal Nehru Technological University, Hyderabad for granting leave under QIP. References [1] Saarivirta MJ, Canon HS. Copper–titanium alloys have high strength. Metal Progr 1959;76:81–4. [2] Cornnie JA, Datta A, Soffa WA. An electron microscopy study of precipitation in Cu–Ti sideband alloys. Metall Trans 1973;4:727–33. [3] Michels HT, Cadoff IB, Levine E. Precipitation hardening in Cu– 3.6wt.% Ti. Metall Trans 1972;3:667–74. [4] Laughlin DE, Cahn JW. Spinodal decomposition in age hardening Cu–Ti alloys. Acta Metall 1975;23:329–39. [5] Datta A, Soffa WA. The structure and properties of age hardened Cu– Ti alloys. Acta Metall 1976;24:987–1001. [6] Dutkiewicz J. Spinodal decomposition, ordering, and discontinuous precipitation in deformed and aged copper–titanium alloys. Met Technol 1978;5:333–40. [7] Saji S, Hornbogen E. Combined recrystallisation and precipitation reactions in a Cu–4wt.%Ti alloy. Z Metall 1978;69:741–6. [8] Nagarjuna S, Balasubramanian K, Sarma DS. Effects of cold work on precipitation hardening of Cu–4.5 mass% Ti alloy. Mater Trans JIM 1995;36(8):1058–66. [9] Nagarjuna S, Balasubramanian K, Sarma DS. Effect of prior cold work on mechanical properties, electrical conductivity and microstructure of aged Cu–Ti alloys. J Mater Sci 1999;34:2929–42. [10] Markandeya R, Nagarjuna S, Sarma DS. Precipitation hardening of Cu–4Ti–1Cd alloy. J Mater Sci 2004;39:1579–87. [11] Markandeya, R., Nagarjuna, S, Sarma, DS. Precipitation hardening of Cu–3Ti–1Cd alloy. Submitted to J Mater Engng and Performance. [12] Markandeya R, Nagarjuna S, Sarma DS. Effect of prior cold work on precipitation hardening of Cu–4Ti–1Cd alloy. J Mater Sci 2006;41:1165. [13] Markandeya R, Nagarjuna S, Sarma DS. Characterization of prior cold worked and age hardened Cu–3Ti–1Cd alloy. Mater Charact 2005;54(5):360–9. [14] Markandeya R, Nagarjuna S, Sarma DS. Precipitation hardening of Cu–Ti–Cr alloys. Mater Sci Eng A 2004;371:291–305. [15] Markandeya R, Nagarjuna S, Sarma DS. Effect of prior cold work on age hardening of Cu–4Ti–1Cr alloy. Mater Sci Eng A 2005;404: 305–13. [16] Properties of wrought copper and copper alloys – C17500 97Cu–0.5Be–2.5Co. In: Metals Hand Book 10th edition, Vol. 2,
R. Markandeya et al. / Materials Characterization 57 (2006) 348–357 Properties and selection: Non-ferrous alloys and pure metals. Materials Park, OH: ASM International, 1990. pp. 288–289. [17] Nagarjuna S, Srinivas M, Balasubramanian K, Sarma DS. On the deformation characteristics of solution treated Cu–Ti alloys. Scr Metall Mater 1995;33:1455–60.
357
[18] Saito K. Deformation twinning in an age-hardened Cu–4 wt.% Ti alloy. J Phys Soc Jpn 1969;27(5):1234–45. [19] Bonfield W, Edwards BC. Precipitation hardening in Cu–1.81% Be–0.28 wt.% Co, Part 3. The effect of prior deformation on the ageing sequence. J Mater Sci 1974;9:415–22.