Microstructural evolution and mechanical properties of a two-phase Cu Ag alloy processed by high-pressure torsion to ultrahigh strains

Transcription

1 Available online at Acta Materialia 59 (2011) Microstructural evolution and mechanical properties of a two-phase Cu Ag alloy processed by high-pressure torsion to ultrahigh strains Y.Z. Tian a, S.D. Wu a, Z.F. Zhang a,, R.B. Figueiredo b, N. Gao c, T.G. Langdon c,d a Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang , China b Department of Metallurgical and Materials Engineering, Federal University of Minas Gerais, Belo Horizonte, MG , Brazil c Materials Research Group, School of Engineering Sciences, University of Southampton, Southampton SO17 1BJ, UK d Departments of Aerospace & Mechanical Engineering and Materials Science, University of Southern California, Los Angeles, CA , USA Received 17 October 2010; received in revised form 24 December 2010; accepted 4 January 2011 Abstract Disks of a coarse-grained Cu 28 wt.% Ag alloy were processed by high-pressure torsion up to 20 revolutions to reveal the microstructural evolution and mechanical properties. The eutectic shows a faster evolution process than the Cu matrix. A banded structure forms in the Cu matrix, and both the eutectic spacing and the band width decrease with increasing shear strain. After 20 revolutions, the substructure may even diminish in the Cu matrix. The microhardness increases with increasing revolutions, and a saturation microhardness is ultimately achieved. After 20 revolutions, the tensile strength was improved to 1420 MPa, and the failure mode of the sample was transferred from necking to full shearing without plasticity. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Cu Ag alloy; High-pressure torsion; Microstructure; Hardness; Fracture 1. Introduction The application of severe plastic deformation (SPD) to coarse-grained metals has now become accepted as a valuable procedure for the fabrication of bulk solids with ultrafine-grained (UFG) or nanocrystalline (NC) sizes [1,2]. Several SPD techniques are available [3], and significant interest has centered on high-pressure torsion (HPT) [4], equal-channel angular pressing (ECAP) [5,6], dynamic plastic deformation (DPD) [7,8] and accumulative rollbonding (ARB) [9,10]. Of these various techniques, HPT is especially attractive, because it is easy to conduct and has the ability to impose exceptionally high strains to produce extremely small grain sizes. Corresponding author. Tel.: address: zhfzhang@imr.ac.cn (Z.F. Zhang). There are several studies where microhardness measurements have been used to evaluate the extent of any microstructural evolution occurring in HPT [11 15]. All these experiments are mutually consistent, and there is a gradual evolution towards a homogeneous distribution of microhardness values with increasing pressure and/or strain. Since experiments on samples processed by ECAP have shown good correlation between microhardness measurements and microstructural observations undertaken using transmission electron microscopy (TEM) [16], it is reasonable to anticipate that the internal microstructure also evolves towards homogeneity when processing by HPT. An evolution towards microstructural homogeneity in HPT has been predicted by making use of strain gradient plasticity modeling [17]. Nevertheless, this result is not consistent with the basic principles of HPT, where the shear strain c is given by the relationship [18,19] /$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi: /j.actamat

2 2784 Y.Z. Tian et al. / Acta Materialia 59 (2011) c ¼ 2pNr h where r and h are the radius and height (or thickness) of the disk, respectively, and N is the number of revolutions. The equivalent von Mises strain is given by e vm ¼ p c ffiffi 3 It follows from Eq. (1) that the shear strain is zero at the center of the disk and, in principle at least, the microstructure across the disk should remain inhomogeneous. It has been suggested that the microhardness homogeneity in HPT disks may arise because of a misalignment of the axes of the anvils or other deviations from the idealized HPT conditions [20]. This apparent dichotomy suggests that more information is critically needed concerning the nature of the flow process occurring in HPT. Three recent reports provide information on the shearing process in HPT. First, experiments on a two-phase duplex stainless steel revealed the development of swirls and vortices during HPT processing where the appearance of these irregularities has similarities to the well-established Kelvin Helmholtz instability observed in fluid flow when there are significant velocity gradients within the liquid [21,22]. Second, the processing of a two-phase Zn Al eutectoid alloy showed the occurrence of agglomeration and banding around the peripheries of the disks during the initial stages of HPT [23]. Third, recent results on a Cu 28 wt.% Ag alloy also revealed an inhomogeneous evolution process and evidence of straining near the center during HPT processing [24]. These results have provided powerful evidence for microstructural evolution during HPT where the decorative phase with great contrast is the hallmark of twophase alloys. Accordingly, it is relevant to reveal the microstructural evolution of HPT processing by tracing the morphological changes of two-phase alloys. Since Cu Ag alloys containing two ductile phases can be processed by conventional cold drawing, high-strength microcomposites have been produced [25]. It is reported that the tensile strength increases continuously with relation to the draw ratio and approaches an extremely high value [26,27]. However, the saturation strength of Cu Ag alloys has not been reported to date, and this may arise owing to drawbacks in conventional processing because of limited strains and cracking at high draw ratios. By contrast, the HPT technique may be applicable for highly strengthening the Cu Ag alloys, since exceptionally high strains can be imposed. A coarse-grained Cu 28 wt.% Ag alloy was selected for this study because this alloy contains two ductile phases and the morphological changes of the decorative eutectic region may be conveniently traced when processing by HPT. Accordingly, the present investigation was initiated to provide further information on the microstructural evolution and the extent of any homogenizing ð1þ ð2þ process in the two-phase Cu Ag alloy during HPT. In addition, the mechanical properties of the Cu Ag alloy processed by HPT to ultrahigh strains were also examined. 2. Experimental material and procedures The Cu 28 wt.% Ag binary alloy was fabricated from % purity Cu and 99.99% purity electrolytic Ag using the procedure described earlier [28]. The alloy is hypoeutectic and generally contains two components: eutectic components composed of Cu and Ag phases and a proeutectic Cu dendrite embedded with Ag precipitates. The Cu and Ag phases always have the cube-on-cube orientation in the proeutectic component, but they have the orientation relationship only in selected eutectic areas [29]. In addition, the dendrite and the eutectic component may also have the orientation relationship [28,30]. Detailed information on the microstructures and plastic deformation of this alloy and a Cu 16 wt.% Ag alloy is now available [28,30]. Disks with diameter 10 mm and thickness mm were prepared, and they were processed by HPT at room temperature under quasi-constrained conditions [4] using the experimental facility and procedure described earlier [31] except only that a lubricant was not placed around the peripheries of the depressions on the upper and lower anvils. The disks were processed through 1, 3, 5, 10 and 20 revolutions under an imposed pressure of 6 GPa. Following HPT, detailed microstructural observations by scanning electron microscopy (SEM) were undertaken at selected positions on transverse sections of the disks (see Fig. 1a) using a LEO SUPRA 35 microscope operated at 20 kv. Images were also recorded by quadrant back scattering detection (QBSD)-SEM to reveal the macroscopic deformation morphologies of the eutectic regions throughout the disks. The microstructure characterization was performed by TEM on an FEI Tecnai F20 operated at 200 kv. Thin foils for TEM characterizations were cut from the sectional areas at selected distances from the center, as shown in Fig. 1b. The TEM samples were first mechanically ground to 30 lm thickness and then dimpled and thinned by ion milling using a GATAN PIPS 691 at 4 kv. It should be noted that the samples were cooled at liquid nitrogen temperature during milling, and the thinning process was discontinued every 10 min to minimize any heat effect on the deformed microstructure. Microhardness and tensile tests were also conducted. Before the microhardness measurements, the disks were ground and mechanically polished to a mirror-like surface. Individual values of the Vickers microhardness were measured across the diameter of each disk in incremental steps of mm up to a distance of 4.5 mm, as depicted in Fig. 1c. Near the center, eight values of Vickers microhardness were also measured. In addition, tensile tests were conducted at an initial strain rate of s 1 using an Instron E1000 testing machine. Tensile specimens with a dog-bone shape with gauge length 2 mm,

3 Y.Z. Tian et al. / Acta Materialia 59 (2011) Fig. 1. Schematic illustration of: (a) the selected positions for microstructure observation on the transverse section of the HPT disk; (b) TEM specimens cut from the disk and tensile specimens cut from the off-center position; (c) TEM specimens cut from the sectional area of the disk; (d) microhardness measurements. width 0.7 mm and thickness 0.5 mm were cut from the disks, as shown in Fig. 1d. After tensile fracture, the samples were observed by SEM to reveal the deformation and fracture morphologies. 3. Experimental results 3.1. Microstructural evolution In this section, the microstructures on the transverse sections of the disks are examined, because the plastic flow of the material is then well revealed. A schematic illustration of the view plane is shown in Fig. 1a, and the selected positions for microstructural characterization are also indicated. Fig. 2a 1 e 1 shows the QBSD-SEM images on the deformation morphologies of the transverse sections of the disks strained through 1, 3, 5, 10 and 20 revolutions of HPT processing, respectively. Fig. 2a 2 e 2 shows the corresponding microstructures near the center of each disk, as indicated by the rectangles. There is a coarse eutectic region near the center of the disk strained to 1 revolution, whereas the eutectic is increasingly refined with increasing distance from the center, as shown in Fig. 2a 1. At a higher magnification, it is found that the eutectic in Fig. 2a 2 has lost the as-cast morphology, and the net-like eutectic is elongated owing to shear deformation. With increasing numbers of revolutions, it is apparent that the eutectic near the center becomes severely sheared and elongated, with these elongated regions lying increasingly parallel to each other, as shown in Fig. 2b 2 d 2. Furthermore, it is clear that the eutectic spacing decreases drastically with increasing numbers of revolutions. By contrast, Fig. 2e 2 shows that the distribution of the eutectic is inhomogeneous near the center of the disk strained to 20 revolutions, and this appears to result from the inhomogeneous shear deformation during the straining process [24]. Since the microstructures removed from the centers of the disks are greatly refined, three positions at distances of 0 mm, 2 mm and 4.5 mm were chosen to obtain detailed information on the microstructural evolution. Detailed QBSD-SEM images on the microstructures of the selected areas are shown in Fig. 3. After 1 revolution, it is found that even the eutectic near the center of the disk has suffered some shear deformation, as shown in Fig. 3a 1. With increasing numbers of revolutions, the eutectic experiences more intense plastic deformation such that after 5 and 10 revolutions the Cu and Ag phases in the eutectic

4 2786 Y.Z. Tian et al. / Acta Materialia 59 (2011) Fig. 2. QBSD-SEM images of the transverse sections of the HPT disks and magnified images of the central regions after straining through: (a 1 and a 2 )1 revolution; (b 1 and b 2 ) 3 revolutions; (c 1 and c 2 ) 5 revolutions; (d 1 and d 2 ) 10 revolutions; (e 1 and e 2 ) 20 revolutions. are severely sheared and elongated, as shown in Fig. 3c 1 and d 1, respectively. By contrast, Fig. 3e 1 shows some shear bands crossing the eutectic and dendrite near the center of the disk strained to 20 revolutions. This latter observation is significant because the theoretical shear strain is zero in the central region according to Eq. (1). It is noted also that the shear bands appear only in selected areas in the peripheral region of the disk after 1 revolution, thereby implying that the central region of the disk strained to 20 revolutions has suffered considerable plastic deformation. The results from the center of the disk strained to different revolutions indicate that plastic flow has occurred in this region. At a distance of 2 mm from the centers of the disks, it is expected that substantial microstructure refinement occurs as revealed in Fig. 3a 2 e 2. After 1 revolution, the eutectic region in Fig. 3a 2 is sheared and elongated, although it remains coarse. With increasing numbers of revolutions, the eutectic is further refined. In addition, a large number of shear bands also appear after 5 revolutions, as shown in Fig. 3c 2. When the disk is further strained to 20 revolutions, it is found that the eutectic has fully transformed into a homogeneous fibrous microstructure, as shown in Fig. 3e 2. At the peripheral region of the disk at 4.5 mm, the shear strain is much higher than at a distance of 2 mm, and it is expected that greater refinement will occur. After 1 revolution, the eutectic is finer at 4.5 mm than at 2 mm, although it remains coarse. After 3 revolutions, however, the eutectic is severely sheared and elongated to a fibrous morphology, although some coarse eutectic remains. Moreover, shear bands are prevalent, as shown in Fig. 3b 3. When the disk is further strained to 5 revolutions, the microstructure is similar to that after 3 revolutions, except that the eutectic is finer and more homogeneous. Finally, it is nearly impossible to distinguish the single eutectic by SEM after 10 and 20 revolutions, as shown in Figs. 3d 3 and 6e 3, respectively. Since the microstructure shows a gradual evolution with increasing distance from the center or numbers of revolutions, TEM was applied to reveal the inherent characteristics. After 1 revolution, it is found that the Cu matrix contains banded structures with the dimensions ranging from 20 to 100 nm, as shown in Fig. 4a. Inside the eutectic, which consists of alternating Cu and Ag phases, the Cu or Ag width is even smaller. Fig. 4b shows a eutectic region where the microstructures are inhomogeneous and the band width ranges from 10 to 50 nm. In addition, some Cu and Ag phases cannot be distinguished owing to severe shear strain. A narrow shear band was also observed in the eutectic region, as revealed in Fig. 4c. Fig. 4d shows the microstructures near the interface between the eutectic and the Cu matrix. It seems that the width of the Cu or

5 Y.Z. Tian et al. / Acta Materialia 59 (2011) Fig. 3. Microstructures taken from the selected positions on the transverse sections of the HPT disks after straining through: (a 1 a 3 ) 1 revolution; (b 1 b 3 ) 3 revolutions; (c 1 c 3 ) 5 revolutions; (d 1 d 3 ) 10 revolutions and (e 1 e 3 ) 20 revolutions. Ag phase in the eutectic is much smaller than the band width of the Cu matrix, and this is reasonable, since the microstructural scale is finer in the as-received eutectic than in the Cu matrix. After 5 revolutions, the microstructures were examined at a distance of 3.5 mm, and they are shown in Fig. 5. Fig. 5a shows an alternating distribution of Cu matrix and eutectic. The width of the Cu matrix (eutectic spacing) can approach 50 nm, wherein only one lamellar boundary was observed. In the lower region, the Cu matrix contains many lamellar boundaries, and the band width ranges from 20 to 50 nm, which is smaller than after only 1 revolution. Inside the eutectic region, in contrast, small domains with several nanometers in size were observed. However, some eutectic regions are still not fully refined, as shown in Fig. 5b. Fig. 5c shows a shear band crossing a eutectic region. Since the eutectic has been severely refined, the microstructures inside the shear band do not show significant contrast to those outside the shear band, as indicated in Fig. 5d. After 20 revolutions, it is known from Fig. 3e 2 that the eutectic is fully refined to a fibrous condition. In this case, the microstructures were inspected at distances of 2.5 and 4.5 mm, and they are shown in Fig. 6. Fig. 6a and b shows the bright field and dark field images of a selected region, respectively, and a shear band can be found. The eutectic spacing was estimated to be smaller than 100 nm, which is much smaller than after 5 revolutions. In some areas the Cu matrix is too fine to be distinguished. Fig. 6c shows alternating Cu matrix and eutectic, with the eutectic indicated by red arrow, and the internal band width is smaller than 50 nm. Inside the eutectic, small domains with several nanometers in size were found, as shown in Fig. 6d. At a distance of 4.5 mm from the center of the disk, Fig. 6e and f shows a shear band 400 nm in width, which is much thicker than that observed in the disk strained through 5 revolutions. Outside the shear band, only limited banded Cu matrix can be distinguished, and the eutectic spacing is even smaller than 20 nm, whereas inside the shear band both the eutectic and the Cu matrix are fully mixed. Detailed information on the nucleation and thickening of shear bands is now available [7]. During conventional cold drawing, it is found that the tensile strength increases with decreasing transverse size of the Cu phase (eutectic spacing) in Cu Ag alloys [27]. Similar results were also obtained in other Cu X (Nb [32], Cr[33], Fe[34]) alloys. In the present study, the eutectic morphology of the Cu Ag alloy after HPT processing varies with the distance from the center. Therefore, the eutectic spacings were measured at distances of 0, 0.25, 0.5, 1, 2, 3 and 4.5 mm on the transverse sections of the disks, as illustrated in Fig. 1a: it should be noted that for

6 2788 Y.Z. Tian et al. / Acta Materialia 59 (2011) Fig. 4. TEM images taken at a distance of 3.5 mm from the center of the HPT disk after 1 revolution: the Ag and Cu phases are indicated by red and blue arrows, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) the disks strained to 10 and 20 revolutions the investigated areas are limited to 2 and 1 mm, respectively. This is because it was impossible to distinguish the single eutectic at greater distances. Since the strains of the disks at different positions and under different revolutions should be different according to Eqs. (1) and (2), the eutectic spacings were plotted as a function of the corresponding equivalent strain, as shown in Fig. 7. It is apparent that the eutectic spacing decreases continuously with increasing equivalent strain. In practice, this is similar to results on the conventional cold-drawn Cu Ag alloys [27], where the tensile strength increases to a high level when the eutectic spacing is very small. It is reasonable to anticipate that the Cu Ag alloy in the present study may exhibit exceptional mechanical properties and accordingly, as described in the following sections, the Vickers microhardness was measured and tensile tests were conducted Variation in Vickers microhardness across the disks The values of the Vickers microhardness recorded across the diameter of each disk are plotted against the distance from the center in Fig. 8, where each point is an average of the eight values recorded at the same distance from the center, as illustrated in Fig. 1c, and the error bar is the standard deviation. It is apparent that the microhardness increases significantly throughout the disk after 1 revolution. Careful inspection shows that the microhardness increases by a factor of 3 at the peripheral region of the disk where this is much higher than for a Cu 8 wt.% Ag alloy processed by ECAP [35]. By contrast, the microhardness near the center of the disk also increases to some extent, although the theoretical shear strain is then zero. This same trend is prevalent in many materials processed by HPT [4,11 15,23,31,36 41]. After 3 revolutions, the microhardness increases to a higher level from the center to the outer region of the disk, with no evidence of any saturation. When the disk is further strained to 5 and 10 revolutions, it is found that the microhardness continues to increase at the selected distances, although the increase is smaller in the center. It is of interest to note that saturation microhardness is not achieved even after 10 revolutions. In terms of the microhardness results of single-phase alloys [19,20,38,39,42] where saturation microhardness generally occurs within an equivalent strain of 50, the present behavior is different and will be discussed in the following section. During conventional cold drawing, the tensile strength increases with relation to the draw ratio, but the saturation value has not been reported so far [25 27]. This may be due to the limitation of the processing ability to high strains and the substantial reduction in the dimensional size.

7 Y.Z. Tian et al. / Acta Materialia 59 (2011) Fig. 5. TEM images taken at a distance of 3.5 mm from the center of the HPT disk after 5 revolutions: the Ag and Cu phases are indicated by red and blue arrows, respectively. (d) is the dark field image of the square region in (c). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) Therefore, the HPT disk was further processed to 20 revolutions to check the possibility of a saturation strength. In contrast to the disk strained through 10 revolutions, it is found that the microhardness continues to increase throughout the disk. Near the center, the microhardness is much higher than after 10 revolutions and, in fact, it is comparable with the microhardness at the peripheral region of the disk after 1 revolution. By contrast, the microhardness gradually levels off with increasing distance from the center, thereby indicating that saturation microhardness has been achieved at the outer region of the disk after 20 revolutions. When the microhardness values recorded in the five disks strained to different revolutions are plotted as a function of the equivalent strain, it is found that the microhardness and the equivalent strain give a unique relationship, apart from only very near the center, as shown in Fig. 9. It is worthwhile examining several trends that is evident from inspection of Fig. 9. First, the Vickers microhardness increases continuously with relation to the equivalent strain over a wide range, and it fails to saturate until the equivalent strain approaches 330. This indicates that the two-phase Cu Ag alloys possesses a stronger strain hardening ability than pure metals [12,19,38 40] during HPT processing. It is also obvious that the strain threshold for a saturation microhardness is much larger than in single-phase alloys [19,20,38 40,42]. Second, the saturation microhardness obtained for the Cu Ag alloy processed by HPT in the present study confirms the existence of a saturation microhardness/strength. Since HPT can impose extremely high strains on a sample, this makes it feasible to investigate the mechanical properties of Cu Ag alloys deformed to ultrahigh strains. It should be noted that a saturation strength of 600 MPa was obtained for Cu 8 wt.% Ag alloy processed by ECAP via route B C for 4 passes, where this is much lower than after processing by conventional cold drawing [25,27]. This low tensile strength appears to arise as a result of the inhibition of the strengthening ability of the eutectic [35]. Third, as shown by the inset in Fig. 9, the microhardness values near the center do not fulfill the relation between microhardness and strain. This is understandable according to Eq. (1), where the shear strain is zero, but nevertheless there is some plastic deformation and strain hardening in the central region. In the conventional cold-drawn Cu Ag alloys, the tensile strength increases continuously as a function of the

8 2790 Y.Z. Tian et al. / Acta Materialia 59 (2011) Fig. 6. TEM images taken at distances of (a d) 2.5 mm and (e and f) 4.5 mm from the center of the HPT disk after 20 revolutions: the red arrows indicate the eutectic region in (c). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) draw ratio, where the eutectic spacing decreases drastically and falls to the submicrometer or nanometer range [27]. In the present study, the microhardness values in Fig. 8 and the corresponding eutectic spacings in Fig. 7 were replotted in Fig. 10. It is evident that the microhardness increases essentially linearly with decreasing eutectic spacing when expressed in a logarithmic form Tensile properties and fracture behavior Tensile engineering stress strain curves before and after HPT are shown in Fig. 11. The as-received disk exhibits high elongation but very low tensile strength, which is characteristic of coarse-grained materials. After 1 revolution, the tensile strength increases significantly to 900 MPa; however, the uniform elongation is extremely small, and the elongation to failure is also reduced drastically in contrast to the as-received sample, owing to a deterioration in the strain hardening ability. When the disks are further strained to 5, 10 and 20 revolutions, the tensile strength continues to increase, but the elongation becomes negligible, displaying a typical brittle fracture behavior. The tensile strength is as high as 1420 MPa after 20 revolutions, and this is comparable

9 Y.Z. Tian et al. / Acta Materialia 59 (2011) Fig. 7. Eutectic spacing taken from the selected positions on the transverse sections of the HPT disks as illustrated in Fig. 1a. The datum points for specimens after 10 and 20 revolutions are limited to distances of 2 mm and 1 mm from the center of the HPT disk, respectively. Fig. 9. The average Vickers microhardness plotted against the equivalent strain for the five HPT disks after straining through different numbers of revolutions. Fig. 8. Variation of the average Vickers microhardness recorded across the diameter of the HPT disks after straining to different revolutions: the broken line shows the average Vickers microhardness of the as-received material. Fig. 10. Vickers microhardness values plotted against the eutectic spacing for the five HPT disks after straining through different numbers of revolutions. with the value in conventional cold drawing to very high draw ratios [27,43]. In view of the abrupt transition of the elongations after different numbers of revolutions, it is anticipated that the fracture mode will change during tensile testing. Accordingly, the fracture surfaces were examined as shown in Fig. 12. A combination of necking and shear fracture is observed in Fig. 12a for the sample after 1 revolution, where this failure behavior is consistent with the tensile stress strain curves and is also similar to materials processed by ECAP for different numbers of passes [6,35]. This indicates that the Cu Ag alloy after 1 revolution of HPT can display only a relatively low tensile strength (900 MPa), following by a weak work-hardening and necking until final shear failure. When the disks are further strained to 5, 10 and 20 revolutions, shear fracture always Fig. 11. Tensile engineering stress strain curves of the Cu Ag alloy before and after HPT.

10 2792 Y.Z. Tian et al. / Acta Materialia 59 (2011) Fig. 12. Tensile fracture morphologies of the Cu Ag alloy processed by HPT for 1, 5, 10 and 20 revolutions: 1r, 5r, 10r and 20r correspond to the numbers of revolutions. occurs abruptly during the tensile test, without any indication of necking, as shown in Fig. 12b d. This typical brittle shear failure is in accordance with the low elongation, which is comparable with Cu Ag and Cu Ag Zr alloys processed by conventional cold drawing [29,44]. In addition, since the microstructural scale of the disks after 20 revolutions is even smaller than 20 nm, the high tensile strength of 1420 MPa may correspond to the limiting strength due to grain refinement. In fact, in a two-phase Cu Nb alloy after heavy drawing, a transition of the fracture mode occurred whereby there was a cup cone fracture at the lower draw ratios of and a tendency towards a shear fracture at a draw ratio of 8.2 [32]. 4. Discussion 4.1. Microstructural evolution characteristics during HPT Microstructural evolution process The primary results for Cu 28 wt.% Ag revealed an inhomogeneous evolution process in the form of vortices during HPT [24]. Furthermore, the present study shows clear evidence of microstructural evolution on the transverse section of the HPT disks by observing the decorative eutectic, as shown in Figs The strong strain hardening ability and the late saturation of microhardness indicate a significant difference between two-phase Cu Ag alloys and single-phase alloys. For single-phase alloys, the microstructural homogeneity is generally obtained by a dislocation-based or twinning process. A recent investigation examined the microstructural homogeneity of various single-phase alloys with stacking fault energies (SFE) ranging from low values to high values when subjected to HPT [38]. It was proposed that the apparent dichotomy between high, low and intermediate SFE materials lies in the nature of the dominant deformation process in achieving grain refinement. Consequently, either a high recovery rate in high-sfe materials or the occurrence of twins, stacking faults and shear bands in low-sfe materials can accelerate the evolution towards a reasonably homogeneous distribution of nanostructured grains. However, it is more challenging in materials where the rate of recovery is relatively slow and, in addition, the formation of twins is difficult. In the present study, in contrast to the single-phase alloys, the microstructural evolution is more difficult because of the presence of the second phase. Based on the results shown in Figs. 4 6, the microstructural evolution process is schematically illustrated in Fig. 13. Fig. 13a shows the as-received microstructure, where the Cu matrix is indicated in gray, and the eutectic is composed of gray Cu phase and white Ag phase. When the shear strain is small, both the eutectic and the Cu matrix are sheared. As a result, the Cu and Ag phases of the eutectic become elongated, and a banded structure forms inside the Cu matrix, as shown in Fig. 13b. A further increase in the shear strain leads to significant refinement, as shown in Fig. 13c, where the eutectic spacing is reduced and the internal band width of the Cu matrix is further narrowed. Note that the band width may be smaller than 50 nm; in contrast, the internal Cu and Ag phases inside the eutectic are further sheared and narrowed, and most of the eutectic is fully refined and saturated where small domains of several nanometers in size are produced. When the shear

11 Y.Z. Tian et al. / Acta Materialia 59 (2011) For Cu Ag alloys, mechanical alloying occurred during conventional cold drawing [45], and there are many observations of substantial mixing of Cu and Ag within the shear bands [10]. It is interesting to note that HPT even leads to the formation of a homogeneous supersaturated solid solution of 12 at.% Fe in Cu [46]. For the present study, it is possible that mechanical alloying may also occur in the microstructural evolution of the Cu Ag alloy during HPT processing, owing to the very intense shearing strain. Fig. 13. Schematic illustration of the microstructural evolution of the Cu Ag alloy under different strain levels. strain is sufficiently large (for example, at a distance of 4.5 mm in the disk after 20 revolutions, as shown in Fig. 6e and f), the width of the eutectic may be reduced and nearly all the eutectic is fully refined. In addition, the eutectic spacing may be decreased to a limit where only one internal band is observed in the Cu matrix, indicating that the substructure has diminished. For this condition, dislocation activity is extremely difficult, and further plastic deformation can be accommodated by inhomogeneous deformation in the form of shear banding, as shown schematically in Fig. 13d. Within the shear band, both the eutectic and the Cu matrix are fully mixed. It was also reported for a Cu Ag eutectic alloy processed by ARB that the inter-lamellar spacing of Cu and Ag phases in the eutectic became more or less saturated after only two cycles of ARB, because heterogeneous deformation by shear banding was the dominant deformation mode when the lamellar distance was very small [10]. Note that the band width of the Cu matrix is significantly smaller than for pure Cu, where the recovery and recrystallization processes lead to a constant grain size and saturation strength. However, in the composites, the continual deformation recovery recrystallization cycle allows further microstructural refinement within the Cu matrix [32]. In addition to the present work on microstructural evolution, there is also a body of research showing the inter-diffusion of elements in two-phase materials during HPT processing, even if the phases are immiscible [10,45 47] Microstructures in the centers of the disks It should be noted that the conventional relationship given in Eq. (1) predicts an absence of any strain at the center of the HPT disk, and yet the present observations, and Fig. 3c 1 e 1, demonstrate the occurrence of a significant plastic straining after rotations through 5, 10 and 20 revolutions. In addition, the microhardness near the center also indicates the occurrence of plastic strain, as shown in Fig. 8. It is also of interest to note that the microhardness in various materials processed by HPT increases in the centers of the disks by comparison with their unprocessed counterparts [4,11 15,31,36 41]. Since the magnitude of the shear strain scales with the radial position, there will be a strain gradient throughout the disk, and geometrically necessary dislocations are therefore required to accommodate the inhomogeneous shear strain along the disk radius [48]. Accordingly, recent work using two separate variants of strain gradient plasticity modeling of HPT shows that plastic strain accumulates initially at a very high rate at the specimen rim while having a relatively low value in a broad area around the center; but on further straining, this difference is gradually reduced as the rate of strain accumulation in the center exceeds that at the rim, thereby leading to a reasonably uniform strain distribution after 5 revolutions [17]. The present results collectively provide further evidence for straining near the disk center, and it is suggested accordingly that the occurrence of strain hardening at the center of the disk may be associated with this strain gradient [24,49]. In addition, it will be assisted by the inhomogeneous nature of the deformation process Strain hardening behavior and strengthening mechanism When single-phase metals and alloys are processed by HPT, the grain size is greatly refined, and the materials experience strain hardening. Detailed experimental results show that these materials can be classified into two types: for an extremely high-sfe material (high-purity Al [12,31,50,51]), the micorhardness is high at the center and low at the peripheral region of the disk, and these high values are gradually reduced with increasing numbers of revolutions; for intermediate- and low-sfe materials (commercial Al and Al alloys [13 15,36], Cu and Cu alloys [37,38,52]), the microhardness is low at the center and high at the outer region of the disk. These different evolution processes are attributed to the rate of recovery, which is

12 2794 Y.Z. Tian et al. / Acta Materialia 59 (2011) dominated by the value of the SFE. When the microhardness values of each material are plotted as a function of the equivalent strain, a saturation microhardness is obtained where the critical equivalent strain for saturation generally does not exceed 50 [19,20,38 40,42,50], thereby indicating a fast rate of homogenization. A recent study on a two-phase Zn Al alloy processed by HPT revealed a different condition [23,53]: the microhardness decreased continuously with increasing equivalent strain and ultimately leveled off, where this unique evolutionary process was due to a significant reduction during processing in the distribution of Zn precipitates which were visible in the Alrich grains in the unprocessed condition. In the present study, Fig. 8 shows that the microhardness values across the diameter of the disk continuously increase with increasing numbers of revolutions, and a saturation microhardness is found only when the equivalent strain is as high as 330, as demonstrated in Fig. 9. This threshold value is much higher than for single-phase alloys [19,20,23,38 40,42,50,51,53]. It is suggested that this is related to the two-phase nature of the structure and the corresponding strengthening mechanisms. The microstructural evolution process includes several different stages, and it is proposed that the strengthening mechanism also evolves. There has been much debate over the strengthening mechanism for the high strengths of Cu X (Ag [27,29], Nb[32], Cr[33], Fe[34]) alloys. A modified rule of mixtures is generally applied to explain the mechanical properties of Cu Ag alloys processed by cold drawing [27,29]. In the present study, the two components of the Cu matrix and the eutectic are considered according to the evolution process as shown in Figs. 2 6, so that the strength of the Cu Ag alloy can be calculated by r Cu Ag ¼ð1 f E Þr M þ f E r E ð3þ where f E is the volume fraction of the eutectic component, and r E and r M are the strengths of the eutectic and the Cu matrix, respectively. When the shear strain is small so that the microstructure is coarse, the strength of both components can be calculated from a Hall Petch type equation, r M ¼ r Cu 0 þ k 1 r 1=2 Cu ð4þ r E ¼ r 0 þ k 2 r 1=2 e ð5þ where r 0 and r Cu 0 are the intrinsic friction stresses of the eutectic and the Cu matrix, respectively. r Cu is the band width of the Cu matrix, r e is the spacing between the Cu or Ag phases inside the eutectic, and k 1 and k 2 are two strengthening coefficients of the Cu Ag alloy. With increasing shear strain, the eutectic spacing and the internal band width are narrowed, and the eutectic is significantly refined to saturation because the eutectic is much finer in the as-received condition and thus experiences a faster evolution to saturation. In that case, the strength of the matrix is again calculated using Eq. (4), but the strength of the eutectic becomes r E ¼ r S ð6þ where r S is the saturation strength of the eutectic component. When the shear strain is significantly large and the eutectic spacing is reduced to such a scale that the substructure is diminished, it becomes extremely difficult to propagate dislocations in the matrix. A model based on the critical stress necessary to propagate dislocations in the matrix between cementite lamellae has been proposed to explain the strengthening mechanism in pearlite [54]. According to this model, the critical stress necessary to propagate dislocations in the matrix is given by [54,55] r M ¼ r Cu 0 þ MAGb 2pk ln k ð7þ b where M is the Taylor factor, A is 1.21 (for mixed dislocations), G is the shear modulus, b is the Burgers vector, and k is the eutectic spacing. These strengthening mechanisms are illustrated schematically in Fig. 14. Fig. 14. The strengthening mechanisms in the Cu Ag alloy corresponding to different microstructure scales.

13 Y.Z. Tian et al. / Acta Materialia 59 (2011) Based on the strengthening mechanisms as delineated above, it is therefore possible to interpret the strain hardening behavior of the Cu Ag alloy as shown in Fig. 9, where this consists of two strong strain hardening stages. Generally, when the equivalent strain is less than 15, the first stage is realized by rapid accumulation of dislocations and the formation of substructures. Thereafter, the continuous decrease in the eutectic spacing and the internal band width of the Cu matrix, together with the further refinement of the eutectic, are responsible for the second stage of strain hardening. It is noted that deformation twins were also observed in the samples strained to different numbers of revolutions, but it appears that they play only a minor role in strengthening owing to their small volume fraction Fracture mode On the macroscale, the Considére criterion governs the onset of localized deformation in 6 r _e where r and e are true stress and strain, respectively. This criterion also predicts that as-processed UFG/NC alloys will lose their strain hardening ability after the onset of yielding. Accordingly, the HPT-processed Cu Ag alloy in this study rarely retains any uniform elongation, as shown in Fig. 11. This is prevalent in Cu Ag and Cu Ag Zr alloys deformed to high strains by conventional cold drawing [30,44]. Further inspection shows that the specimens processed through 5, 10 and 20 revolutions ruptured by shearing at different shear angles of 50, 47 and 54, respectively, as revealed in Fig. 12. In a recent report on iron processed by HPT, the specimens macroscopically failed through shear fracture at an angle of 45 [56], and two nanostructured pure Ni samples with multi- and bi-modal grain size distributions displayed a shear fracture mode with an angle of 50 [57]. All these phenomena exhibit similarities to bulk metallic glasses (BMG) during tensile testing [58] where the specimens generally fail in a shear fracture mode with shear angles >45. A unified tensile fracture criterion was proposed to describe the tensile fracture behavior of various BMG as well as a variety of other brittle materials [59]. It is suggested that the tensile fracture angle of each material depends on the ratio of a = s 0 /r 0, where s 0 and r 0 are the critical shear fracture stress and the normal cleavage fracture stress of the material, respectively. For the present Cu Ag alloy processed by HPT for 5, 10, and 20 revolutions, with refining of the grain size the slip stress in the interior of grains increases so that necking becomes more difficult, as observed in Fig. 12. Meanwhile, the critical shear stress s 0 along the localized shear bands is also enhanced with grain refinement. Therefore, the Cu Ag alloy becomes more brittle and displays an obvious shear failure mode under tensile loading. In this case, the shear fracture is attributed to the nanostructures, which lead to an increased ratio of a = s 0 /r 0 in contrast to the as-received counterpart [59,60]. The increased ratio of a = s 0 /r 0 means that the sample fails in a shear mode with a shear angle >45, as reported in some other UFG or NC materials [57,61]. Although the HPT-processed Cu Ag alloy fractured by shearing and possesses little elongation, detailed observations on the fracture surface in Fig. 12 reveal the ductile features as numerous dimples that are visible throughout the fracture surface. This is similar to the features of veins and cores observed on the tensile fracture surfaces of BMG [62]. 5. Summary and conclusions (1) A Cu 28 wt.% Ag alloy was processed by HPT through up to 20 revolutions, and the microstructural evolution and mechanical properties were investigated. A nanostructured Cu Ag microcomposite and a high tensile strength of 1420 MPa were achieved after 20 revolutions of HPT. (2) The eutectic component shows a faster evolution process than the Cu matrix. Banded structures are found in the Cu matrix, and both the eutectic spacing and the internal band width decrease with increasing shear strain. The substructure diminishes in the Cu matrix after 20 revolutions. (3) When the micohardness values are plotted against the equivalent strain, there is a saturation in the microhardness value at a strain of 330. It is shown that multiple strengthening mechanisms correspond to this two-stage strain hardening behavior. (4) There is a transition in the fracture mode from necking to shearing with increasing numbers of HPT resolutions. The measured fracture shear angles are >45, indicating that shear fracture dominates the failure mode due to the considerable grain refinement in the Cu Ag alloy. Acknowledgements The authors are grateful to Prof. J.D. Embury, Prof. J.T.M. De Hosson and Prof. Y. Brechet for discussions during their visit to IMR. This work was financially supported by the National Natural Science Foundation of China (NSFC) under Grant Nos , and , the National Basic Research Program of China under Grant No. 2010CB and the Royal Society of the UK under International Joint Project No. JP References [1] Valiev RZ, Islamgaliev RK, Alexandrov IV. Prog Mater Sci 2000;45:103. [2] Zhu YT, Lowe TC, Langdon TG. Scripta Mater 2004;51:825. [3] Valiev RZ, Estrin Y, Horita Z, Langdon TG, Zehetbauer MJ, Zhu YT. JOM 2006;58(4):33.

14 2796 Y.Z. Tian et al. / Acta Materialia 59 (2011) [4] Zhilyaev AP, Langdon TG. Prog Mater Sci 2008;53:893. [5] Furukawa M, Horita Z, Nemoto M, Langdon TG. J Mater Sci 2001;36:2835. [6] Valiev RZ, Langdon TG. Prog Mater Sci 2006;51:881. [7] Hong CS, Tao NR, Huang X, Lu K. Acta Mater 2010;58:3103. [8] Li YS, Zhang Y, Tao NR, Lu K. Acta Mater 2009;57:761. [9] Saito Y, Utsunomiya H, Tsuji N, Sakai T. Acta Mater 1999;47:579. [10] Ohsaki S, Kato S, Tsuji N, Ohkubo T, Hono K. Acta Mater 2007;55:2885. [11] Zhilyaev AP, Nurislamova GV, Kim BK, Baró MD, Szpunar JA, Langdon TG. Acta Mater 2003;51:753. [12] Xu C, Horita Z, Langdon TG. Acta Mater 2007;55:203. [13] Zhilyaev AP, McNelley TR, Langdon TG. J Mater Sci 2007;42:1517. [14] Xu C, Horita Z, Langdon TG. Acta Mater 2008;56:5168. [15] Xu C, Langdon TG. Mater Sci Eng A 2009;503:71. [16] Xu C, Furukawa M, Horita Z, Langdon TG. Mater Sci Eng A 2005;398:66. [17] Estrin Y, Molotnikov A, Davies CHJ, Lapovok R. J Mech Phys Solids 2008;56:1186. [18] Valiev RZ, Ivanisenko Yu V, Rauch EF, Baudelet B. Acta Mater 1996;44:4705. [19] Wetscher F, Vorhauer A, Stock R, Pippan R. Mater Sci Eng A 2004; :809. [20] Vorhauer A, Pippan R. Scripta Mater 2004;51:921. [21] Cao Y, Wang YB, Alhajeri SN, Liao XZ, Zheng WL, Ringer SP, et al. J Mater Sci 2010;45:765. [22] Cao Y, Kawasaki M, Wang YB, Alhajeri SN, Liao XZ, Zheng WL, et al. J Mater Sci 2010;45:4545. [23] Kawasaki M, Ahn B, Langdon TG. Acta Mater 2010;58:919. [24] Tian YZ, An XH, Wu SD, Zhang ZF, Figueiredo RB, Gao N, et al. Scripta Mater 2010;63:65. [25] Freudenberger J, Grünberger W, Botcharova E, Gaganov A, Schultz L. Adv Eng Mater 2002;4:677. [26] Embury JD, Fisher RM. Acta Metall 1966;14:147. [27] Benghalem A, Morris DG. Acta Mater 1997;45:397. [28] Tian YZ, Zhang ZF, Wang ZG. Philos Mag 2009;89:1715. [29] Han K, Vasquez AA, Xin Y, Kalu PN. Acta Mater 2003;51:767. [30] Tian YZ, Zhang ZF. Mater Sci Eng A 2009;508:209. [31] Kawasaki M, Langdon TG. Mater Sci Eng A 2008;498:341. [32] Spitzig WA, Pelton AR, Laabs FC. Acta Metall 1987;35:2427. [33] Adachi K, Tsubokawa S, Takeuchi T, Suzuki HG. J Jpn Inst Met 1997;61:397. [34] Funkenbusch PD, Courtney TH. Scripta Metall 1981;15:1349. [35] Tian YZ, Duan QQ, Yang HJ, Zou HF, Yang G, Wu SD, et al. Metall Mater Trans A 2010;41:2290. [36] Loucif A, Figueiredo RB, Baudin T, Brisset F, Langdon TG. Mater Sci Eng A 2010;527:4864. [37] Horita Z, Langdon TG. Mater Sci Eng A 2005; :422. [38] An XH, Wu SD, Zhang ZF, Figueiredo RB, Gao N, Langdon TG. Scripta Mater 2010;63:560. [39] Edalati K, Horita Z. Mater Trans 2010;51:1051. [40] Zhang HW, Huang X, Hansen N. Acta Mater 2008;56:5451. [41] Bayramoglu S, Gür CH, Alexandrov IV, Abramova MM. Mater Sci Eng A 2010;527:927. [42] Harai Y, Edalati K, Horita Z, Langdon TG. Acta Mater 2009;57:1147. [43] Sakai Y, Schneider-Muntau H-J. Acta Mater 1997;45:1017. [44] Freudenberger J, Lyubimova J, Gaganov A, Witte H, Hickman AL, Jones H, et al. Mater Sci Eng A 2010;527:2004. [45] Raabe D, Ohsaki S, Hono K. Acta Mater 2009;57:5245. [46] Quelennec X, Menand A, Le Breton JM, Pippan R, Sauvage X. Philos Mag 2010;90:1179. [47] Sauvage X, Genevois C, Da Costa G, Pantsyrny V. Scripta Mater 2009;61:660. [48] Ashby MF. Philos Mag 1970;21:399. [49] Todaka Y, Umemoto M, Yin J, Liu Z, Tsuchiya K. Mater Sci Eng A 2007;462:264. [50] Ito Y, Horita Z. Mater Sci Eng A 2009;503:32. [51] Xu C, Horita Z, Langdon TG. Mater Trans 2010;51:2. [52] Balogh L, Ungár T, Zhao Y, Zhu YT, Horita Z, Xu C, et al. Acta Mater 2008;56:809. [53] Kawasaki M, Ahn B, Langdon TG. Mater Sci Eng A 2010;527:7008. [54] Sevillano JG. Strength of metals and alloys. In: Haasen P, Gerold V, Kowtorz G, editors. Proc. ICSMA 5. Oxford: Pergamon Press; p [55] Verhoeven JD, Chumbley LS, Laabs FC, Spitzig WA. Acta Metall Mater 1991;39:2825. [56] Hohenwarter A, Pippan R. Mater Sci Eng A 2010;527:2649. [57] Zhao YH, Topping T, Bingert JF, Thornton JJ, Dangelewicz AM, Li Y, et al. Adv Mater 2008;20:3028. [58] Zhang ZF, He G, Eckert J, Schultz L. Phys Rev Lett 2003;91: [59] Zhang ZF, Eckert J. Phys Rev Lett 2005;94: [60] Zhang ZF, Eckert J. Adv Eng Mater 2007;9:143. [61] Fang DR, Duan QQ, Zhao NQ, Li JJ, Wu SD, Zhang ZF. Mater Sci Eng A 2007;459:137. [62] Zhang ZF, Eckert J, Schultz L. Acta Mater 2003;51:1167.