On the variation of mechanical properties with solute content in Cu–Ti alloys

Materials Science and Engineering A259 (1999) 34 – 42

On the variation of mechanical properties with solute content in Cu–Ti alloys S. Nagarjuna a,*, M. Srinivas a, K. Balasubramanian b, D.S. Sarma c a

b

Defence Metallurgical Research Laboratory, Kanchanbagh-PO, Hyderabad-500 058, India Non-Ferrous Materials Technology De6elopment Centre, Kanchanbagh-PO, Hyderabad-500 058, India c Department of Metallurgical Engineering, Banaras Hindu Uni6ersity, Varanasi-221 005, India Received 21 October 1997; received in revised form 11 August 1998

Abstract The variation of mechanical properties and electrical conductivity of Cu – Ti alloys of four compositions, viz. Cu –1.5 wt%Ti, Cu–2.7 wt%Ti, Cu–4.5 wt%Ti, and Cu–5.4 wt%Ti, have been studied in solution treated (ST), solution treated +peak aged (ST+ PA), and solution treated +cold worked + peak aged (ST + CW +PA) conditions. In the ST condition, Ti is found to be a potential solid solution strengthener of copper showing greater effect than other elements like Zn, Ni, Al, Si, Be, and Sn. Solid solution strengthening in Cu–Ti alloys is attributed to the interaction of titanium atoms with screw dislocations and the effective interaction is more due to modulus mismatch than size misfit. Further, a marked change in the linear variation of tensile strength and elongation with Ti content is observed at about 4.0 wt%Ti beyond which, tensile strength increases sharply while elongation decreases further, which is attributed to fine scale precipitation formed during quenching of Cu – 4.5 Ti and Cu – 5.4 Ti alloys. On the other hand, hardness and tensile properties increase linearly up to 5.4 wt%Ti in the peak aged condition with or without prior cold work, due to uniform precipitation of Cu4Ti, b l phase in all the four alloys. The increase in yield and tensile strengths due to solid solution strengthening, cold work, and precipitation have been determined quantitatively in ST + CW +PA alloys. While electrical conductivity is less, the mechanical properties of Cu – Ti alloys are comparable with those of commercial Cu –Be alloys. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Cu – Ti alloys; Solid solution strengthening; Prior cold work; Ageing; Cu4Ti, b l precipitate; Mechanical properties and electrical conductivity

1. Introduction Binary Cu–Ti alloys have good potential as a substitute for expensive and toxic Cu – Be alloys. The mechanisms of spinodal decomposition and precipitation strengthening in Cu – Ti alloys have been studied extensively [1–9]. It was reported earlier by us that hardness and yield strength in solution treated (ST) Cu – Ti alloys increase linearly up to about 4.0 wt%Ti beyond which, a sharp increase is observed with further additions of Ti, due to fine scale precipitation in the form of modulations in Cu–4.5 Ti and Cu4Ti, b l precipitate in Cu– 5.4 Ti alloy formed during quenching [10]. Further, a similar behaviour has been observed in the case of fatigue strength [11] and electrical resistivity [12] as * Corresponding author. Tel.: +91 40 4440051/4442408; fax: + 91 40 4440683.

well. On the other hand, variation of fatigue strength and electrical resistivity in peak aged (PA) condition has been reported to be linear up to 5.4 wt%Ti and this has been attributed to uniform precipitation of ordered, metastable, and coherent Cu4Ti, b l phase in all the four Cu–Ti alloys [11,12]. However, little work has been reported on the variation of tensile strength and elongation in ST condition and mechanical properties (hardness, yield and tensile strengths, and elongation) in aged condition with and without prior cold work (CW). The aim of the present study has been to investigate the variation of mechanical properties with Ti content and correlation of yield and tensile strengths with volume fraction of the Cu4Ti, b l precipitate in Cu–Ti alloys. The results obtained on the effect of Ti content on hardness and tensile properties in ST, ST+PA, and ST+CW+ PA conditions are presented in this paper. Further, tensile strength and electrical conductivity of

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Cu–Ti alloys are compared with those of Cu–Be alloys.

2. Experimental A 30 kg melt of each of the four Cu – Ti alloys with the nominal composition (in wt%) of 1.5, 3.0, 4.3, and 5.5 Ti have been made in a Stokes vacuum induction melting (VIM) furnace with oxygen free electronic (OFE) copper and Cu – 26 wt%Ti master alloy as charge materials. The ingots were homogenised at 850°C for 24 h and analysed for Ti content. The analysed composition of the ingots (in wt%Ti) is 1.5, 2.7, 4.5, and 5.4 and oxygen (in ppm) is 5.7, 6.0, 6.6, and 6.4, balance being copper. The homogenised ingots were hot forged and rolled at 850°C into 10 mm thick flats and 12 mm F rods. Specimens from the hot rolled flats were ST at different temperatures to achieve a constant grain size (75 mm), then aged at different temperatures for different times and their hardness (HV) was measured. The details of solution treatment and ageing (time, temperature, percentage cold work, and alloy composition) have been published elsewhere [7 – 10]. Cylindrical tensile samples having 25 mm gauge length and 4.0 mm gauge diameter were tested for tensile properties in ST as well as PA (450°C/16 h) conditions, at an ambient temperature and at a nominal strain rate of 10 − 3 s − 1 using INSTRON 1185 ball-screw driven universal testing system. The ST specimens were cold rolled giving 80% deformation. Hardness (HV) of the cold rolled specimens was measured after ageing at different temperatures for different times. Flat tensile samples of 25 mm gauge length, 6 mm width, and 1 mm thickness were made according to the ASTM specification E 8M-89b (sub-size specimen) [13] from the cold rolled strips with 90% deformation, PA at 400°C, and tested for tensile properties.

Fig. 1. Variation of hardness and yield strength with Ti content in Cu – Ti alloys [10].

3. Results and discussion

3.1. Solid solution strengthening The variation of hardness and yield strength with Ti content in ST Cu – Ti alloys as reported earlier [10], is reproduced in Fig. 1. Both hardness and yield strength increased linearly up to about 4.0 wt%Ti beyond which, a sharp increase was observed with further additions of titanium. The sharp rise in hardness and yield strength was attributed to fine scale precipitation in the form of modulations in Cu – 4.5 Ti and Cu4Ti, b l precipitate in Cu–5.4 Ti alloy formed during quenching itself. The variation of 1.0% proof stress as a function of solute content (at%), is plotted for ST Cu – 1.5 Ti and Cu–2.7 Ti alloys and compared with the data reported in

Fig. 2. Comparison of solid solution strengthening in Cu–Ti with other Cu base alloys [14].

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Table 1 Values of constant A in Eq. (1) for different Cu alloys Alloy

A

Electronic configuration of the solute

Difference in valence

Cu Cu–Zn Cu–Ni Cu–Al Cu–Si Cu–Be Cu–Sn Cu–Ti

— 0.12 0.26 0.35 0.44 0.86 1.76 5.00

(Ar) 3d104s1 (Ar) 3d104s2 (Ar) 3d84s2 (Ne) 3s23p1 (Ne) 3s23p2 1s22s2 (Kr) 4d10 5s25p2 (Ar) 3d24s2

— 1 2 2 3 3 3 3

Electronic configuration of: Ne: 1s22s22p6; Ar: 1s22s22p63s23p6; Kr: 1s22s22p63s23p63d104s24p6. Table 2 Misfit parameters for Cu alloys Alloy

oG%

ob

os = oG%−3ob

% of effect due to size

Ref.

Cu–Zn Cu–Ni Cu–Al Cu–Si Cu–Sn Cu–Ti

−0.38 +0.48 −0.61 −0.76 −1.18 −1.65

+0.056 −0.031 +0.064 +0.020 +0.282 +0.105

0.55 0.57 0.80 0.82 2.03 1.97

13 16 24 07 42 16

Fleischer [16] Fleischer [16] Fleischer [16] Fleischer [16] Fleischer [16] Present work

oG% = 1/G(dG/dc)(1+1/2G · dG/dc )−1 where oG% is Modulus mismatch parameter; G, the shear modulus; c, the solute content in at%. ob =1/b(db/dc) where ob is size misfit parameter; b, the lattice parameter; c, the solute content in at%. os = sum of mismatch parameters

literature [14] for other Cu base alloys such as Cu–Zn, Cu–Ni, Cu–Al, Cu – Si, Cu – Be, and Cu – Sn in Fig. 2. The linear increase in proof stress with the solute content can be represented as follows: sa = sCu + A · c

(1)

where sa and sCu are the 1.0% proof stress of the alloy and pure copper, c, the concentration of solute element in at%, and A, the slope of the straight line. The value of A, listed in Table 1, shows that the slope A for Ti is the highest (5.00) and decreases for Sn, Be, Si, Al, Ni, and Zn, with the lowest value of 0.12 for Zn, thus indicating that strengthening of copper due to Ti additions is quite significant. Solid solution strengthening in dilute substitutional solid solutions was related to atom size misfit parameter [15], effective modulus [16], and difference in valence of solute and solvent metals [17]. The atom size misfit parameter (ob = 1/b(db/dc) calculated for Cu – Ti alloys is +0.105 and the modulus mismatch parameter [oG% = 1/G(dG/dc)(1+1/2G · dG/dc ) − 1] is − 1.65. The ob and oG% values for various other solutes in copper reported by Fleischer [16] are presented in Table 2 and compared with Cu – Ti solid solution. The percentage of effect in solid solution strengthening in Cu – Ti alloys due to size misfit is 16% which is in line with other solid solutions (Table 2). Fleischer [16], after studying the relation between hardening and different misfit parameters, concluded that modulus mismatch parameter con-

tributes to a larger extent than atom size misfit parameter to the solid solution strengthening, and the solution hardening is the result of interaction of solute atoms with screw dislocations. It is evident from Table 2 that the behaviour of Cu–Ti alloys is similar to that of other Cu alloys. Hence, it can be concluded that solid solution strengthening in Cu–Ti alloys could be due to interaction of titanium atoms with screw dislocations and the effective interaction is more due to modulus mismatch than size misfit. The fact that though Sn and Ti are equally effective in terms of os the sum of mismatch parameters, i.e.(oG% − 3ob), as shown in Table 2, Ti is more effective than Sn in solid solution strengthening (Fig. 2). This suggests that other factors like tendency to form clusters (which is well established in high Ti alloys due to spinodal decomposition) also contributes significantly to solid solution strengthening in Cu–Ti alloys. Table 1 also lists the values of the difference in valence for the above copper alloys. It is found that the slope A increases with increasing value of the difference in valence, which is in agreement with the behaviour reported for electrical resistivity in Cu–Ti alloys [12].

3.2. Variation of mechanical properties in aged conditions The variation of hardness, yield and tensile strength, and elongation plotted as a function of alloying content

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in ST, ST +PA, and ST +CW + PA conditions is shown in Figs. 3 – 6, respectively. It is evident from Figs. 3–5 that in solution treated condition, the properties have shown a bi-linear behaviour with a transition point at about 4.0 wt%Ti. The sharp rise in hardness and strength was attributed to fine scale precipitation in the form of modulations in Cu – 4.5 Ti and Cu4Ti, b l precipitate in Cu– 5.4 Ti alloy formed during quenching itself [10]. On the other hand, a single linear relationship, i.e. gradual increase in properties with increasing titanium content, has been observed in ST+PA and

Fig. 5. Effect of Ti content on tensile strength of Cu –Ti alloys.

Fig. 3. Variation of hardness with Ti content in Cu–Ti alloys.

Fig. 6. Effect of Ti content on elongation in Cu – Ti alloys.

Fig. 4. Variation of yield strength with Ti content in Cu– Ti alloys.

ST+CW+ PA conditions. The difference in behaviour could be attributed to the uniform precipitation of ordered, metastable, and coherent Cu4Ti, b l phase in the four alloys [11,12]. In the precipitation hardening alloys, there are three causes of hardening[18]: (i) coherency strain hardening; (ii) chemical hardening i.e. when the dislocation cuts through the precipitate; (iii) dispersion hardening, i.e. when the dislocation goes round or over the precipitate. The precipitation of particles having a slight misfit in the matrix gives rise to stress fields which hinder the

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movement of gliding dislocations. For the dislocations to pass through the internal stress fields, additional stress is required and this is termed as coherency strain hardening. When the precipitate particles are small in size and coherent, they are deformable and the dislocations cut and pass through them, and the resultant hardening is known as chemical hardening. When the precipitates become coarse in the later stages of precipitation, dislocation line moves between the precipitate particles to form loops around them and the hardening is known as dispersion hardening. In the PA Cu–Ti alloys, the maximum strength is achieved as a result of both coherency strain hardening and chemical hardening, while dispersion hardening contributes to the strength in the over aged condition. The ductility parameter, i.e. elongation, has also shown a bi-linear relationship with a sharp drop beyond 4.0 wt%Ti in ST condition (Fig. 6), while it decreased linearly and gradually up to 5.4 wt%Ti in ST + PA and ST+CW +PA conditions. The sharp decrease in ductility beyond 4.0 wt%Ti in ST condition could be due to the modulated structure and Cu4Ti, b l precipitate. The decrease in elongation in the aged alloys is nearly 35 and 90%, respectively, in ST+ PA and ST+CW+ PA conditions. The decrease in elongation in the aged conditions is attributed to the precipitation of Cu4Ti, b l phase. The SEM fractographs of the four alloys in ST +PA and ST + CW+ PA conditions are shown in Figs. 7 and 8, respectively. It is evident from the fractographs that the failure is by dimple rupture and equiaxed dimples of varied size and depth are seen in all cases. The large dimples in Fig. 7a and Fig. 7b contain ripple marks and tear ridges on their

surfaces which could be due to void growth by tearing. Alloys having low ductility (Fig. 7c and Fig. 7d) show shallow and small equiaxed dimples without ridge markings indicating that the linking of voids has taken place at relatively small strains due to large population of precipitates. In ST + CW+ PA condition (Fig. 8), all alloys have shown shallow and equiaxed dimples, as compared to the as-quenched and peak aged condition (Fig. 7). Further, it is to be noted that the void size decreases with increasing alloy content which can again be related to the increasing volume fraction of the precipitates with increasing Ti content.

3.3. Correlation of strength with 6olume fraction of Cu4Ti, b 1 precipitate Fig. 9 shows the variation of volume fraction of Cu4Ti, b l precipitate with Ti content. The volume fraction of the precipitate was computed using experimentally measured density values, as reported earlier [12] and are plotted with Ti content here (there could be a small error in the values of the volume fraction of the precipitate due to the possible errors in the measurement of densities). The volume fraction of the precipitate increased linearly with Ti content in both ST+PA and ST + CW+ PA conditions. The volume fraction of the precipitate is found to be higher in PA alloys with prior cold deformation than that of as-quenched and PA ones. The observed behaviour is based on two fundamental phenomena: (i) deformation characteristics; (ii) heterogeneous nucleation. According to Datta and Soffa [3], during ageing of the solution treated alloy, homogeneous and continuous precipitation oc-

Fig. 7. Fractographs of as-quenched and peak aged Cu – Ti alloys.(a) Cu – 1.5 Ti; (b) Cu – 2.7 Ti; (c) Cu – 4.5 Ti; (d) Cu – 5.4 Ti.

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Fig. 8. Fractographs of as-quenched, cold worked, and peak aged Cu – Ti alloys.(a) Cu – 1.5 Ti; (b) Cu – 2.7 Ti; (c) Cu – 4.5 Ti; (d) Cu–5.4 Ti.

curs and heterogeneous nucleation on structural singularities such as grain boundaries and dislocations is not observed in a Cu – 4.0 wt%Ti alloy. However, Dutkiewicz [5] concluded that in a highly deformed Cu–4.29 wt%Ti alloy, heterogeneous precipitation of transitional b l phase occurs on dislocations and other structural defects. Therefore, nucleation of the intermediate precipitate in deformed Cu – Ti alloys occurs at crystal defects such as dislocations that act as heterogeneous sites. The quantity of precipitate nucleated is proportional to the number of nucleation sites (i.e. number of dislocations) available. Since the number of dislocations is increased by means of cold plastic deformation according to the equation: so =si +aGbr 1/2, the quantity or volume fraction of the precipitate nucleated also increases. An annealed metal contains about 106 –108 dislocations /cm2, while a severely plastically deformed metal contains about 1012 dislocations /cm2 [19]. Now, in the present case, the 90% cold deformation given to the as-quenched Cu – Ti alloys by means of cold rolling, increases the dislocation density to about 1012 dislocations /cm2. Consequently after ageing, the quantity (volume fraction) of the precipitate nucleated will also be higher, i.e. proportional to 1012 dislocations /cm2. Hence, our observation that the volume fraction of the precipitate in PA alloys with prior cold deformation is higher than that of as-quenched ones is justified. When the straight lines are extrapolated backwards, the line corresponding to PA condition meets the X-axis at the point (0.55, 0) and that of CW +PA condition at (0.45, 0) indicating that the boundary composition limit of Ti content, below which,

precipitation hardening does not occur in Cu–Ti alloys. This observation is in agreement with the reported results of Saarivirta and Cannon [1] on mechanical properties and Nagarjuna et al. [12] on resistivity. The dependence of yield and tensile strengths on volume fraction of Cu4Ti, b l precipitate is shown in Fig. 10. Yield as well as tensile strength increased linearly with the volume fraction of the precipitate in both ST + PA as well as ST + CW+ PA conditions. It is interesting to note from Fig. 10 that the rate of increase in strength with volume fraction of Cu4Ti, b l

Fig. 9. Variation of volume fraction of Cu4Ti, b l precipitate with Ti content in Cu – Ti alloys.

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Fig. 10. Variation of yield and tensile strength with volume fraction of Cu4Ti, b l precipitate in Cu – Ti alloys.

precipitate is nearly the same for both yield and tensile strengths and this indicates that hardening is more due to particle shearing than Orowan looping. The higher strength values in ST +CW +PA condition as compared to those in as-quenched and PA state are the result of enhanced dislocation density due to CW as well as increased volume fraction of Cu4Ti, b l precipitate due to CW and ageing. It has been reported in our earlier work [7,8] that neither recovery nor recrystallisation occurs at the peak ageing temperature and time in prior cold deformed alloys and therefore, the effects of cold work persist in the yield and tensile strengths as well. Further, strengthening in ST+CW +PA Cu–4.5 and Cu–5.4 Ti alloys is also attributed to mechanical twins caused by the fine scale precipitation formed during quenching itself [20]. The mechanical properties of Cu–Ti alloys showing the strength contribution by different strengthening mechanisms are listed in Table 3. It is evident from Table 3 that different mechanisms, viz. solid solution strengthening, grain size strengthening (reported elsewhere [10]), hardening due to CW, and precipitation hardening are contributing to the strength of Cu–Ti alloys. Since grain size is kept constant in the present study, the contribution to strengthening by grain size is the same in all the alloys. In the as-quenched and PA condition, the Cu4Ti, b l precipitate contributes significantly to the strength of Cu–1.5 Ti (200%) and Cu – 2.7 Ti (140%) alloys. However, it is

small in Cu–4.5 Ti (59%) and Cu–5.4 Ti (34%) alloys, as the fine scale precipitation formed during quenching itself is already existing in these alloys prior to ageing. It is also clear from this table that the strength contribution due to the increased volume fraction of the Cu4Ti, b l precipitate in ST+ CW+ PA condition, though increases with Ti content, is nearly the same in Cu–4.5 and Cu–5.4 Ti alloys, as the increase in volume fraction of the precipitate itself is constant. Further, it is also less than the contribution made by the CW. It is therefore, concluded that cold work contributes considerable amount of strength (\80%) and the rest by Cu4Ti, b l precipitate in ST+ CW+ PA Cu–Ti alloys.

3.4. Correlation between strength and electrical conducti6ity Fig. 11 shows variation of 1.0% proof stress with electrical resistivity for different Cu base alloys for 1.0 at% solute. The resistivity and proof stress values for Cu base alloys except Cu–Ti, were taken from references [12] and [14], respectively. The increase in proof stress from Cu to Cu–Zn and Cu–Ni is marginal while it is significant for Cu–Be, Cu–Sn, and Cu–Ti alloys. This indicates that Ti is a potential strengthener as well as resistivity causing element in solid solution with copper, showing greater effect than other elements like Zn, Ni, Be, and Sn. This behaviour is attributed to the

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Table 3 Mechanical properties of Cu–Ti alloys showing the strength increase (Ds) due to different mechanisms, viz. solid solution strengthening, precipitation hardening, cold work, and increase in volume fraction of Cu4Ti, b l precipitate (DVfp) Alloy

Condition Vfp

sy (MPa)

su (MPa)

El (%)

DVfp

Dsy (MPa)

Ds py (MPa)

Ds C y (MPa)

Dsu (MPa)

Ds pu (MPa)

Ds C u (MPa)

Cu–1.5 Ti S SA SCA

— 0.06 0.09

112 350 670

292 520 760

41 23 9

— 0.06 0.03

42a 238 320

— — 40

— — 280

72a 228 240

— — 50

— — 190

Cu–2.7 Ti S SA SCA

— 0.14 0.20

192 460 950

430 680 1000

36 22 3.5

— 0.14 0.06

120a 268 490

— — 90

— — 400

210b 250 320

— — 60

— — 260

Cu–4.5 Ti S SA SCA

— 0.28 0.38

440 700 1280

680 890 1380

29 20 2

— 0.28 0.10

370b 260 580

— — 110

— — 470

460b 210 490

— — 110

— — 380

Cu–5.4 Ti S SA SCA

— 0.33 0.43

590 790 1400

780 930 1450

23 15 1.5

— 0.33 0.10

520b 200 610

— — 100

— — 510

560b 150 520

— — 150

— — 370

Yield strength of Cu: 70 MPa; tensile strength of Cu: 220 MPa. S, solution treated (ST); SA, ST+peak aged (PA) at 450oC; SCA, ST+cold worked by 90%+PA at 400oC. Dsy, increase in yield strength; Dsu, increase in ultimate tensile strength; Ds py , increase in yield strength due to increased volume fraction of the p Cu4Ti, b l precipitate after ST+cold work (90%)+peak ageing; Ds C y , increase in yield strength due to cold work (90%); Ds u, increase in tensile l C strength due to increased volume fraction of the Cu4Ti, b precipitate after ST+cold work (90%)+peak ageing; Ds u , increase in tensile strength due to cold work (90%). a Increase in strength due to solid solution strengthening. b Values include both solid solution strengthening and strengthening due to fine scale precipitation formed during quenching.

differences in size factor, elastic modulus, crystal structure, and valence of copper and the solutes. The tensile strength and electrical conductivity of Cu–Ti alloys compared with those of commercial Cu–Be alloys [21] are shown in Fig. 12. It is clear from this figure that while tensile strength is comparable, the electrical conductivity of Cu – Ti alloys is less than that of Cu–Be alloys. Though the electrical conductivity of Cu–Ti alloys is low, preliminary investigations carried out indicate that the Cu – 4.5 wt%Ti alloy exhibited the non-sparking property required for non-sparking tools satisfactorily. Cu – Ti alloys therefore, can be used as a substitute to Cu – Be alloys in those applications which require higher strength and lower conductivity. Further, in order to improve the electrical conductivity of Cu – Ti alloys significantly without foregoing the gains in strength, ternary additions such as boron may be required which is being considered in the future investigations.

tions and the effective interaction is more due to modulus mismatch than size misfit. The slope of the straight lines showing solid solution strengthening in copper by different solutes, increases in the order: Zn, Ni, Al, Si, Be, Sn, and Ti and is maximum for Ti (5.0).

4. Conclusions 1. Titanium is a potential solid solution strengthening element in copper resulting in higher strength than that of Zn, Ni, Al, Si, Be, and Sn. Solid solution strengthening in Cu – Ti alloys is attributed to the interaction of titanium atoms with screw disloca-

Fig. 11. Relation between 1.0% proof stress and electrical resistivity for 1.0 at% solute of different elements in solid solution with copper.

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alloy meets the requirements of non-sparking tools. Hence, Cu–Ti alloys can be used as a substitute for Cu–Be alloys in those applications which require high strength and low conductivity.

Acknowledgements The financial support of the Defence Research and Development Organisation is gratefully acknowledged. One of the authors (S. Nagarjuna) is grateful to Dr K.K. Sharma, Director-II, DMRL for encouragement and support.

References [1] [2] [3] [4]

Fig. 12. Relation between tensile strength and electrical conductivity of Cu – Ti alloys compared with Cu–Be alloys [21].

[5] [6] [7] [8]

2. In solution treated condition, a marked change is observed in hardness, strength, and elongation at about 4.0 wt%Ti beyond which strength increases sharply and elongation decreases further, with increasing Ti content. This is attributed to fine scale precipitation formed during quenching itself, in Cu– 4.5 and Cu–5.4 Ti alloys. 3. Contrary to the above, a linear variation (without any sharp changes) in hardness and tensile properties is found in PA condition with or without prior CW. This is attributed to the uniform precipitation of ordered, metastable, and coherent Cu4Ti, b l phase in all the four Cu – Ti alloys. 4. Yield and tensile strength increase significantly with volume fraction of Cu4Ti, b l precipitate. The Cu4Ti, b l precipitate in as-quenched and PA condition and the CW in ST+ CW +PA condition contribute significantly to the strength of Cu – Ti alloys. 5. Strength and electrical resistivity of copper are increased significantly by Ti additions. While tensile strength is comparable, the electrical conductivity of Cu–Ti alloys is less than that of commercial Cu–Be alloys in peak aged conditions. Still, Cu – 4.5 wt%Ti

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