Biaxial deformation behavior and formability of precipitation hardened ultra-fine grained (UFG) Cu–Cr–Zr alloy

Author’s Accepted Manuscript Biaxial deformation behavior and formability of precipitation hardened ultra-fine grained (UFG) cucr-zr alloy Onur Saray www.elsevier.com/locate/msea

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S0921-5093(16)30020-X http://dx.doi.org/10.1016/j.msea.2016.01.021 MSA33208

To appear in: Materials Science & Engineering A Received date: 22 October 2015 Revised date: 6 January 2016 Accepted date: 6 January 2016 Cite this article as: Onur Saray, Biaxial deformation behavior and formability of precipitation hardened ultra-fine grained (UFG) cu-cr-zr alloy, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2016.01.021 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Biaxial deformation behavior and formability of precipitation hardened ultra-fine grained (UFG) Cu-Cr-Zr alloy Onur Saray Bursa Technical University, Department of Mechanical Engineering, 16190 Bursa, Turkey

Abstract The combined effects of ultrafine-grained (UFG) microstructure and precipitation on the formability and biaxial deformation behavior of a Cu-Cr-Zr alloy were investigated. The UFG microstructure formation results in good formability with an Erichsen index (Ei) of 4.05 mm compared to that of peak-aged coarse grained (CG) alloy (3.95 mm). Aging heat treatments increase strength and formability of the UFG alloy simultaneously. Biaxial deformation behavior is found to be dependent on the strain hardenability. Excellent strain hardenability of the CG alloy brought about higher punch displacement within the membrane stretching regime. However, deformation localization with the early onset of necking is evident in the UFG alloy. Subsequent aging treatments decrease deformation localization behavior of UFG alloy with increasing aging durations. Results also show that both grain refinement and aging increased the punch load due to enhanced strength. A linear relationship is generated based on punch load vs. punch displacement curve slope to predict ultimate tensile strength (UTS) with high accuracy. It is concluded that synergetic effect of UFG microstructure formation and subsequent aging provides a simple and effective procedure to produce Cu-Cr-Zr alloy for applications where balance of strength and formability are needed. Keywords: Ultra-fine grained materials; formability; aging; Erichsen test; biaxial deformation; equal channel angular extrusion *Corresponding

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1. Introduction Binary Cu-Cr and ternary Cu-Cr-Zr alloys are generally utilized in critical functional applications including railway contact elements, spot welding electrodes and thermonuclear reactor parts due to their high strength, good fatigue resistance, excellent electrical and thermal conductivity, good thermal stability at high temperatures, outstanding resistance to corrosion and ease of fabrication [1-6]. Aging heat treatment as a primary strengthening method is employed to enhance the strength of these alloys with an adequate electrical conductivity [1, 4, 7-9]. Generally, tensile strength values in the range of 450-500 MPa accompanied with an electrical conductivity of around 75% IACS can be achieved after peak-aging heat treatment [5, 9, 10]. Main aspects of the precipitation hardening characteristics and their effect on the strength, ductility and electrical conductivity of Cu-Cr-Zr alloys have been widely investigated [1, 7, 1114] . In previous studies, however, it has been shown that the strengthening by conventional aging is limited due to the low solubility of Cr and Zr in the coarse grained (CG) Cu matrix [5, 15]. Some efforts have been undertaken using plastic deformation with classical metal forming applications for further enhancement of the mechanical properties of the Cu-Cr-(Zr) alloys [1618]. However, improvements in strength and hardness have been found limited after application of these methods [5, 16-18]. Grain refinement down to sub-micron levels using severe plastic deformation (SPD) methods has been considered as a good strategy to achieve extraordinary strength enhancements [19, 20]. Application of the SPD to Cu-Cr-(Zr) alloys has been resulted in outstanding improvement of the mechanical properties. Many reports points out that synergetic effect of grain refinement and extensive dislocation accumulation accompanied with the formation of the nanoscaled precipitates cause these alloys to reach exceptional levels of mechanical strength [5-7, 10, 21-26]. Also it has been shown that, these strength enhancements can be achieved with an 2

acceptable electrical conductivity by applying appropriate post-SPD aging heat treatments. [7, 10, 23, 26-29]. Moreover, homogeneously dispersed nano-scaled precipitates introduced by aging of ultra-fine grained (UFG) Cu-Cr-Zr alloys have been found effective to prevent from early onset of plastic instability [6, 25]. Such a combination of the high strength, improved plastic stability and reasonable electrical conductivity obtained by the combined effect of SPD and aging may lead to consider UFG Cu-Cr-Zr alloys as the materials of choice for fabricating micro-electromechanical (MEMs) components with micro-metal forming applications [27, 28, 30-32]. Because, in order to prevent from the size-effects and to ensure reliable property control, grain size of the material to be used in micro-forming of the MEMs must be smaller than the smallest dimension of the components [28, 32]. Moreover, improved plastic stability may be expected to be rewarding to form crack or defect free components with complicated geometrical features [28, 31]. Also, high strength of the UFG Cu-Cr-Zr alloy may meet expectations to reduce component dimensions, weight and/or energy consumption without a load bearing loss [28, 31]. However, to the author’s best knowledge, no study available in the literature addresses the formability and biaxial deformation behavior of aged UFG Cu-Cr and/or Cu-Cr-Zr alloys. Therefore, the present work has investigated the effect of aging on the formability and biaxial deformation behavior of a Cu-Cr-Zr alloy before and after the formation of UFG via equal channel angular extrusion (ECAE). Effects of the formation of UFG microstructure and subsequent aging heat treatment conditions on the microstructure and mechanical properties were studied and correlated with the biaxial deformation mechanisms and formability. 2. Experimental procedure A Cu-Cr-Zr alloy with a composition of Cu-0.80Cr-0.080Zr (wt. %) was used in present study. As-received samples were solution treated at 1020 ºC for 20 min and then quenched into

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water in order to obtain a supersaturated copper matrix. Aging treatments were carried out at 475 ºC for 5 h in a furnace and subsequently quenched in water. The aging temperature and aging time for the supersaturated alloy were chosen

to reach the high hardness along with adequate

electrical conductivity based on the results presented in [5, 33]. Hardness and electrical conductivity of CG alloy after this aging treatment (475 ºC for 5 h) were determined with preliminary tests as165 Hv0.5 and 76 %IACS, respectively. Billets machined from supersaturated alloy in dimensions of 20 mm x 20 mm x 120 mm were subjected to ECAE at room temperature in a die with a sharp 90 channel cross-section angle at a rate of 2.5 mm s-1. ECAE processes were performed up to eight passes using route-BC (8Bc) in which the billets are rotated 90 in the same direction around their longitudinal axis between successive passes. ECAE processed billets were then subjected to age hardening heat treatment at 450 for 30 min, 90 min and 240 min. These aging conditions were selected based on the results reported in [29], where aging characteristics of the UFG Cu-Cr-Zr alloy by means of variations of hardness and electrical conductivity were determined. In the current study, repeatability of these results was also checked with preliminary tests, and heat treatment conditions were determined to cover various states of aged UFG microstructure. According to results of preliminary tests, aging at 450°C for 30 min leads to peak aging condition with hardness and electrical conductivity of 229 Hv0.5 and 36 IACS%, respectively. Aging at 450°C for 240 min yields the condition where a reasonable electrical conductivity for industrial applications of 74 IACS% obtained with a hardness of 185Hv0.5. Aging at 450°C for 90 min stands for an intermediate state of hardness (215 Hv0.5) and electrical conductivity (71 IACS%) [29]. By this way, it is aimed to observe

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variations in biaxial deformation behavior and formability of the UFG Cu-Cr-Zr alloys in conditions that may be relevant to industrial applications of the UFG alloy. Microstructure of the ECAE processed UFG samples were characterized with transmission electron microscope (TEM) operated at 200 kV. For the TEM investigations, 0.5 mm thick discs were sectioned from the transverse plane of the extruded billets, and then mechanically ground and polished down to 0.15 mm thick foils (Fig. 1(a)). Large electrontransparent areas were obtained in these foils by conventional twin-jet polishing utilizing a 5 % perchloric acid solution under an applied potential of 25 V at -40 C. Tensile tests were conducted using dog-bone shaped tension samples with dimensions of 1.5 mm x 3 mm x 26 mm machined from the ECAE-processed billets with their tensile axis aligned with the direction of extrusion (Fig. 1(a)). The tests were performed using an Instron3382 electro-mechanical load frame at a strain rate of 5.4 x 10-4 s-1. Strain was measured using a video-type extensometer. For each case, three experiments were conducted on companion specimens to check the repeatability of the results. In addition, hardness measurements were performed on transverse plane using a Vickers micro-hardness tester with a load of 500 g for the dwell time 10 s (Fig. 1(a)). Formability tests were performed on the samples with the dimensions of 20 mm x 20 mm x 0.9±0.01 mm using a miniaturized Erichsen test die (Fig. 1(b)). Erichsen samples sectioned from the transverse plane of the extruded billets using the wire-EDM technique (Fig. 1(a)). After sectioning, surfaces of the samples were ground using emery paper down to a grid size of 1000 and then finished with 1µm alumina solution before Erichsen tests in order to prevent form surface defects that can act as a crack initiation points. The tests were conducted without lubrication with a punch speed of 0.01 mm s-1. Load–displacement data were collected during the

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tests, and Erichsen index (Ei) and the load corresponding to this index (FEi) were evaluated from the load–displacement data. Morphological features of the Erichsen tested samples were examined using a ZEISS-Evo50 scanning electron microscope (SEM) operated in the secondary electron mode at 15 keV.

3. Results 3.1. Microstructure Solution treatment and quenching of the Cu-Cr-Zr alloy resulted in formation of a microstructure consisted of coarse copper grains with a mean grain size of 52 m. ECAE processing of the quenched alloy transformed the CG microstructure into a UFG microstructure (Fig. 2(a)). Grain morphology of the UFG structure are mainly equiaxed and more or less homogeneous through the microstructure. This is a common feature of FCC and BCC materials subjected to ECAE fallowing the route-Bc [28, 34-37]. It is observed from Fig. 2(a) that dislocations mainly accumulate at the vicinity of the grain boundaries while dislocation density of grain interiors is lower compared to grain boundaries. Grain size of the UFG microstructure was determined to be 180±20 nm using linear intercept method. It has been well known that UFG materials fabricated with SPD methods contain both high angle grain boundaries (HAGBs) and low angle grain boundaries (LAGBs) [38]. The mean grain size reported here is determined by accounting both types of grain boundaries. Aging treatment of the UFG Cu-Cr-Zr alloy affects the UFG microstructure. Dislocation density somehow decreased and some dislocation free grains formed according to the TEM observations (Fig. 2(b)-(g)). Also, grain boundaries became sharper by increasing aging duration (Fig. 2 (b)-(g)). Peak aging heat treatment at 450°C for 30 min somehow coarsened the grain size

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of the as-ECAE processed microstructure to 643±68 nm (Fig. 2(b)). Moreover, uniformly distributed very fine particles precipitated

at the grain interiors after the heat treatment (Fig.

2(c)). Increasing aging duration somehow increased the grain size to 765±110 nm and 980±250 nm after aging treatments for 90 min (Fig. 2(d)-(e)) and 240 min (Fig. 2(f)-(g)), respectively. Hence, it can be said that grain sizes are still in the range of UFG structures after these aging treatments. However, increasing aging time also resulted in over-aging of the UFG microstructure. It is evident from Fig. 2(e) and Fig. 2(g) that particles were coarsened with increasing aging duration to 90 min and 240 min. Also, coarse precipitates more likely formed at the sites close to the grain boundaries (Fig. 2(f)). This variation became more pronounced after the aging of the UFG alloy for 240 min Fig. 2(g).

3.3 Uniaxial deformation behavior and mechanical properties Engineering stress- strain curves of CG and UFG Cu-Cr-Zr alloy were represented in Fig. 3. Mechanical properties (yield strength y, ultimate tensile strength UTS, uniform elongation u, and elongation to failure f) were summarized in Table 1. It is seen that CG alloy features a strain hardening dominated uniaxial deformation behavior in both quenched and aging treated conditions (Fig. 3). Subsequent aging treatment enhances low strength of the quenched alloy with a decline of ductility. Yield strength and UTS of the quenched alloy increase from 218 MPa and 274 MPa to 295 MPa and 380 MPa, respectively. However, uniform elongation and elongation to failure decrease from 22% and 36% to 9% and 16%, respectively, (Table 1). Formation of the UFG microstructure strongly affects the uniaxial deformation behavior. Mainly, as-ECAE processed alloy reaches to UTS without a considerable elongation of %2 after yielding (Fig. 3). Hence, localized deformation with early onset of macroscopic necking

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dominates deformation behavior of the UFG alloy. Formation of the UFG microstructure enhances strength of the alloy. Yield strength and UTS of the UFG alloy were determined to be 516 MPa and 572 MPa, respectively, which are higher than those obtained by aging treatment of CG alloy. Also elongation to failure of the UFG alloy is 15% which is close to that of aging treated CG alloy (16 %). Peak-aging of the UFG alloy further increases the strength. Yield strength and UTS of this condition are 618 MPa and 627 MPa, respectively. These values are roughly 2 times higher than those of the quenched CG alloy. Extending the aging time slightly decreases strength of the UFG alloy. The matter becomes more pronounced with increasing aging durations. After aging for 240 min, yield strength and UTS of the UFG alloy are 557 MPa and 601 MPa, respectively. It is important to note that, these values are still higher than those of the as ECAE processed condition (Table 1). Also, aging heat treatment duration is found to be inoperative on uniaxial deformation behavior of the UFG structure (Fig. 3). So that, uniform elongations of the aging treated UFG alloy are in the range of the 0.9%-2% (Table 1). This may indicate that, localized deformation with early onset of necking is evident for the aging treated UFG alloy regardless to aging conditions (Table 1, Fig. 3). Peak-aging decreases the elongation to failure of the UFG alloy down to 9.5%. However, extending aging treatment durations considerably increases this value. For instance, aging treatment for 240 min causes UFG alloy to reflect an elongation to failure of 17%, which is slightly higher than that of the peak-aged CG alloy (Table 1).

3.4 Biaxial deformation behavior and formability Load-displacement curves from Erichsen tests are shown in Fig. 4(a). Erichsen index (Ei) and punch load at Erichsen indexes (FEi) are summarized in Table. 2. Formerly, deformation stages of a sample subjected to Erichsen test has been outlined in several publications [27, 30, 398

41]. These deformation stages comprise elastic bending, transition to plastic bending with microyielding, yield surface propagation, membrane stretching, deformation localization and fracture [27, 30, 39-41]. It has been shown that transition between these deformation stages changes F-X curve slope [27, 30]. In order to distinguish these deformation stages from each other, first order derivatives of punch load (F) with respect to the displacement (X) were calculated and represented in Fig. 4(b). Among the deformation stages of Erichsen test, membrane stretching regime is the most distinguishing one determining the formability and/or Erichsen index of the sample. In this stage, thickness of sample decreases under the effect of biaxial tension stresses acting on the dome wall [27, 30, 39-41]. Hence, plastic flow characteristics of the tested materials have been found to be strongly effective on membrane stretching regime [27, 30, 42]. In this point of view, greater emphasis was given to that regime in order to understand the effect of the grain refinement and/or aging treatments on the formability and biaxial deformation behavior of the alloy. Displacements occurring within the membrane stretching regime where the curve slope remains nearly constant and/or reflected a slight increase are presented on Fig. 4(b) with solid symbols. As a common feature, load-displacement curves of Cu-Cr-Zr samples in various processing conditions reflect typical deformation stages of the Erichsen test (Fig. 4). However, applied processing conditions are found to be effective on the membrane stretching regime of the biaxial deformation (Fig. 4(b)). Excellent formability by means of the highest EI of 4.68 mm was detected for the quenched CG alloy (Table. 2). Displacement within the membrane stretching regime of quenched alloy was determined to be 2.4 mm which is about 50% of the EI (Fig. 4(b)). Subsequent aging treatment slightly decreased the EI of CG alloy to 3.95 mm. However, displacement within the membrane stretching regime comprised 55% of the Ei. Thus, it is important to note that contribution of the membrane stretching regime on the biaxial deformation 9

slightly increases after aging treatment. Formation of the UFG microstructure strongly affects the deformation behavior of the Cu-Cr-Zr alloy. The EI of UFG alloy is 4.05 mm which is slightly higher than that of the aging treated CG alloy. However, the membrane stretching regime is considerably contracted to 0.3 mm, which is about 7% of the EI. Aging treatments applied for various durations have minor effects on the Ei values. Aging for 30 min yielded an Ei of 4.15 mm which is very close to that of the UFG alloy. Extending the aging time to 90 min and 240 min increased the EI to 4.20 mm and 4.23 mm, respectively. Also, displacements within the membrane stretching regime take values of 0.2 mm, 0.3 mm and 0.8 mm after aging treatments for 30 min, 90 min and 240 min, respectively (Fig. 4(b)). CG alloy exhibits lowest value of FEI of 3535 N among all applied processing conditions (Table 2). This value considerably increased after both aging treatment and UFG microstructure formation to about 5391 N and to 6955 N, respectively. Aging treatment of UFG alloy further increased the FEI value to 8417 N (Table 2). But, longer aging durations of 90 min and 240 min brought about slightly lower FEI values of 8291 N and 7604 N, respectively (Table 2). It may be important to emphasize that these FEI values are still higher than those of the UFG alloy. The results indicate an apparent similarity in trends of variation of FEi and strength with various processing conditions. Deformation mode and homogeneity of the Erichsen tested CG and UFG samples were studied by inspection of the morphological features of the dome free surfaces (Fig. 5 and Fig. 6). Biaxial tension stresses acting on the CG Erichsen samples cause roughening of the dome free surface after quenched and aging treated conditions (Fig. 5(a)-(d)). Mirror finished pre-tests surface appearance transforms into a roughened and grainy appearance by emerging of grains and grain boundaries of the CG structures after Erichsen tests (Fig. 5(a)-(d)). At the grain interiors, continuous deformation bands are evident (Fig. 5(b)). Variation in orientation of these bands by 10

passing the grain boundaries indicates that dislocation slip is the main mechanism of deformation (Fig.5 (b), (d)). These morphological transformation on dome free-surface so-called “orange-peel effect” is mainly attributed to surface level differences between the neighboring grains [43]. Orange peel effect mainly occurs due to crystallographic misorientations leading to variation in deformation characteristics [27, 30, 40, 43]. Fracture of the quenched and aged CG samples occurs with necking (Fig. 5(a)-(d)). The cracks are mainly formed at the grain interiors of the quenched alloy (Fig. 5(b)). However, cracks are also visible through grain boundaries of the aged alloy (Fig. 5(d)). Dome free surface appearance of the Erichsen tested UFG samples does not reflect a considerable roughening compared to that of the CG samples (Fig. 6(a)-(b)). Simply, roughening with orange peel effect was mostly eliminated by the formation of the UFG microstructure as a result of grain refinement [27, 30] (Fig. 6(a)). Instead, shallow deformation bands locally form at the necking region (Fig. 6(b)). Aging of the UFG samples mainly reflected very similar feature with that of the as processed sample (Fig. 6(c)-(h)). Generally, dome free surfaces of the aged UFG alloys reflect a necking region where deformation bands were formed (Fig. 6(d), (f) and (h)). Qualitatively, increasing the aging temperature leads to enlargement of the necking region (Fig. 6(d), (f) and (h)). In other words, localized necking behavior occurred after aging for 30 min (Fig. 6(d)) transforms into a diffuse necking behavior with increasing aging temperature (Fig. 6(h)).

4.

Discussion

4.1. Biaxial deformation behavior Membrane stretching regime of the Erichsen test comprises uniform reduction of the sample thickness with the aid of the biaxial tension stresses [27, 30, 42, 44]. One can expect a reduction in load-bearing capacity of the sample by decreasing dome wall thickness [27, 30]. 11

However, dF/dX values remain nearly stable with the punch displacement (Fig. 4(b)). This may indicate that reduction in load baring capacity is balanced with the

strain hardening while the

dome thickness is decreasing [27, 30, 45]. Hence, membrane stretching regime is directly related to the strain hardenability and consequently dislocation interaction during plastic deformation of the sample [27, 30, 40]. Membrane stretching regime of the CG sample constitutes a great portion (nearly 50%) of the Ei (Fig.4 (b)). After aging treatment, contribution of the membrane stretching to the Ei of the CG alloy is very close to that of the quenched condition (Fig. 4(b)). This may be attributed to the excellent strain hardenability of the CG samples before and after aging heat treatment (Fig. 3, Table 1). As can be understood from the uniaxial tension test results, uniform elongation of these samples is nearly half of the elongation to failure (Fig. 3, Table 1). Catastrophic decrease in the strain hardenability of the alloy after formation of the UFG microstructure may cause the occurrence of a very limited displacement within the membrane stretching regime of the Erichsen test (Fig. 4(b)). The same situation is also evident after aging treatment of the UFG alloy. Fig. 4(b) clearly indicates that displacement in the membrane stretching regime of the UFG sample slightly decreases after at the 450 C for 30 min. Similarly, extending aging duration caused a slight increase in the displacement within this regime (Fig. 4(a)-(b)) as a reflection of scarce enhancement in uniform elongations (Table 1). Saray et al. [27] indicated previously that the ratio of the uniform elongation to elongation to failure follows similar trend with the ratio of displacement in membrane stretching regime to Ei for the ECAE processed UFG IF-steel [27]. Current results obtained with CG and UFG Cu-Cr-Zr alloy is also in consistent with this finding [27]. Limited displacement within membrane stretching basically indicates that load bearing capability of the UFG samples decreases by the early onset of necking. Localized deformation is a common trend in UFG pure metals and alloys produced with SPD

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methods [5, 24, 46-48]. This characteristic behavior of the UFG material is attributed to the loss of strain hardenability after SPD [5, 24, 46-48]. The microstructure of UFG materials mainly consists of refined grains surrounded by strain-induced dislocation boundaries (SIB) [24, 46-48]. Therefore, the free path of dislocations is lower than that of CG microstructure (Fig. 2). Also, these boundaries act as a trap for the mobile dislocations [49]. Thus, the formation of UFG microstructure decreases the possibility of both dislocation–dislocation and dislocation–boundary interactions, which in turn limited strain hardenability and deformation localization [48-50]. Dome free surface appearances of the Erichsen tested UFG samples also support the occurrence of the deformation localization. Fig. 7(a)-(h) indicate that, dome free surface of the Erichsen tested UFG samples can be characterized by the formation of deformation bands at necking regions while rest of the sample reflects minor variations from the initially polished surface. It can be qualitatively observed that these areas are narrower and/or sharper in conditions with limited uniform elongation as in the case of peak-aged UFG sample (450C for 30 min) (Fig. 6(d)). This finding is also in consistence with the former studies on the biaxial deformation behavior of the UFG metals like IF-steel [27, 30] and pure copper [32].

4.2. Formability Decrease in the Ei of CG Cu-Cr-Zr alloy after aging treatment can be attributed to the enhanced pining effect of fine coherent precipitates leading to formation of dislocation pile-ups. These pile-ups may act as a necking initiation points leading to decrease in uniaxial and biaxial ductility. This idea may be supported by the fracture behavior of the CG alloys (Fig. 5(a)-(d)). Fracture of the aged CG sample occurs through grain boundaries (Fig. 5(d)). This may be due to precipitations at the grain boundary sites [15, 24]. Particles precipitated along grain boundaries 13

may weaken the grain boundaries due to the more effective dislocation accumulation at the vicinity of these sites. Eventually, they may act as the easy crack initiation regions after aging treatment. Decrease in Ei after the formation of UFG structure may be related to early onset of necking leading to localized deformation as explained above. The occurrence of deformation at the local areas may impede the contribution of the sample to Ei and formability. In turn, fracture of UFG sample occurs with lower Ei values. Contrary to the CG alloy, aging of the UFG structure to peak hardness (at 450°C for 30 min) does not considerably affect in the Ei. Also, increasing in aging duration enhances the Ei values (Fig. 4, Table 2). Generally, two important microstructural mechanisms may play role simultaneously during aging treatment of the UFG structure. The first one is the formation of the nano-sized precipitates (Fig. 2(c), (e) and 2(g)) and the second one is recovery and recrystallization of the UFG structure as a result of the thermal energy input during aging treatments (Fig. 2(b), (d) and (f)). Precipitates may be considered as a stress concentration points leading to increase in cracking tendency of the structure. On the other hand, recovery and/or recrystallization reduce the cracking tendency of the structure due to decrease in internal energy of the UFG microstructure. Hence, a balance between these two mechanisms makes Ei nearly stable after peak-aging (Table 2). This balance may be broken by extending the aging treatment durations. As can be seen from the TEM micrographs (Fig. 2), a considerable increase in grain size and decrease in the dislocation density are evident by prolonging aging treatment duration (Fig. 2(b), (d) and (f)). This basically indicates more effective occurrence of softening mechanisms (recovery and recrystallization), which in turns a slight increase in the Ei (Fig. 4, Table 2). Also, from TEM micrographs, increasing aging duration to 90 min and 240 min leads to an apparent coarsening of the precipitates (Fig. 2(e), Fig. 2(g)). This may contribute to increase in the Ei by decreasing the dislocation pining effects of the precipitates. 14

Variation of strength and FEi values reflects the same trends with respect to applied thermomechanical treatments as an expected result. It is well established that aging of the CG Cu-Cr-Zr alloy enhances the strength of the alloy due to the precipitation of the coherent Cu-rich particles in the microstructure [1, 7-9, 15]. The strengthening mechanism is mainly based on the resistance of precipitates to dislocation motion and/or slip mechanism due to the interaction of the dislocations with the stress fields around coherent particles [1, 7-9, 15]. Thus, the FEi value of solution treated alloy considerably increases in accordance to the strengthening of the CG alloy after aging treatment (Fig. 3 and Fig. 4). On the other hand, increases in strength and FEi values after ECAE processes are mainly related to the formation of UFG microstructure (Fig. 2(a)). It is well accepted that Hall-Petch strengthening via grain refinement down to sub-micron size range and dislocation accumulation leading to work hardening are two major microstructural evolutions which contribute to reach exceptional levels of strength in UFG materials [28, 31, 42-45]. Peakaging of the UFG alloy (at 450C for 30 min) forms coherent Cu-rich precipitates in optimum size and distribution (Fig. 2(b)-(c)). This further enhances the strength and consequently FEi values via incorporation of precipitation hardening (Table 1 and Table 2). However, extending the aging time makes softening mechanisms more effective. This can be understood from the microstructure obtained after aging of UFG structure for 90 min (Fig. 2(c)) and 240 min (Fig. 2(d)), where coarsening of the grains and decrease of dislocation density can be clearly seen in UFG microstructures. Decrease in yield strength and UTS also supports this idea (Table 1). On this point of view, it may be concluded that the reason of decreasing the FEi after aging treatment for 90 min and 240 min may be related to thermal input leading to the occurrence recovery and/or recrystallization mechanisms. Another important reason to decrease in the FEi of UFG alloy with increasing aging time may be occurrence of over-aging (Fig. 2(e), Fig. 2(g)). It is well known that

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the uniform size and distribution of the precipitations are two important parameters being effective on the resulting strength of the precipitation hardened materials. Extending aging duration coarsens the precipitation size and increase inter-particle distance (Fig. 2(e), Fig. 2(g)). Because of these two effects, stress necessary for dislocations to loop around a precipitate decreases. In result, resistance to plastic flow and consequently the FEi and overall strength decrease.

4.3. Estimation of UTS using Miniaturized Erichsen test data: Miniaturized Erichsen tests can also be utilized to estimate the uniaxial tensile strength of materials. This estimation has been found necessary and/or useful for applications where running a standard tensile test is impractical due to the inadequate sample dimensions like irritated materials used in nuclear power plants. Up to now, several correlations have been proposed in order to estimate UTS by using the data obtained from small punch testing. In these correlations, the ratio of the maximum punch load (Pmax) to initial sample thickness (t) or squared initial sample thickness (t2) was used to estimate the UTS [39, 45, 51-53]. It is well known that the necking of the sample occurs if applied stress level reaches to UTS in uniaxial tension tests. In miniaturized Erichsen tests (or small punch tests), on the other hand, necking of the sample occurs before the applied punch load reaches to maximum punch load. Hence, correlations based on the maximum punch load may be contradictory since the point of maximum load does not represent the necking situation as in the case of UTS in uniaxial tension tests [45]. According to findings of the current study, sample attains to necking by the end of the membrane stretching regime (Fig. 4). Because, the slope of the F-X curve remains nearly constant if the strain hardening can compensates the load bearing loss due to the thickness reduction within membrane stretching regime (Fig. 4(b)). If the necking occurs, then strain 16

hardening cannot compensate the loss in load bearing capacity due to the locally thinned regions. Hence, slope of the F-X curve sharply decreases by the neck formation (Fig. 4(b)). On this point of view, the sensitivity of the curve slope within the membrane stretching regime to the necking may be correlated with uniform elongation stage of the uniaxial tension test that ends by reaching to UTS. On the other hand, it was reported that the F-X data is most sensitive to sample thickness[54]. As can be predicted, increasing the sample thickness leads to increase in slope of F-X curve within the stretching regime. Because, the same increase in punch displacement can be achieved with a higher force increment during the testing of thicker samples and vice versa. This may require normalization of the curve slope to make it independent to sample thickness. On this basis, variation of the UTS with the maximum value of normalized F-X curve slope within the membrane stretching regime is plotted as shown in Fig. 6. This figure clearly indicates the existence of a linear relationship that can be represented as follows: ( ) Where,

and

(1) are constants, ( ) is the maximum slope of the curve within the membrane

stretching regime and t is the initial thickness of the sample. The term

is used to normalize the

term “( ) ” with respect to sample thickness. Proposed Expression 1 correlates the UTS with a high coefficient of determination (R2) of 0.962. Also, constants α1 and α2 of the Expression 1 is determined to be 0.242 and 0, respectively, for the UFG Cu-Cr-Zr alloy (Fig. 7). In order to exhibit the material dependency of the proposed correlation, results of Ref [27] is also adapted to the Expression 1 in Fig. 7. Adapted data was generated by miniaturized Erichsen testing of subsequently annealed UFG IF-steel using the same testing die, specimen dimensions and loading conditions [27]. It is clear from Fig. 7 that the proposed expression can

17

correlate UTS of UFG IF-steels with a reasonable coefficient of determination (R2) of 0.941. Constants α1 and α2 of the Expression 1 is determined to be 0.3953 and 338.22, respectively, for the UFG IF-steel (Fig. 7) [27]. This may indicate that the proposed Expression 1 is capable to predict the UTS of various UFG-materials with different material dependent variables like crystallographic structure, strengthening mechanism and strain hardening capabilities based on the data obtained from the miniaturized Erichsen tests.

5. Conclusions In this study, synergetic effects of ultra-fine grained (UFG) microstructure formation and precipitation hardening on the biaxial deformation behavior and formability of Cu-Cr-Zr alloy are investigated. The main findings and conclusion of the study can be outlined as follows: 1. ECAE process significantly refined the microstructure of the Cu-Cr-Zr alloy and formed an UFG microstructure with a mean grain size of 180 nm. Aging treatments applied at 450°C for 30 min formed very fine and homogeneously distributed precipitates. Extending the aging heat treatment duration caused over-aging of the microstructure leading to increase in precipitate size. Recovery and recrystallization mechanisms simultaneously took place during the aging heat treatments resulted with a decrease in dislocation density and an increase in grain size of UFG structure. 2. Formation of the UFG structure gradually enhanced strength of the alloy with the expenses of ductility and strain hardenability. Strength enhancement obtained by UFG microstructure formation was more effective than that obtained with peak-aging heat treatment of the CG alloy. Subsequent aging treatments applied to UFG alloy further increased the strength. However uniaxial tension ductility of the UFG alloy was slightly affected by the aging heat treatment conditions. 18

3. Biaxial deformation mechanism of the Cu-Cr-Zr alloy was strongly affected by UFG microstructure formation. The CG alloy deformed by exhibiting a large membrane stretching region in Erichsen tests with aid of strain hardening. On the other hand, contribution of the strain hardening to the deformation was limited in UFG structure. Thus, localized deformation behavior by formation of contracted membrane stretching region was evident. Subsequent aging heat treatments enlarged the membrane stretching regime and enhanced contribution of the strain hardening to the biaxial deformation. The UFG microstructure formation resulted in a good formability with an Erichsen index (Ei) of 4.05 mm, which is slightly higher than that of peak-aged coarse grained (CG) alloy (3.95 mm). Also, subsequent peak-aging treatment (at 450°C for 30 min) of UFG alloy considerably enhanced the Ei to 4.15 mm which is higher than that of the peak-aging treated CG alloy. The EI of the UFG alloy further increased by longer aging heat treatments, and it reached to 4.23 mm after aging at 450°C for 240 min. Also, 1.9 and 1.6 times higher yield strength and UTS than those of the peak-aged CG alloy accompanied with reasonable electrical conductivity of 74 %IACS can be achieved after aging treatment of UFG alloy at this condition. Hence, a good combination of higher strength, better formability with a resonable electrical conductivity was achieved after aging heat treatment of the UFG alloy at 450°C for 240 min. 4. A linear and high accuracy correlation of the UTS was obtained when the punch force vs. punch displacement (F-X) curve slope within membrane stretching regime is normalized by the sample thickness. 5. It can be concluded that high strength, adequate electrical conductivity and good formability properties achieved by aging heat treated UFG Cu-Cr-Zr alloys make them a good candidate to be used in micro-forming applications.

19

Acknowledgements I would like to thank Dr. Gencaga Purcek and Harun Yanar for their contributions on performing the ECAP experiments and discussing the results of this study. Also, I would like to thank Sağlam Metal Inc., Istanbul for their support in kindly supplying the Cu-Cr-Zr alloy.

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References [1] I.S. Batra, G.K. Dey, U.D. Kulkarni, S. Banerjee, Microstructure and properties of a Cu–Cr– Zr alloy, J. Nucl. Mater., 299 (2001) 91-100. [2] J. Davis, G. Kalinin, Material properties and design requirements for copper alloys used in ITER, J. Nucl. Mater., 258 (1998) 323-328. [3] M. Kulczyk, B. Zysk, M. Lewandowska, K.J. Kurzydłowski, Grain refinement in CuCrZr by SPD processing, Phys. Status. Solidi., 207 (2010) 1136-1138. [4] Q. Liu, X. Zhang, Y. Ge, J. Wang, J.-Z. Cui, Effect of processing and heat treatment on behavior of Cu-Cr-Zr alloys to railway contact wire, Metall. Mater. Trans. A, 37 (2006) 3233-3238. [5] G. Purcek, H. Yanar, O. Saray, I. Kraman, H.J. Maier, Effect of precipitation on mechanical and wear properties of ultrafine-grained Cu-Cr-Zr alloy, Wear, 311 (2014) 149-158. [6] N. Takata, Y. Ohtake, K. Kita, K. Kitagawa, N. Tsuji, Increasing the ductility of ultrafinegrained copper alloy by introducing fine precipitates, Scripta Mater., 60 (2009) 590-593. [7] K.X. Wei, W. Wei, F. Wang, Q.B. Du, I.V. Alexandrov, J. Hu, Microstructure, mechanical properties and electrical conductivity of industrial Cu–0.5%Cr alloy processed by severe plastic deformation, Mat. Sci. Eng. A-Struct., 528 (2011) 1478-1484. [8] I. Shakhova, Z. Yanushkevich, I. Fedorova, A. Belyakov, R. Kaibyshev, Grain refinement in a Cu–Cr–Zr alloy during multidirectional forging, Mat. Sci. Eng. A-Struct., 606 (2014) 380389. [9] D.J. Edwards, B.N. Singh, S. Tähtinen, Effect of heat treatments on precipitate microstructure and mechanical properties of a CuCrZr alloy, J. Nucl. Mater., 367–370, Part B (2007) 904909.

21

[10] R.K. Islamgaliev, K.M. Nesterov, J. Bourgon, Y. Champion, R.Z. Valiev, Nanostructured Cu-Cr alloy with high strength and electrical conductivity, J. Appl. Phys., 115 (2014) 194301. [11] E. Batawi, D. Morris, M. Morris, Effect of small alloying additions on behaviour of rapidly solidified Cu–Cr alloys, Mater. Sci. Tech., 6 (1990) 892-899. [12] N. Tang, D. Taplin, G. Dunlop, Precipitation and aging in high-conductivity Cu–Cr alloys with additions of zirconium and magnesium, Mater. Sci. Tech., 1 (1985) 270-275. [13] K. Leedy, J. Stubbins, B. Singh, F. Garner, Fatigue behavior of copper and selected copper alloys for high heat flux applications, J. Nucl. Mater., 233 (1996) 547-552. [14] A. Tavassoli, Materials design data for fusion reactors, J. Nucl. Mater., 258 (1998) 85-96. [15] P. Liu, J.H. Su, Q.M. Dong, H.J. Li, Microstructure and properties of Cu-Cr-Zr alloy after rapidly solidified, aging and solid solution aging, J. Mater. Sci. Technol., 21 (2005) 475478. [16] N. Gao, T. Tiainen, Y. Ji, L. Laakso, Control of microstructures and properties of a phosphorus-containing Cu-0.6 Wt.% Cr alloy through precipitation treatment, Journal of Materials Engineering and Performance, 9 (2000) 623-629. [17] H. Zhou, J. Zhong, X. Zhou, Z. Zhao, Q. Li, Microstructure and properties of Cu–1.0 Cr–0.2 Zr–0.03 Fe alloy, Mat. Sci. Eng. A-Struct., 498 (2008) 225-230. [18] G. Weatherly, P. Humble, D. Borland, Precipitation in a Cu-0, 55 wt.% Cr alloy, Acta Metall. Mater., 27 (1979) 1815-1828. [19] Y. Huang, J. Robson, P. Prangnell, The formation of nanograin structures and accelerated room-temperature theta precipitation in a severely deformed Al–4wt.% Cu alloy, Acta Mater., 58 (2010) 1643-1657.

22

[20] W. Kim, C. Chung, D. Ma, S. Hong, H. Kim, Optimization of strength and ductility of 2024 Al by equal channel angular pressing (ECAP) and post-ECAP aging, Scripta Mater., 49 (2003) 333-338. [21] A. Vinogradov, K. Kitagawa, V. Kopylov, Mater. Sci. Forum., Trans. Tech. Publ., 2006, 811-816. [22] C. Xu, Q. Wang, M. Zheng, J. Zhu, J. Li, M. Huang, Q. Jia, Z. Du, Microstructure and properties of ultra-fine grain Cu–Cr alloy prepared by equal-channel angular pressing, Mat. Sci. Eng. A-Struct., 459 (2007) 303-308. [23] A. Vinogradov, Y. Suzuki, T. Ishida, K. Kitagawa, V.I. Kopylov, Effect of chemical composition on structure and properties of ultrafine grained Cu-Cr-Zr alloys produced by equal-channel angular pressing, Mater. Trans., 45 (2004) 2187-2191. [24] A. Vinogradov, V. Patlan, Y. Suzuki, K. Kitagawa, V. Kopylov, Structure and properties of ultra-fine grain Cu–Cr–Zr alloy produced by equal-channel angular pressing, Acta Mater., 50 (2002) 1639-1651. [25] K.V. León, M. Munoz-Morris, D.G. Morris, Optimisation of strength and ductility of Cu– Cr–Zr by combining severe plastic deformation and precipitation, Mat. Sci. Eng. A-Struct., 536 (2012) 181-189. [26] R. Mishnev, I. Shakhova, A. Belyakov, R. Kaibyshev, Deformation microstructures, strengthening mechanisms, and electrical conductivity in a Cu–Cr–Zr alloy, Mat. Sci. Eng. A-Struct., 629 (2015) 29-40. [27] O. Saray, G. Purcek, I. Karaman, H. Maier, Improvement of formability of ultrafine-grained materials by post-SPD annealing, Mat. Sci. Eng. A-Struct., 619 (2014) 119-128.

23

[28] R.Z. Valiev, M.J. Zehetbauer, Y. Estrin, H.W. Höppel, Y. Ivanisenko, H. Hahn, G. Wilde, H.J. Roven, X. Sauvage, T.G. Langdon, The innovation potential of bulk nanostructured materials, Adv. Eng. Mater., 9 (2007) 527-533. [29] G. Purcek, H. Yanar, M. Demirtas, Y. Alemdag, D.V. Shangina, S.V. Dobatkin, Optimization of strength, ductility and electrical conductivity of CU-Cr-Zr alloy by combining multi-route ECAP and aging, Mat. Sci. Eng. A-Struct., 649 (2015) 114-122. [30] O. Saray, G. Purcek, I. Karaman, H.J. Maier, Formability of ultrafine-grained interstitial-free steels, Metallurgical and Materials Transactions a-Physical Metallurgy and Materials Science, 44A (2013) 4194-4206. [31] Y. Champion, Y. Bréchet, Effect of grain size reduction and geometrical confinement in fine grained copper: Potential applications as a material for reversible electrical contacts, Adv. Eng. Mater., 12 (2010) 798-802. [32] E.C. Moreno-Valle, M.A. Monclus, J.M. Molina-Aldareguia, N. Enikeev, I. Sabirov, Biaxial deformation behavior and enhanced formability of ultrafine-grained pure copper, Metall. Mater. Trans. A, 44 (2013) 2399-2408. [33] I. Sağlam, D. Özyürek, K. Çetinkaya, Effect of ageing treatment on wear properties and electrical conductivity of Cu-Cr-Zr alloy, Bull. Mater. Sci., 34 (2011) 1465-1470. [34] Y. Iwahashi, Z. Horita, M. Nemoto, T.G. Langdon, An investigation of microstructural evolution during equal-channel angular pressing, Acta Mater., 45 (1997) 4733-4741. [35] R.Z. Valiev, A.V. Korznikov, R.R. Mulyukov, Structure and properties of ultrafine-grained materials produced by severe plastic deformation, Mat. Sci. Eng. A-Struct., 168 (1993) 141148. [36] R. Valiev, Y. Estrin, Z. Horita, T. Langdon, M. Zechetbauer, Y. Zhu, Producing bulk ultrafine-grained materials by severe plastic deformation, JOM, 58 (2006) 33-39. 24

[37] R.Z. Valiev, T.G. Langdon, Principles of equal-channel angular pressing as a processing tool for grain refinement, Prog. Mater. Sci., 51 (2006) 881-981. [38] N. Hansen, X. Huang, D.A. Hughes, Microstructural evolution and hardening parameters, Mat. Sci. Eng. A-Struct., 317 (2001) 3-11. [39] E. Fleury, J.S. Ha, Small punch tests to estimate the mechanical properties of steels for steam power plant: I. Mechanical strength, Int. J. Pres. Ves. Pip., 75 (1998) 699-706. [40] D. Sekban, O. Saray, S. Aktarer, G. Purcek, Z. Ma, Microstructure, mechanical properties and formability of friction stir processed interstitial-free steel, Mat. Sci. Eng. A-Struct., 642 (2015) 57-64. [41] S.D. Norris, J.D. Parker, Deformation processes during disc bend loading, Mater. Sci. Tech., 12 (1996) 163-170. [42] M. Abendroth, M. Kuna, Determination of deformation and failure properties of ductile materials by means of the small punch test and neural networks, Comp. Mater. Sci., 28 (2003) 633-644. [43] R. Becker, Effects of strain localization on surface roughening during sheet forming, Acta Mater., 46 (1998) 1385-1401. [44] M. Abendroth, M. Kuna, Determination of ductile material properties by means of the small punch test and neural networks, Adv. Eng. Mater., 6 (2004) 536-540. [45] T.E. García, C. Rodríguez, F.J. Belzunce, C. Suárez, Estimation of the mechanical properties of metallic materials by means of the small punch test, J. Alloy. Compd., 582 (2014) 708717. [46] C.Y. Yu, P.W. Kao, C.P. Chang, Transition of tensile deformation behaviors in ultrafinegrained aluminum, Acta Mater., 53 (2005) 4019-4028.

25

[47] G. Purcek, O. Saray, I. Karaman, H.J. Maier, High strength and high ductility of ultrafinegrained, interstitial-free steel produced by ecae and annealing, Metall. Mater. Trans. A, 43 (2012) 1884-1894. [48] Y. Wang, M. Chen, F. Zhou, E. Ma, High tensile ductility in a nanostructured metal, Nature, 419 (2002) 912-915. [49] Y.M. Wang, E. Ma, Three strategies to achieve uniform tensile deformation in a nanostructured metal, Acta Mater., 52 (2004) 1699-1709. [50] T.S. Wang, Z. Li, B. Zhang, X.J. Zhang, J.M. Deng, F.C. Zhang, High tensile ductility and high strength in ultrafine-grained low-carbon steel, Mat. Sci. Eng. A-Struct., 527 (2010) 2798-2801. [51] K. Kumar, A. Pooleery, K. Madhusoodanan, R. Singh, J. Chakravartty, R. Shriwastaw, B. Dutta, R. Sinha, Evaluation of ultimate tensile strength using Miniature Disk Bend Test, J. Nucl. Mater., 461 (2015) 100-111. [52] M. Suzuki, M. Eto, K. Fukaya, Y. Nishiyama, T. Kodaira, T. Oku, M. Adachi, A. Umino, I. Takahashi, T. Misawa, Y. Hamaguchi, Evaluation of toughness degradation by small punch (SP) tests for neutron-irradiated 214Cr-1Mo steel, J. Nucl. Mater., 179–181, Part 1 (1991) 441-444. [53] Y. Ruan, P. Spätig, M. Victoria, Assessment of mechanical properties of the martensitic steel EUROFER97 by means of punch tests, J. Nucl. Mater., 307–311, Part 1 (2002) 236239. [54] G.E. Lucas, A. Okada, M. Kiritani, Parametric analysis of the disc bend test, J. Nucl. Mater., 141 (1986) 532-535.

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Table captions: Table 1. Mechanical properties of CG and UFG Cu-Cr-Zr alloy from uniaxial tension tests (y: yield strength, UTS: ultimate tensile strength, u: uniform elongation, f: elongation to failure). Table 2. Mechanical properties of CG and UFG Cu-Cr-Zr alloy from Erichsen tests (Ei : Erichsen index, FEi: load corresponding to Erichsen index, dF/dX)m: Slope of the F-X curve within the membrane stretching regime)

Table 1. Mechanical properties of CG and UFG Cu-Cr-Zr alloy from uniaxial tension tests (y: yield strength, UTS: ultimate tensile strength, u: uniform elongation, f: elongation to failure).

Condition

Aging

Aging

temperature

time

y

u

f

(%)

(%)

UTS (MPa) (MPa)

(min)

(C) Quenched

-

-

218±11

274±19

22.0±3.0

36.0±4.0

Aged

475

300

295±2

380±18

9.0±1.0

16.0±2.0

-

-

516±8

572±18

2.0±0.1

15.0±1.0

30

618±11

627±8

0.9±0.1

9.5±1.1

90

601±5

611±4

1.0±0.1

11.1±1.5

240

557±16

601±11

1.6±0.2

17.0±3.0

CG

Asprocessed UFG Aged

450

27

Table 2. Mechanical properties of CG and UFG Cu-Cr-Zr alloy from Erichsen tests (Ei: Erichsen index, FEi: punch load corresponding to Erichsen index)

Aging

Aging

temperature

time

(°C)

(min)

Quenched

-

Aged As-processed

Condition

Ei (mm)

FEi (N)

-

4.68±0.12

3535±61

475

300

3.95±0.17

5391±71

-

-

4.05±0.12

6955±37

30

4.15±0.09

8417±58

90

4.20±0.02

8291±16

240

4.23±0.11

7604±38

CG

UFG Aged

450

Figure captions: Fig.1. Schematic illustrations showing: (a) the position of the samples in the ECAE processed billet and (b) miniaturized Erichsen test die.

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Fig.2 Microstructure of ultra-fine grained Cu-Cr-Zr alloy in the condition of (a) as 8Bc ECAE processed, and subsequently aging heat treated at 450 C for: (b)-(c) 30 min, (d)-(e) 90 min and (f)-(g) 240 min.

29

30

Fig.3. Engineering stress-engineering strain curves of CG and UFG Cu-Cr-Zr alloys.

31

Fig.4. (a) Punch load (F) vs. displacement (X) curves of Cu-Cr-Zr alloy obtained via Erichsen test and (b) dF/dX vs. X curves representing displacement within the membrane stretching regime.

32

33

Fig.5. Dome free surface appearance of the Erichsen tested CG samples in (a) – (b) quenched condition, and (c)-(d) aging heat treated condition.

Fig.6. Dome free surface appearance of the Erichsen tested UFG samples in (a)-(b) as-processed condition, and aging heat treated at 450°C for (c)-(d) for 30 min, (e)-(f) 90 min, (g)-(h) 240 min.

34

35

Fig.7. Correlation between UTS and (dF/dX)mt-1.

36