A. The effect of isothermal and aging treatment on properties

Figures 1(a) and 1(b) show the variation of microhardness and electrical conductivity of Cu–10Ni–3Al–0.8Si alloy after isothermal treatment at different temperatures (Procedure-1). Both the microhardness and electrical conductivity increased with the isothermal time and decreased with the increase of isothermal temperature. The variation of microhardness and electrical conductivity of the isothermally treated alloys after being aged at 500 °C for 8 h (Procedure-2) is shown in Figs. 1(c) and 1(d). After aging, the sample isothermally treated at 650 °C had the lowest microhardness and highest electrical conductivity at every isothermal treatment time than those treated at other temperatures.

FIG. 1.

Variation of microhardness and electrical conductivity. (a) Variation of microhardness with Procedure-1; (b) variation of electrical conductivity with Procedure-1; (c) variation of microhardness with Procedure-2; and (d) variation of electrical conductivity with Procedure-2.

It was worthwhile to note that compared with the isothermally treated Cu–10Ni–3Al–0.8Si alloy, the microhardness and electrical conductivity of the alloy increased after isothermally treated and aged at 500 °C for 8 h, and the rise of microhardness and electrical conductivity also increased with the isothermally treated temperature (as shown in Table I), which indicated that the degree of super saturated solid solution increased with isothermal temperature.

TABLE I.

Influence of aging on microhardness and electrical conductivity.

It was generally believed that solid solute elements (Ni, Al, and Si) have strong electron scattering effect and secondary phases show weak electron scattering effect. Therefore, precipitation of secondary phases from the supersaturated solid solution led to the decrease of the amount of solid solute elements and increase of electrical conductivity. ^{Reference Blatt19,Reference Lei, Li, Xiao, Pang, Xiang, Qiu and Xiao20} When the alloy was isothermally treated at different temperatures for 900 s after solid solution treatment [Fig. 1(b)], the increment of electrical conductivity decreased with isothermal temperature, indicating that the amount of remained solute elements and the degree of supersaturated solid solution increased with isothermal temperature. As a result, the amount of precipitation increased with isothermal temperature after isothermal treatment and aging.

B. Evolution of microstructure during isothermal treatment and aging

1. Isothermal treatment at 550 °C

The microstructure and corresponding SADPs of Cu–10Ni–3Al–0.8Si alloy after solution treatment at 950 °C for 4 h and isothermal treatment at 550 °C for 30 s are shown in Fig. 2. Two kinds of nanoscale particles, one with zero-contrast line and the other without zero-contrast line, were found [Fig. 2(a)]. The corresponding SADP with the beam direction along [001]_Cu is shown in Fig. 2(b). Two kinds of precipitates were identified: one was L1₂ ordered Ni₃Al (FCC structure) with a lattice parameter of 0.358 nm and the other was DO₂₂ ordered δ-Ni₂Si (lattice parameter: a = 7.065 nm; b = 5.001 nm; and c = 3.731 nm). The HRTEM micrograph displayed in Figs. 2(c) and 2(d) showed that the particles with zero-contrast line were Ni₃Al and those without zero-contrast line were δ-Ni₂Si, respectively. The zero-contrast lines were perpendicular to g vector, indicating that Ni₃Al was coherent with the matrix. Precipitation free zone (PFZ) was observed around the large δ-Ni₂Si particles. When the alloy was isothermally treated at 550 °C, both Ni₂Si and Ni₃Al nucleated in the matrix and the Ni₂Si coarsen faster than Ni₃Al as isothermal time prolonged. The growth of Ni₂Si resulted in the decrease of the amount of solute nickel atoms in the matrix around the Ni₂Si particle, and the diffusion rate of solute elements was not fast enough at 550 °C. Therefore, the precipitation of Ni₃Al around large Ni₂Si particle was hindered, due to the lack of Ni atoms. The grain size of Ni₃Al particles was about 20 nm and that of Ni₂Si particles was about 50 nm [Fig. 2(a)]. The orientation relationships between the precipitates and matrix can be determined as ${[001]_{{\rm{Cu}}}}{\left\| {{{[001]}_{{\rm{N}}{{\rm{i}}_3}{\rm{Al}}}}\left\| {[001]} \right.} \right._\delta }$ , ${(110)_{{\rm{Cu}}}}{\left\| {{{(110)}_{{\rm{N}}{{\rm{i}}_3}{\rm{Al}}}}\left\| {(010)} \right.} \right._\delta }$ , and ${(1\bar 10)_{{\rm{Cu}}}}\left\| {{{(1\bar 10)}_{{\rm{N}}{{\rm{i}}_3}{\rm{Al}}}}} \right\|{(100)_\delta }$ .

FIG. 2.

Micrograph and SADPs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 550 °C for 30 s. (a) Bright field micrograph; (b) SADP of (a) with beam direction along [001]_Cu; (c) HRTEM of Ni₃Al; and (d) HRTEM of δ-Ni₂Si.

The microstructure of Cu–10Ni–3Al–0.8Si alloy as solution-treated at 950 °C for 4 h and isothermally treated at 550 °C for 900 s is displayed in Figs. 3(a) and 3(b) and the corresponding SADPs of Figs. 3(a) and 3(b) with beam direction along [001]_Cu and [112]_Cu are shown in Figs. 3(c) and 3(d), which indicated that both Ni₃Al and δ-Ni₂Si phases existed. The size of Ni₃Al particles with zero-contrast line was about 40 nm and remained coherent with the matrix, and that of large incoherent δ-Ni₂Si particles was about 100 nm [Fig. 3(a)]. However, DP occurred at the grain boundary after isothermally aged at 550 °C for 900 s [Fig. 3(b)]. Figures 3(e) and 3(f) show the bright field micrograph and the corresponding SADP of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 550 °C for 900 s and then aged at 500 °C for 8 h. The size of coherent Ni₃Al precipitates was still about 40 nm and δ-Ni₂Si with a size of about 50 nm precipitated in the matrix, however, the discontinuous δ-Ni₂Si precipitates with a size of about 200 nm appeared at the grain boundary after aging.

FIG. 3.

Micrographs and SADPs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 550 °C. (a) Bright field micrograph in the grain after isothermal treatment at 550 °C for 900 s; (b) bright field micrograph at the grain boundary after isothermal treatment at 550 °C for 900 s; (c) SADP of (a) with beam direction along [001]_Cu; (d) SADP of (b) with beam direction along [112]_Cu; (e) bright field micrograph after isothermal treatment at 550 °C for 900 s and aging at 500 °C for 8 h; and (f) SADP of (e) with beam direction along [112]_Cu.

2. Isothermal treatment at 650 °C

Figure 4 presents the bright field micrographs and corresponding SADPs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 650 °C for 30 s. The SADP indicated that both δ-Ni₂Si and Ni₃Al phases precipitated in the matrix [Fig. 4(b)]. The average size of δ-Ni₂Si particles with zero-contrast line was about 60 nm and a large quantity of coherent Ni₃Al particles with zero-contrast line was about 25 nm [Fig. 4(a)].

FIG. 4.

Micrograph and SADPs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 650 °C for 30 s. (a) Bright field micrograph and (b) SADP of (a) with beam direction along [001]_Cu.

Figure 5 shows the bright field micrographs, SADP, and HRTEMs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 650 °C for 900 s. Both Ni₃Al and δ-Ni₂Si phases precipitated in the matrix. The size of Ni₃Al particles with zero-contrast line was about 50 nm [Fig. 5(a)] and the HRTEM showed that those Ni₃Al particles remained coherent with the matrix [Fig. 5(d)]. The large δ-Ni₂Si precipitates with an average size of about 250 nm were observed in the matrix and the coarse δ-Ni₂Si precipitates with a size of about 400 nm precipitated at the grain boundary [Fig. 5(b)]. The HRTEM of the δ-Ni₂Si particle is shown in Fig. 5(e). After aging at 500 °C for 8 h, the size of discontinuous δ-Ni₂Si precipitates at the grain boundary was about 500 nm and that of coherent Ni₃Al particles was still about 50 nm; however, the fine δ-Ni₂Si particles precipitated during aging were about 70 nm [Fig. 5(f)].

FIG. 5.

Micrographs, SADPs, and HRTEMs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 650 °C. (a) Bright field micrograph in the grain after isothermal treatment at 650 °C for 900 s; (b) bright field micrograph at the grain boundary after isothermal treatment at 650 °C for 900 s; (c) SADP of (a) with beam direction along [112]_Cu; (d) HRTEM of Ni₃Al; (e) HRTEM of δ-Ni₂Si; and (f) bright field micrograph after isothermal treatment at 650 °C for 900 s and aging at 500 °C for 8 h.

3. Isothermal treatment at 750 °C

The bright field micrograph and corresponding SADPs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 750 °C for 30 s are shown in Fig. 6, which indicated that both δ-Ni₂Si and Ni₃Al with a size of about 15 nm precipitated in the matrix.

FIG. 6.

Micrograph and SADPs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 750 °C for 30 s.

Figure 7 presents the microstructure and corresponding SADPs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 750 °C for 900 s. The large δ-Ni₂Si with a size of about 400 nm and Ni₃Al with a size of about 100 nm precipitated [Fig. 7(a)]. DP began to nucleate at the grain boundary and cellular structure was observed after isothermally treated at 750 °C for 900 s [Fig. 7(b)]. The corresponding SADPs of Figs. 7(a) and 7(b) showed that both Ni₃Al and δ-Ni₂Si phases appeared in the matrix [Figs. 7(c) and 7(d)]. The diffraction of Ni₃Al in the SADPs was very weak, because the solid solution temperature for Ni₃Al was about 780 °C and precipitation of metastable Ni₃Al decreased obviously during isothermal treatment at 750 °C. ^{Reference Alexander21,Reference Cho, Kim and Lee22} However, after isothermal treatment at 750 °C for 900 s and aging at 500 °C for 8 h, the size of discontinuous δ-Ni₂Si precipitates at the grain boundary was about 150 nm, while a great number of fine δ-Ni₂Si and Ni₃Al particles precipitated during aging at 500 °C, and the size of those fine δ-Ni₂Si and coherent Ni₃Al particles inside the grain was about 40 nm [Figs. 7(e) and 7(f)].

FIG. 7.

Micrographs and SADPs of Cu–10Ni–3Al–0.8Si alloy isothermally treated at 750 °C. (a) Bright field micrograph of the grain after isothermal treatment at 750 °C for 900 s; (b) bright field micrograph at the grain boundary after isothermal treatment at 750 °C for 900 s; (c) SADP of (a) with beam direction along [001]_Cu; (d) SADP of (a) with beam direction along [112]_Cu; (e) bright field micrograph after isothermal treatment at 750 °C for 900 s and aging at 500 °C for 8 h; and (f) SADP of (e) with beam direction along [001]_Cu.

C. The quench sensitivity

The TTP curves of Cu–10Ni–3Al–0.8Si alloy based on microhardness experiments after isothermal and aging treatment are presented in Fig. 8. The TTP curves were fitted by the following equation ^{Reference Robinson, Cudd, Tanner and Dolan23} :

(1) $$C(T) = {k_1}{k_2}\exp \left[ {{{{k_3}\mathop {{k_4}}\nolimits^2 } \over {RT{{({k_4} - T)}^2}}}} \right]\exp \left[ {{{{k_5}} \over {RT}}} \right]\quad {\rm{,}}$$

where C(T) was the time, T was the isothermal temperature, R was the gas molar constant, and k ₁, k ₂, k ₃, k ₄, and k ₅ were constants, which were related to the natural logarithm of the volume fraction of untransformed precipitates, the reciprocal of the number of nucleation sites, the formation energy of nucleation, solid solution temperature, and the diffusion activation energy, respectively. The fitted coefficients of TTP curves are listed in Table II.

FIG. 8.

The TTP curves of Cu–10Ni–3Al–0.8Si alloy.

TABLE II.

The fitted coefficients of TTP curves.

The TTP curves were typical C-type ones and the fitted nose temperature was about 662 °C. The TTP curves could be divided into three regions: the nose temperature region, low temperature region, and high temperature region. The alloys presented high quench sensitivity as quenched from the nose temperature. The quench sensitivity of the alloy decreased continually when the quench temperature was above or below 650 °C.

Quench sensitivity was closely related to the decomposition ability of supersaturated solid solution during isothermal treatment. The precipitation in Cu–10Ni–3Al–0.8Si alloy during isothermal and aging treatment was a thermally activated process. The precipitation process was determined by both the diffusion rate of atoms and the nucleation rate of precipitates. The diffusion rate of atoms increased with the increasing isothermal temperature, which accelerated the growth rate of precipitates. However, the nucleation rate decreased with the increasing isothermal temperature, which leaded to the decrease of the amount of precipitates.

Both Ni₃Al and Ni₂Si precipitated during isothermal treatment and the degree of supersaturated solid solution increased with isothermal temperature. As a result, fine Ni₃Al and Ni₂Si particles would continue to precipitate during aging and the amount of precipitates increased with the isothermal temperature. The interactions between those precipitates and the matrix increased the strength of Cu–10Ni–3Al–0.8Si alloy, which can be described by the Orowan looping mechanism. The increment of strength increased with volume fraction of precipitates and decreased with particle size, which could be calculated as follows ^{Reference Torma, Kovács, Turmezey, Ungár and Kovács24} :

(2) $${\sigma _{0.2}} =\; 0.85m{G \over {2\pi {{(1 - \nu )}^{1/2}}}}{\left( {{f \over 2}} \right)^{1/2}} \times {{b\ln (R\pi /2{r_0})} \over {R\left[ {1 - (\pi /2){{(f/2)}^{1/2}}} \right]}}$$

where m is the Taylor factor of copper, G is the shear modulus of copper, R is the average radius of precipitates, b is the Burgers vector, ν is the Poisson's ratio, r ₀ is the cut-off radius of the dislocation core, and f is the volume fraction of precipitates.

The microhardness of Cu–10Ni–3Al–0.8Si alloy after solution treatment was about 135.6 HV, which increased to 249.9 HV after isothermal treatment at 550 °C for 900 s, then aged at 500 °C for 8 h; while the isothermal temperature increased to 750 °C, the hardness increased 257.4 HV after isothermal treatment and aging; when the alloy was isothermally treated at 650 °C, the hardness only increased to 226.3 HV [Fig. 1]. According to the bright field micrographs of Fig. 3(e), the particle size of fine Ni₃Al and Ni₂Si was about 40 and 50 nm, while that of discontinuous Ni₂Si was about 200 nm after isothermal treatment at 550 °C and aging. The increment of strength was calculated to be 304.8 MPa. As the isothermal temperature increased to 650 °C, the size of fine Ni₃Al and Ni₂Si was about 50 and 70 nm [Fig. 5(f)], and the size of discontinuous Ni₂Si was about 500 nm, thus the strength of Cu–10Ni–3Al–0.8Si alloy only increased by 248.5 MPa. Further increasing the isothermal temperature to 750 °C, the size of fine Ni₃Al and Ni₂Si was about 20 nm [Fig. 7(e)], while the size of discontinuous Ni₂Si was about 150 nm, as a result, the strength increased by 337.5 MPa. The calculation results indicated that Cu–10Ni–3Al–0.8Si alloy showed lowest strengthening response after isothermally treated at 650 °C for 900 s and then aged at 500 °C for 8 h, which agreed well with our experimental results. The Cu–10Ni–3Al–0.8Si alloy showed lowest strength response after isothermal treatment at the nose temperature region and aging, because the incubation of precipitation was shorter and the nucleation and growth rate of those precipitates were faster. In addition, a plenty of discontinuous δ-Ni₂Si phases precipitated at the grain boundary after isothermal treatment at the nose temperature and grew up quickly after aging, resulting in low microhardness value. Therefore, the Cu–10Ni–3Al–0.8Si alloy presented high quench sensitivity at the nose temperature region.