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Microstructural evolution of copper clad steel bimetallic wire
Materials Science and Engineering A 528 (2011) 2974–2981
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Microstructural evolution of copper clad steel bimetallic wire T.T. Sasaki a , M. Barkey b , G.B. Thompson a,∗ , Y. Syarif c , D. Fox c a
The University of Alabama, Department of Metallurgical & Materials Engineering, Tuscaloosa, AL 35487-0202, United States The University of Alabama, Department of Aerospace Engineering & Mechanics, Tuscaloosa, AL 35487-0202, United States c Fushi Copperweld, Inc., Fayetteville, TN 37334-7249, United States b
a r t i c l e
i n f o
Article history: Received 7 September 2010 Accepted 8 December 2010 Available online 14 December 2010 Keywords: Copper clad steel Drawing Texture Finite element modeling
a b s t r a c t We investigated the microstructure of two different bimetallic wires of Copper Clad Low Carbon Steel Wire (LCSW), which had a 1006 steel core, and Copper Clad High Carbon Steel Wire (HCSW), which had a 1055 steel core. The HCSW generally showed higher hardness than LCSW because of the pearlitic grain structure. A low temperature annealing at 720 ◦ C to the drawn HCSW caused a significant reduction of hardness, which was as low as that of an annealed LCSW. In general, both LCSW and HCSW showed strong global textured features after drawing, with the steel having a strong 1 1 0 fiber texture and the copper having a 1 1 1 –1 1 2 deformation direction. At the interface, a grain size discrepancy at the steel–copper interface was observed. Post-drawing, the LCSW copper grains exhibited refined grain sizes near the interface and has been explained in terms of shear strain gradient. The HCSW did not exhibit this copper grain size distribution but did exhibit a coarsening of the steel grains near the interface after a subsequent 720 ◦ C heat treatment. This is attributed to the large localized stress concentration at the perimeter of the steel region during the drawing process. The strain induced regions at the steel–copper interface have been simulated by finite element modeling. These grain size discrepancies caused the smooth variation in nanohardness across the interface. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Multi-component wires offer the ability to blend the properties of different metals into a single wire form to achieve improved performance such as electrical conduction or mechanical strength. In addition, these composite wires can assist in cost reduction by using less expensive wires as support or filler material in place of higher cost, single elemental wires [1–15]. In particular, copper clad steel (CCS) bimetallic wires can be a low cost alternative to single copper wires and has applications as the center conductor in high frequency coaxial cables, power transmission, and tracer wires for buried ceramics [10–15]. The copper sheath provides the low electrical resistance while the center core steel provides high structural strength, which can be tuned by the carbon content of the steel wire and heat treatment. To date, a number of studies on the structure–property relationship have been done for either Cu or steel single wires. In elemental metal drawn wire, the microstructure generally consists of filamentary grains elongated along the drawing direction with textured features [16]. In steel wires, the grains normally exhibit a strong
∗ Corresponding author at: The University of Alabama, Department of Metallurgical & Materials Engineering, 301 7th Avenue, 116 Houser Building, Tuscaloosa, AL 35487-0202, United States. Tel.: +1 205 348 1589; fax: +1 205 348 2164. E-mail address: [email protected] (G.B. Thompson). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.12.032
1 1 0 fiber texture [17–20]. In some specific processing routes, the deviation from the 1 1 0 fiber texture can be observed at the perimeter region for the wires [21,22]; similar variations in texture have been observed on the surface of rolled sheets [23–25]. On the other hand, Cu grains exhibit a major 1 1 1 and a minor 1 0 0 fiber texture in the as-drawn wire condition [16,26,27]. Texture deviations at the perimeter region and texture changes by annealing of the wire have also been reported [27–29]. There have been reports on CCS wires, sheets and extruded bars, focusing on the processing coupled with finite element modeling (FEM) [10–15]. To date, however, there has been limited work addressing the microstructure coupled to the simultaneous co-deformation behavior that occurs in the CCS bimetallic wires. Since steel and copper have different mechanical properties, these metals at the interface may deform differently to stay “connected” to each other during drawing. In the present study, we have selected two different CCS wires using either low- and high-carbon steel as the core rod and compared the microstructure evolution during the drawing process in terms of texture formation, grain structure, and local microstructure features at the steel–copper interface.
2. Experimental procedure Two CCS wires, one with a low carbon 1006 steel and the other with a high carbon 1055 steel have been clad to an oxygen-free cop-
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Table 1 Processing route of CCS bimetallic wires of (a) LCSW and (b) HCSW. The ‘D’ designates drawing stage and ‘HT’ designates the heat treatment stage. Diameter (mm)
Cu thickness (mm)
Drawing ◦
Heat treatment
Temp. ( C) (a) As-Clad D-1 HT-1 D-2 HT-2 Fine drawn
9.5 4.1 4.1 2.0 2.0 1.0
0.95 0.41 0.41 0.20 0.20 0.10
(b) As-Clad D-1 D-2 HT-1 Fine drawn HT-2
9.5 4.6 2.9 2.9 1.2 1.2
0.28 0.14 0.09 0.09 0.04 0.04
Rate (m/s)
Strain
55
7.1
1.7
55
10.2
1.4
55
20.3
1.4
55 55
7.1 10.2
1.5 2.4
55
15.2
1.8
per (99.95%) metal strip. The copper-strip was continuously formed around a moving steel rod during the cladding operation. The clad wire had an initial wire diameter of 9.5 mm with the steel core diameter being 7.62 mm for 1006 and 9.01 mm for 1055. Hereafter, the CCS wires are described as Low Carbon Steel Wire (LCSW) for the 1006 core steel and High Carbon Steel Wire (HCSW) for the 1055 core steel. Table 1(a) and (b) summarizes the processing conditions of the LCSW and HCSW. The wire diameter was reduced from 9.5 to 1.0 mm by a combination of three major drawing steps (D-1, D-2 and fine drawing) and intermediate heat treatments (HT1 and HT-2). The relative reduction of each metal, copper or steel, is consistent during drawing, e.g. both metals reduce at the same reduction fraction relative to their initial diameter, or, alternatively, the volume is conserved and one metal does not reduce any more than the other to compensate for discrepancies in mechanical property differences. The strain introduced during the drawing process was calculated by = ln(A0 /A), where A0 and A is the area before and after drawing. The bulk hardness values of the wires were measured by a Vickers hardness tester. The variation in the hardness across the copper–steel interface was measured by nanoindentation using a Hysitron TI900 TriboIndenter. Thirty indents were made with a spacing of 1 m between indents under a 4000 N load. The hardness measurements for both tests were performed on the normal plane to the drawing direction. Microstructure characterization was carried out using a field emission gun-scanning electron microscope (SEM), JEOL JSM7000F, equipped with an Oxford electron backscattered diffraction (EBSD) system and INCA software. The texture was measured by EBSD and is presented in the form of inverse pole figure map (IPF map) and inverse pole figure (IPF). The EBSD analyses were carried out from both normal and parallel planes to the drawing direction. The finite element modeling (FEM) was conducted using ABAQUS Explicit 6.9-1. An axi-symmetric model was constructed for the bimetallic wire using 4-noded axi-symmetric reduced integration elements. Table 2 summarizes the initial size of the wire, steel diameter and copper thickness for LCSW and HCSW. The model was made of elements tabulated in Table 2. For LCSW, 48 elements of steel and 12 elements of copper were used in the radial direction, respectively. For HCSW, 56 elements of steel and 4 elements of copper were used in the radial direction, respectively. The difference in number of elements was done because of the experimental differences in steel and copper diameters and thicknesses, respectively, between the LCSW and HCSW. The boundary conditions of the wire were set to an initial velocity of the wire at 15 m/s. The pressure at the lower end of the wire was 150 and 200 MPa and a pressure on the upper end of the wire was 10 and 50 MPa for
Temp. (◦ C)
Time (min)
950
15
720
480
950
21
720
480
Table 2 Initial size of wire, steel diameter and copper thickness for LCSW and HCSW used for FEM modeling.
Wire Total Diameter (mm) Steel diameter (mm) Copper thickness (mm) Length (mm) Element Radial direction (mm) Axial direction (mm)
LCSW
HCSW
2.052 1.642 0.2052 26.3
2.906 2.724 0.0908 26.3
0.0171 0.0496
0.0227 0.0496
the LCSW and HCSW, respectively. The pressure was applied in the vertical direction. Radial displacements were restrained for nodes on the axi-symmetry line. The die was modeled using 2-noded axi-symmetric rigid elements. The bearing length of the die was taken as 1 mm and the bell angle and back relief angle were 60◦ which mimics the experimental set-up. The interaction between the surface of the die and the contact surface of the wire was modeled with a coefficient of friction of 0.01. The die was initially offset from the surface of the wire by 0.0224 and 0.0455 mm and moved radially into the wire by a distance of 0.239 and 0.2 mm for LCSW and HCSW, respectively. After the wire was moved through the die, the reduction of area was 20%. The step time of the analysis was 0.003 s and the time that the die was moved into place was 0.0001 s. 3. Results Fig. 1 shows variations in Vickers hardness for the LCSW and HCSW as a function of the wire diameter. Note that the open and closed plots show the Vickers hardness values in the as-drawn and heat-treated conditions, respectively. In both wires, the hardness of the steel and copper increases after drawing because of work hardening and is recovered, i.e. lower hardness, partially or fully upon the heat treatments. The hardness values of the heat treated samples are slightly lower than the As-Clad conditions for both metals indicating that the sample is slightly work hardened during the cladding process. After each draw, the hardness of the LCSW was approximately the same, which is consistent that the strain was approximately equivalent among each draw, Table 1. The HCSW increased the hardness between the D-1 and D-2 draws. This is attributed to the increase in the strain between these two draws. Interestingly, the fine-drawn condition of the HCSW had a higher hardness value than the D-2 condition, even though the lower strain was introduced compared to the HCSW in the D-2
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Fig. 1. Variation in Vickers hardness as a function of wire diameter in (a) LCSW and (b) HCSW. Note that open and closed circles show the Vickers hardness in as-drawn condition and the one after heat treatment, respectively.
condition, Table 1 (b). The LCSW recovers with increasing temperature. In contrast, the HCSW exhibited a higher amount of recovery during a low temperature heat treatment, HT-2, of D-2, as seen in Fig. 1. Fig. 2 shows typical IPF maps obtained from normal and parallel planes to the drawing direction of the LCSW in the As-Clad, HT-2 and fine-drawn conditions. As shown in Fig. 2(a) and (b), both ferrite and copper grains are equiaxed in the As-Clad condition. The ferrite grains have an average size of approximately 8.7 m.
Upon HT-2 annealing, as shown in Fig. 2(c) and (d), the copper grains are recrystallized and exhibit annealing twins as indicated by black arrows. Their average grain size is approximately 153 m. The ferrite grains in the HT-2 condition are equiaxed and are slightly refined to 6.1 m as compared to those in the As-Clad condition. In the Fine drawn condition, the filamentary grain microstructure from the drawing process is clearly present, Fig. 2(e) and (f). The drawing process causes a loss of the distinct equiaxed grains. A similar filamentary microstructure, as in Fig. 2(e) and (f), was observed for the D-1 and D-2 LCSW wires. Fig. 3 shows IPFs obtained from the steel region at each processing step for the LCSW. Note that the IPFs are obtained from normal and parallel planes to the drawing direction. Fig. 3(a) confirms ferrite grains do not show any obvious textured feature in the As-Clad condition. During the D-1 drawing step, a strong 1 1 0 fiber texture is developed for the ferrite, Fig. 3(b). During the HT-1 heat treatment, the 1 1 0 fiber texture becomes diffused as seen in Fig. 3(c). But the 1 1 0 direction of the ferrite grains, Fig. 3(c), are qualitatively more aligned to the drawing direction as compared to the As-Clad sample, Fig. 3(a). The D-2 draw again recovers the strong 1 1 0 fiber texture, Fig. 3(d). This strong fiber texture is subsequently lost post HT-2, Fig. 3(e). Upon the fine drawing, the 1 1 0 fiber texture is recovered, Fig. 3(f), but slightly diffused compared to the one in D-1 and D-2 conditions. The quantification of the texture feature is difficult for the Cu grains in the As-Clad, HT-1 and HT-2 conditions since the large grain sizes limited the counting statistics over the field of view studied, as seen in Fig. 2. However, obvious deformation textures can be seen in the drawn wires. The IPF maps, Fig. 4(a) and (b) shows that the copper grains for the LCSW in the D-1 condition developed a tendency for the 1 1 1 –1 1 2 directions, i.e. the distribution along the 1 1 1 –1 0 0 symmetry boundary of the unit angle, aligning with the drawing direction. In the D-2 condition, the majority of the grains are present with their major 1 1 2 direction parallel to the drawing direction and the remaining grains indicating a minor 1 0 0 alignment, as seen in Fig. 4(c) and (d). The Fine drawn texture, Fig. 4(e) and (f), shows the copper grains also form diffuse texture along 1 1 1 –1 0 0 symmetry boundary of the unit triangle. Compared to copper grains in the D-1 and D-2 conditions, these copper grains in the fine-drawn condition showed slightly stronger 1 0 0 fiber texture.
Fig. 2. Inverse pole figure (IPF) maps obtained from copper–steel interface in LCSW. (a) and (b) are those obtained from As-Clad wire taken from normal (⊥) and parallel () planes to drawing direction. (c) and (d) are heat treated (HT-2) wire before fine drawing taken from normal and parallel direction to drawing direction. (e) and (f) are from fine drawn wire taken from normal and parallel direction to drawing direction, respectively.
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Fig. 3. Inverse pole figure obtained from the steel region in LCSW. (a) is obtained from normal (⊥) and parallel () planes to the drawing direction in As-Clad, (b) is obtained from normal and parallel planes to the drawing direction in D-1, and (c) is obtained from normal and parallel planes to the drawing direction in HT-1, (d) is obtained from normal and parallel planes to the drawing direction in D-2, and (e) is obtained from normal and parallel planes to the drawing direction in HT-2, (f) is obtained from normal and parallel planes to the drawing direction in fine drawn.
Fig. 5 is backscattered electron SEM micrographs of the HCSWs in HT-1, fine-drawn and HT-2 conditions observed from the parallel plane with respect to the drawing direction. In Fig. 5(a), the steel region in the HT-1 consists of equiaxed ferrite and pearlite grains. Upon Fine drawing, the ferrite grains are elongated along the drawing direction, Fig. 5(b), but the copper grains are equiaxed compared to the steel grains. The subsequent heat treatment of HT-2, Fig. 5(c), the ferrite grains near the Cu interface are coarser than those further away from the interface. Unlike the HT-1 condition, the HT-2 microstructure does not show the pearlite grain
Fig. 4. Inverse pole figure obtained from the Cu region in LCSW. (a) is obtained from normal (⊥) and parallel () planes to the drawing direction in D-1, (b) is obtained from normal and parallel planes to the drawing direction in D-2, and (c) is obtained from normal and parallel planes to the drawing direction in fine drawn.
structure. The Fe3 C phase, the spotty dark features, is present as spherical particles mainly along the grain boundaries. Fig. 6 shows typical inverse pole figure (IPF) maps obtained from normal and parallel planes to the drawing direction of HCSW in the As-Clad, Fine drawn and HT-2 conditions. As shown in Fig. 6(a) and (b), both steel and copper regions in the As-Clad sample consists of equiaxed grains similar to those reported above for the LCSW in AsClad condition, Fig. 2(a) and (b). The ferrite grains and copper grains have an average size of approximately 5.7 m and 67 m, respectively. In the Fine drawn condition, Fig. 6(c) and (d), both ferrite and copper grains have the typical filamentary grain structure formed from the drawing process. The filamentary copper grains consist of the equiaxed grains with low angle boundaries. Similar fine filamentary grain structures were observed for the other D-1 and D-2 drawn wires for the HCSW. Upon HT-2 heat treatment, Fig. 6(e) and (f), the steel grains recover the equiaxed grain structure. As noted previously, the grains near the copper interface are coarser and elongated along the drawing direction; this will be developed further in Section 4. Fig. 7 is the inverse pole figures (IPFs) obtained from the steel region in the HCSW at each processing step. Note that these IPFs are obtained from normal and parallel planes to the drawing direction. The steel region does not show any obvious textured features in the As-Clad condition, Fig. 7(a). The D-1 drawing substantially develops the 1 1 0 fiber texture, Fig. 7(b). As shown in Fig. 7(c), the 1 1 0 fiber texture becomes more intense with the subsequent D-2 draw. During the HT-1 heat treatment, the drawn 1 1 0 fiber texture becomes more diffused, Fig. 7(d), although the 1 1 0 fiber texture is still observable. The strong 1 1 0 fiber texture is recovered with the Fine drawing as shown in Fig. 7(e). The 1 1 0 strong texture in the HT-2, Fig. 7(f), is still retained even after heat treatment. Similar to the LCSW, the large copper grains in the heat treated HCSW limited the counting statistics in the EBSD scans as seen in Fig. 8. Again, the IPFs for copper grains in Fig. 8 are obtained from normal and parallel planes to the drawing direction. In the D-1 condition, the majority of the grains are present with their 1 1 1 direction parallel to the drawing direction, Fig. 8(a). The copper grains in D-2 condition develop a relatively diffused texture as compared to those in the D-1 condition. A slightly higher density of the 1 1 w textures parallel to the drawing direction developed for D-2 condition, Fig. 8(b). Similar to the LCSW, a tendency for the 1 1 1 –1 1 2 (distributed along the 1 1 1 –1 0 0 symmetry boundary of the unit angle) is aligned with the drawing direction of the 1 1 1 fiber texture. In the Fine drawn condition, the copper grains form a diffuse texture for the major 1 1 1 plus the minor 1 0 0 fiber orientations, Fig. 8(c). For the Fine drawn condition, the copper grains in the HCSW show stronger 1 0 0 fiber texture compared to the ones in the D-1 and D-2 conditions similar to those in the LCSW.
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Fig. 5. Backscattered electron SEM images of the HCSWs in (a) HT-1, (b) fine drawn and (c) HT-2 conditions. Note that these images are taken from the parallel plane to the drawing direction, and the drawing direction is horizontal direction.
Fig. 6. Inverse pole figure (IPF) maps obtained from Fe/Cu interface in HCSW. (a) and (b) are those obtained from As-Clad wire taken from normal (⊥) and parallel () planes to drawing direction. (c) and (d) are from the sample in fine drawn wire taken from normal and parallel planes to drawing direction, respectively. (e) and (f) are obtained from HT-2 heat treated wire after fine drawing taken from normal and parallel planes to drawing direction.
Fig. 9(a) shows the variation in nanohardness across the interface for the LCSW in the As-Clad and Fine drawn conditions. In the As-Clad condition, the average nanohardness values are 2.4 and 1.4 GPa in the steel and copper regions, respectively. A discontinuous change at the steel–copper interface exists in the As-Clad wire. In the Fine drawn condition, the steel and copper regions have a
higher nanohardness value of 3.9 and 2.1 GPa, respectively, and the nanohardness smoothly varies across the interface. Fig. 9(b) shows variations in the nanohardness across the interface for the HCSW in the As-Clad, Fine drawn and HT-2 conditions. Similar to the LCSW, the As-Clad HCSW shows a discontinuous change at the steel–copper interface with average nanohardness values of 3.6
Fig. 7. Inverse pole figure obtained from the steel region in HCSW. (a) is obtained from normal (⊥) and parallel () planes to the drawing direction in As-Clad, (b) is obtained from normal and parallel planes to the drawing direction in D-1, and (c) is obtained from normal and parallel planes to the drawing direction in D-2, (d) is obtained from normal and parallel planes to the drawing direction in HT-1, and (e) is obtained from normal and parallel planes to the drawing direction in fine drawn, (f) is obtained from normal and parallel planes to the drawing direction in HT-2.
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Fig. 8. Inverse pole figure obtained from the Cu region in LCSW. (a) is obtained from normal (⊥) and parallel () planes to the drawing direction in D-1, (b) is obtained from normal and parallel planes to the drawing direction in D-2, and (c) is obtained from normal and parallel planes to the drawing direction in fine drawn.
and 1.6 GPa in the steel and copper regions, respectively. In the Fine drawn condition, the steel and copper region’s nanohardness values increased to 5.7 and 2.0 GPa, respectively. Unlike the Fine drawn LCSW, the nanohardness change at the interface is not smoothly varying in the HCSW. Only post HT-2 does the HCSW nanohardness value across the interface smoothly vary from 2.9 to 2.3 GPa between the steel and copper, respectively. Fig. 10 shows the magnified IPF maps obtained from copper–steel interface of the LCSW and HCSW in the Fine drawn condition normal to the drawing direction. In the LCSW, Fig. 10(a), the grain size of the copper near the interface is a few microns, which is finer than those away from the interface. In contrast, the steel grains near the interface have similar grain sizes at and away from the interface. The HCSW’s copper and steel grains near and far from the interface exhibit similar sizes, Fig. 10(b), which are 1.2 m. Fig. 11(a) and (b) shows the stress distribution in the axial direction (S22 component) analyzed by FEM for Fine drawn LCSW and HCSW, respectively. Note that each figure shows only one-half of the wire and is radially symmetric on the other side. LCSW and HCSW showed similar stress distribution along drawing direction after the drawing. While the stress is distributed rather uniformly in the copper strip region, a large stress concentration is observed at the perimeter region of the steel as compared to the middle section in the steel. The region between the center and perimeter regions is free from a significant stress concentration. As compared to the LCSW, HCSW showed higher and localized stress concentration at the perimeter region and a larger stress gradient is observed from the perimeter to middle section of the steel region in the HCSW as compared to LCSW, Fig. 11(a) and (b). On the other hand, the logarithmic shear strain (LE12 component) distribution at the cop-
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Fig. 9. Variation in hardness across the interface measured by nanoindentation in (a) LCSW and (b) HCSW.
per strip is different between the LCSW and HCSW. In the LCSW, the copper strip near the copper–steel interface is more strained compared to that away from the interface, Fig. 11(c). Since the perimeter region of the steel wire is strained as well as the copper strip in the LCSW, Fig. 11(c) the copper and low carbon steel are co-deformed during drawing. In HCSW, the copper strip is subjected to a higher magnitude of shear strain as compared to that in the LCSW, Fig. 11(d). The perimeter region of the steel is subjected to a large stress concentration along the drawing direction and experiences smaller shear strain than the middle section in the high carbon steel and copper strip.
Fig. 10. Magnified inverse pole figure (IPF) maps obtained from steel–copper interface in (a) LCSW and (b) HCSW in fine drawn condition. Note that the image was obtained from the plane normal to the drawing direction.
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Fig. 11. Stress distribution of along the drawing direction (SS22 component) for (a) LCSW and (b) HCSW analyzed by FEM. Shear strain distribution (LE12 component) for (c) LCSW and (d) HCSW analyzed by FEM.
4. Discussion In the present study, we have investigated the microstructure evolution in two different bimetallic wires of LCSW and HCSW. Although previous work has reported the drawing behavior of copper and steel single wires individually [16–29], this paper has investigated the simultaneous co-deformation behavior between the two wires. Upon drawing, the hardness of each material increased because of work hardening. As shown in Fig. 1(a), the steel region in the LCSW had almost the same Vickers hardness values of approximately 250 VHN in the D-1, D-2 and Fine drawn conditions, respectively. On the other hand, the Vickers hardness of the steel region in HCSW increased up to 466 HV with subsequent draws to the Fine drawn condition. The higher hardness in the drawn HCSW is attributed to the finer filamentary microstructure, Fig. 6(c) and (d), which is formed as a consequence of the drawing of the wires with a pearlitic grain structure. Upon heat treating, there was a loss of the bulk hardness because of recrystallization as evident by the formation of equiaxed, finer grains, Fig. 6. The hardness values of the HCSW wire in the As-Clad and HT-1 conditions were nearly twice as large as those of the LCSW’s As-Clad and HT-1 conditions. Although the HT-2 was equivalent in time and temperature between the LCSW and HCSW, the HCSW plotted in Fig. 1(b) exhibited a greater percentage loss in Vickers hardness by the HT-2 heat treatment. For the HCSW, the different heat treatment temperature between the HT-1 to HT-2 resulted in the microstructure change from pearlite grains to the ferrite grains with the spherical Fe3 C precipitates along the grain boundaries as shown in Fig. 6(a), (b), (e) and (f). Therefore, the significant decrease in the hardness is attributed to the formation from the pearlite to ferrite grain structure. In the steel region, the main texture observed in the as-drawn LCSW and HCSW is the 1 1 0 fiber texture, as shown in Figs. 3 and 7. The 1 1 0 fiber texture in the Fine drawn LCSW is slightly diffused as compared to those in D-1 and D-2 conditions, although this was not the case for the HCSW. Since the strength difference is smaller between the low carbon steel and copper than between the high carbon steel and copper, the perimeter of the low carbon steel was easier to co-deform with copper and experienced larger shear deformation originating from the friction between the wire and the die upon drawing as compared to that of the high
carbon steel. Thus, the diffused texture could be explained by the shear deformation. The strong 1 1 0 fiber texture in the drawn sample was generally weakened in severity because of recrystallization during the post heat treatments, HT-1 and HT-2, for both the LCSW and HCSW. The recrystallization was evident by the change to an equiaxed, refined grain structure post heat treatment, Figs. 2 and 6. However, the 1 1 0 fiber texture feature still remained rather stronger in the HT-2 condition as compared to the HT-1 condition, particularly for the HCSW. This is attributed to the slower recrystallization kinetics because of (1) the lower heat treatment temperature in HT-2 than HT-1 for both LCSW and HCSW and (2) the Fe3 C particles in the grain boundaries for the HCSW that can inhibit (abnormal) grain growth for the heat treated condition studied. In the Cu region, a 1 1 1 –1 1 2 deformation direction developed with decreasing wire diameter after the D-1 and D-2 draws for both the LCSW and HCSW. While the global texture of drawn copper wire consists of a major 1 1 1 plus minor 1 0 0 fiber texture [16,21,27,28], this 1 1 1 –1 1 2 behavior is similar to the textured feature seen in the perimeter and middle sections of wire drawn copper [21,27,28]. Since the clad copper strip would be subjected to shear deformation as shown in Fig. 11(c) and (d), this could explain the evolution of the 1 1 1 –1 1 2 texture in these wires. On the other hand, the HCSW and LCSW in the Fine drawn condition showed the 1 1 1 fiber with intense 1 0 0 fiber texture indicating the occurrence of recrystallization during the Fine draw [28]. The copper grains in the HCSW are much finer, and rather equiaxed compared to the ones in the LCSW, Figs. 5 and 6. Since the high carbon steel exhibits higher hardness than LCSW, higher flow stress would be required when the fine drawing is conducted under similar conditions for these wires as shown in Table 1. In addition, the copper strip in the HCSW is thinner than that in the LCSW. Thus, the copper strip in the HCSW likely experienced a higher temperature because of the larger frictional heat between the die and the wire. Since the frictional heat was reported to cause the recrystallization in single copper wire [30,31], the larger amount of recrystallization would increase the intensity of the 1 0 0 fiber in the Fine drawn HCSW, Fig. 7, as compared to the LCSW. Arguably the most intriguing observation in the co-deformation microstructures is the grain size discrepancies at the copper–steel interfaces. The nanoindentation measurement showed that there was a discontinuous nanohardness change across the copper–steel interface in both As-Clad LCSW and HCSW, Fig. 9(a) and (b). Interestingly, a smoothly varying hardness across the copper–steel interface is observed in the Fine drawn LCSW in contrast to an abrupt hardness variation in the Fine drawn HCSW. As shown in the magnified IPF maps in Fig. 10(a), the copper grains near the interface were slightly smaller than those away from the interface for the LCSW while the HCSW’s copper grains were approximately the same size and shape at and away from the interface, Fig. 10(b). The grain refinement in the LCSW could contribute to an increase in hardness for the copper because of a Hall–Petch effect [32,33]. This increase towards the steel value would result in the smooth variation in the hardness in the fine drawn LCSW. The slightly refined copper grain near the interface are attributed to the shear strain gradient within the copper strip as depicted in FEM image of Fig. 11(c). This grain size discrepancy was also observed for the copper–steel interface in HCSW after the HT-2 treatment. As previously noted in Fig. 6(f), the grains at the steel region near the interface after HT-2 are coarser than those further from the interface. The difference in the stress localization shown by the FEM in Fig. 11(a) and (b), could also explain this grain size discrepancy. The HCSW experienced high and localized stress concentration at the perimeter of steel region during Fine drawn prior to the HT2. The highly localized stress in the fine drawn HCSW facilitated
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faster recrystallization and grain growth for the steel grains near the interface as compared to those far from the interface which caused the smooth hardness variation in the HT-2 condition. 5. Summary A series of LCSW and HCSW were subjected to drawing and heat treatments. Their hardness measurement and microstructure characterization delivered following conclusions: (1) The Vickers hardness measurement showed the HCSW had a higher strength as compared to the LCSW. However, the strength of the HCSW significantly decreased by the subsequent heat treatment of HT-2 to values lower than those of the drawn LCSW. This is attributed to the dispersion of the spherical Fe3 C particles along the grain boundary rather than forming a pearlitic grain structure. (2) Both LCSW and HCSW showed strong textured feature after drawing, and the texture is weakened by the heat treatment process because of recrystallization and grain growth. In the steel region, the strong 1 1 0 fiber texture could be observed in the as-drawn wires. With decreasing the wire diameter, the 1 1 0 fiber texture is diffuse in the LCSW while this was not the case for HCSW. In the copper region, a 1 1 1 –1 1 2 deformation direction, which is distributed along the 1 1 1 –1 0 0 symmetry boundary of the unit angle, developed for both LCSW and HCSW. The diffused 1 1 0 texture and the 1 1 2 deformation direction in the copper region have been explained by the shear deformation of the wire using FEM. (3) There is a grain size discrepancy at the steel–copper interface. In the Fine drawn LCSW, copper grains are slightly finer near the copper–steel interface than those away from the interface. This is because of the shear strain gradient in the copper region upon Fine draw LCSW shown by FEM. Although there was no such grain size discrepancy in the Fine drawn HCSW, the subsequent HT-2 heat treatment resulted in the coarsening of the steel grains near the interface. This is attributed to a large localized stress concentration at the perimeter of the steel region during the Fine drawing process. These grain size discrepancies caused the smooth variation in nanohardness.
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Acknowledgement This research was performed through The University of Alabama’s Materials Metrology Research Consortium through a grant from Fushi Copperweld, Inc. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16]
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructural evolution of copper clad steel bimetallic wire T.T. Sasaki a , M. Barkey b , G.B. Thompson a,∗ , Y. Syarif c , D. Fox c a
The University of Alabama, Department of Metallurgical & Materials Engineering, Tuscaloosa, AL 35487-0202, United States The University of Alabama, Department of Aerospace Engineering & Mechanics, Tuscaloosa, AL 35487-0202, United States c Fushi Copperweld, Inc., Fayetteville, TN 37334-7249, United States b
a r t i c l e
i n f o
Article history: Received 7 September 2010 Accepted 8 December 2010 Available online 14 December 2010 Keywords: Copper clad steel Drawing Texture Finite element modeling
a b s t r a c t We investigated the microstructure of two different bimetallic wires of Copper Clad Low Carbon Steel Wire (LCSW), which had a 1006 steel core, and Copper Clad High Carbon Steel Wire (HCSW), which had a 1055 steel core. The HCSW generally showed higher hardness than LCSW because of the pearlitic grain structure. A low temperature annealing at 720 ◦ C to the drawn HCSW caused a significant reduction of hardness, which was as low as that of an annealed LCSW. In general, both LCSW and HCSW showed strong global textured features after drawing, with the steel having a strong 1 1 0 fiber texture and the copper having a 1 1 1 –1 1 2 deformation direction. At the interface, a grain size discrepancy at the steel–copper interface was observed. Post-drawing, the LCSW copper grains exhibited refined grain sizes near the interface and has been explained in terms of shear strain gradient. The HCSW did not exhibit this copper grain size distribution but did exhibit a coarsening of the steel grains near the interface after a subsequent 720 ◦ C heat treatment. This is attributed to the large localized stress concentration at the perimeter of the steel region during the drawing process. The strain induced regions at the steel–copper interface have been simulated by finite element modeling. These grain size discrepancies caused the smooth variation in nanohardness across the interface. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Multi-component wires offer the ability to blend the properties of different metals into a single wire form to achieve improved performance such as electrical conduction or mechanical strength. In addition, these composite wires can assist in cost reduction by using less expensive wires as support or filler material in place of higher cost, single elemental wires [1–15]. In particular, copper clad steel (CCS) bimetallic wires can be a low cost alternative to single copper wires and has applications as the center conductor in high frequency coaxial cables, power transmission, and tracer wires for buried ceramics [10–15]. The copper sheath provides the low electrical resistance while the center core steel provides high structural strength, which can be tuned by the carbon content of the steel wire and heat treatment. To date, a number of studies on the structure–property relationship have been done for either Cu or steel single wires. In elemental metal drawn wire, the microstructure generally consists of filamentary grains elongated along the drawing direction with textured features [16]. In steel wires, the grains normally exhibit a strong
∗ Corresponding author at: The University of Alabama, Department of Metallurgical & Materials Engineering, 301 7th Avenue, 116 Houser Building, Tuscaloosa, AL 35487-0202, United States. Tel.: +1 205 348 1589; fax: +1 205 348 2164. E-mail address: [email protected] (G.B. Thompson). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.12.032
1 1 0 fiber texture [17–20]. In some specific processing routes, the deviation from the 1 1 0 fiber texture can be observed at the perimeter region for the wires [21,22]; similar variations in texture have been observed on the surface of rolled sheets [23–25]. On the other hand, Cu grains exhibit a major 1 1 1 and a minor 1 0 0 fiber texture in the as-drawn wire condition [16,26,27]. Texture deviations at the perimeter region and texture changes by annealing of the wire have also been reported [27–29]. There have been reports on CCS wires, sheets and extruded bars, focusing on the processing coupled with finite element modeling (FEM) [10–15]. To date, however, there has been limited work addressing the microstructure coupled to the simultaneous co-deformation behavior that occurs in the CCS bimetallic wires. Since steel and copper have different mechanical properties, these metals at the interface may deform differently to stay “connected” to each other during drawing. In the present study, we have selected two different CCS wires using either low- and high-carbon steel as the core rod and compared the microstructure evolution during the drawing process in terms of texture formation, grain structure, and local microstructure features at the steel–copper interface.
2. Experimental procedure Two CCS wires, one with a low carbon 1006 steel and the other with a high carbon 1055 steel have been clad to an oxygen-free cop-
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Table 1 Processing route of CCS bimetallic wires of (a) LCSW and (b) HCSW. The ‘D’ designates drawing stage and ‘HT’ designates the heat treatment stage. Diameter (mm)
Cu thickness (mm)
Drawing ◦
Heat treatment
Temp. ( C) (a) As-Clad D-1 HT-1 D-2 HT-2 Fine drawn
9.5 4.1 4.1 2.0 2.0 1.0
0.95 0.41 0.41 0.20 0.20 0.10
(b) As-Clad D-1 D-2 HT-1 Fine drawn HT-2
9.5 4.6 2.9 2.9 1.2 1.2
0.28 0.14 0.09 0.09 0.04 0.04
Rate (m/s)
Strain
55
7.1
1.7
55
10.2
1.4
55
20.3
1.4
55 55
7.1 10.2
1.5 2.4
55
15.2
1.8
per (99.95%) metal strip. The copper-strip was continuously formed around a moving steel rod during the cladding operation. The clad wire had an initial wire diameter of 9.5 mm with the steel core diameter being 7.62 mm for 1006 and 9.01 mm for 1055. Hereafter, the CCS wires are described as Low Carbon Steel Wire (LCSW) for the 1006 core steel and High Carbon Steel Wire (HCSW) for the 1055 core steel. Table 1(a) and (b) summarizes the processing conditions of the LCSW and HCSW. The wire diameter was reduced from 9.5 to 1.0 mm by a combination of three major drawing steps (D-1, D-2 and fine drawing) and intermediate heat treatments (HT1 and HT-2). The relative reduction of each metal, copper or steel, is consistent during drawing, e.g. both metals reduce at the same reduction fraction relative to their initial diameter, or, alternatively, the volume is conserved and one metal does not reduce any more than the other to compensate for discrepancies in mechanical property differences. The strain introduced during the drawing process was calculated by = ln(A0 /A), where A0 and A is the area before and after drawing. The bulk hardness values of the wires were measured by a Vickers hardness tester. The variation in the hardness across the copper–steel interface was measured by nanoindentation using a Hysitron TI900 TriboIndenter. Thirty indents were made with a spacing of 1 m between indents under a 4000 N load. The hardness measurements for both tests were performed on the normal plane to the drawing direction. Microstructure characterization was carried out using a field emission gun-scanning electron microscope (SEM), JEOL JSM7000F, equipped with an Oxford electron backscattered diffraction (EBSD) system and INCA software. The texture was measured by EBSD and is presented in the form of inverse pole figure map (IPF map) and inverse pole figure (IPF). The EBSD analyses were carried out from both normal and parallel planes to the drawing direction. The finite element modeling (FEM) was conducted using ABAQUS Explicit 6.9-1. An axi-symmetric model was constructed for the bimetallic wire using 4-noded axi-symmetric reduced integration elements. Table 2 summarizes the initial size of the wire, steel diameter and copper thickness for LCSW and HCSW. The model was made of elements tabulated in Table 2. For LCSW, 48 elements of steel and 12 elements of copper were used in the radial direction, respectively. For HCSW, 56 elements of steel and 4 elements of copper were used in the radial direction, respectively. The difference in number of elements was done because of the experimental differences in steel and copper diameters and thicknesses, respectively, between the LCSW and HCSW. The boundary conditions of the wire were set to an initial velocity of the wire at 15 m/s. The pressure at the lower end of the wire was 150 and 200 MPa and a pressure on the upper end of the wire was 10 and 50 MPa for
Temp. (◦ C)
Time (min)
950
15
720
480
950
21
720
480
Table 2 Initial size of wire, steel diameter and copper thickness for LCSW and HCSW used for FEM modeling.
Wire Total Diameter (mm) Steel diameter (mm) Copper thickness (mm) Length (mm) Element Radial direction (mm) Axial direction (mm)
LCSW
HCSW
2.052 1.642 0.2052 26.3
2.906 2.724 0.0908 26.3
0.0171 0.0496
0.0227 0.0496
the LCSW and HCSW, respectively. The pressure was applied in the vertical direction. Radial displacements were restrained for nodes on the axi-symmetry line. The die was modeled using 2-noded axi-symmetric rigid elements. The bearing length of the die was taken as 1 mm and the bell angle and back relief angle were 60◦ which mimics the experimental set-up. The interaction between the surface of the die and the contact surface of the wire was modeled with a coefficient of friction of 0.01. The die was initially offset from the surface of the wire by 0.0224 and 0.0455 mm and moved radially into the wire by a distance of 0.239 and 0.2 mm for LCSW and HCSW, respectively. After the wire was moved through the die, the reduction of area was 20%. The step time of the analysis was 0.003 s and the time that the die was moved into place was 0.0001 s. 3. Results Fig. 1 shows variations in Vickers hardness for the LCSW and HCSW as a function of the wire diameter. Note that the open and closed plots show the Vickers hardness values in the as-drawn and heat-treated conditions, respectively. In both wires, the hardness of the steel and copper increases after drawing because of work hardening and is recovered, i.e. lower hardness, partially or fully upon the heat treatments. The hardness values of the heat treated samples are slightly lower than the As-Clad conditions for both metals indicating that the sample is slightly work hardened during the cladding process. After each draw, the hardness of the LCSW was approximately the same, which is consistent that the strain was approximately equivalent among each draw, Table 1. The HCSW increased the hardness between the D-1 and D-2 draws. This is attributed to the increase in the strain between these two draws. Interestingly, the fine-drawn condition of the HCSW had a higher hardness value than the D-2 condition, even though the lower strain was introduced compared to the HCSW in the D-2
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Fig. 1. Variation in Vickers hardness as a function of wire diameter in (a) LCSW and (b) HCSW. Note that open and closed circles show the Vickers hardness in as-drawn condition and the one after heat treatment, respectively.
condition, Table 1 (b). The LCSW recovers with increasing temperature. In contrast, the HCSW exhibited a higher amount of recovery during a low temperature heat treatment, HT-2, of D-2, as seen in Fig. 1. Fig. 2 shows typical IPF maps obtained from normal and parallel planes to the drawing direction of the LCSW in the As-Clad, HT-2 and fine-drawn conditions. As shown in Fig. 2(a) and (b), both ferrite and copper grains are equiaxed in the As-Clad condition. The ferrite grains have an average size of approximately 8.7 m.
Upon HT-2 annealing, as shown in Fig. 2(c) and (d), the copper grains are recrystallized and exhibit annealing twins as indicated by black arrows. Their average grain size is approximately 153 m. The ferrite grains in the HT-2 condition are equiaxed and are slightly refined to 6.1 m as compared to those in the As-Clad condition. In the Fine drawn condition, the filamentary grain microstructure from the drawing process is clearly present, Fig. 2(e) and (f). The drawing process causes a loss of the distinct equiaxed grains. A similar filamentary microstructure, as in Fig. 2(e) and (f), was observed for the D-1 and D-2 LCSW wires. Fig. 3 shows IPFs obtained from the steel region at each processing step for the LCSW. Note that the IPFs are obtained from normal and parallel planes to the drawing direction. Fig. 3(a) confirms ferrite grains do not show any obvious textured feature in the As-Clad condition. During the D-1 drawing step, a strong 1 1 0 fiber texture is developed for the ferrite, Fig. 3(b). During the HT-1 heat treatment, the 1 1 0 fiber texture becomes diffused as seen in Fig. 3(c). But the 1 1 0 direction of the ferrite grains, Fig. 3(c), are qualitatively more aligned to the drawing direction as compared to the As-Clad sample, Fig. 3(a). The D-2 draw again recovers the strong 1 1 0 fiber texture, Fig. 3(d). This strong fiber texture is subsequently lost post HT-2, Fig. 3(e). Upon the fine drawing, the 1 1 0 fiber texture is recovered, Fig. 3(f), but slightly diffused compared to the one in D-1 and D-2 conditions. The quantification of the texture feature is difficult for the Cu grains in the As-Clad, HT-1 and HT-2 conditions since the large grain sizes limited the counting statistics over the field of view studied, as seen in Fig. 2. However, obvious deformation textures can be seen in the drawn wires. The IPF maps, Fig. 4(a) and (b) shows that the copper grains for the LCSW in the D-1 condition developed a tendency for the 1 1 1 –1 1 2 directions, i.e. the distribution along the 1 1 1 –1 0 0 symmetry boundary of the unit angle, aligning with the drawing direction. In the D-2 condition, the majority of the grains are present with their major 1 1 2 direction parallel to the drawing direction and the remaining grains indicating a minor 1 0 0 alignment, as seen in Fig. 4(c) and (d). The Fine drawn texture, Fig. 4(e) and (f), shows the copper grains also form diffuse texture along 1 1 1 –1 0 0 symmetry boundary of the unit triangle. Compared to copper grains in the D-1 and D-2 conditions, these copper grains in the fine-drawn condition showed slightly stronger 1 0 0 fiber texture.
Fig. 2. Inverse pole figure (IPF) maps obtained from copper–steel interface in LCSW. (a) and (b) are those obtained from As-Clad wire taken from normal (⊥) and parallel () planes to drawing direction. (c) and (d) are heat treated (HT-2) wire before fine drawing taken from normal and parallel direction to drawing direction. (e) and (f) are from fine drawn wire taken from normal and parallel direction to drawing direction, respectively.
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Fig. 3. Inverse pole figure obtained from the steel region in LCSW. (a) is obtained from normal (⊥) and parallel () planes to the drawing direction in As-Clad, (b) is obtained from normal and parallel planes to the drawing direction in D-1, and (c) is obtained from normal and parallel planes to the drawing direction in HT-1, (d) is obtained from normal and parallel planes to the drawing direction in D-2, and (e) is obtained from normal and parallel planes to the drawing direction in HT-2, (f) is obtained from normal and parallel planes to the drawing direction in fine drawn.
Fig. 5 is backscattered electron SEM micrographs of the HCSWs in HT-1, fine-drawn and HT-2 conditions observed from the parallel plane with respect to the drawing direction. In Fig. 5(a), the steel region in the HT-1 consists of equiaxed ferrite and pearlite grains. Upon Fine drawing, the ferrite grains are elongated along the drawing direction, Fig. 5(b), but the copper grains are equiaxed compared to the steel grains. The subsequent heat treatment of HT-2, Fig. 5(c), the ferrite grains near the Cu interface are coarser than those further away from the interface. Unlike the HT-1 condition, the HT-2 microstructure does not show the pearlite grain
Fig. 4. Inverse pole figure obtained from the Cu region in LCSW. (a) is obtained from normal (⊥) and parallel () planes to the drawing direction in D-1, (b) is obtained from normal and parallel planes to the drawing direction in D-2, and (c) is obtained from normal and parallel planes to the drawing direction in fine drawn.
structure. The Fe3 C phase, the spotty dark features, is present as spherical particles mainly along the grain boundaries. Fig. 6 shows typical inverse pole figure (IPF) maps obtained from normal and parallel planes to the drawing direction of HCSW in the As-Clad, Fine drawn and HT-2 conditions. As shown in Fig. 6(a) and (b), both steel and copper regions in the As-Clad sample consists of equiaxed grains similar to those reported above for the LCSW in AsClad condition, Fig. 2(a) and (b). The ferrite grains and copper grains have an average size of approximately 5.7 m and 67 m, respectively. In the Fine drawn condition, Fig. 6(c) and (d), both ferrite and copper grains have the typical filamentary grain structure formed from the drawing process. The filamentary copper grains consist of the equiaxed grains with low angle boundaries. Similar fine filamentary grain structures were observed for the other D-1 and D-2 drawn wires for the HCSW. Upon HT-2 heat treatment, Fig. 6(e) and (f), the steel grains recover the equiaxed grain structure. As noted previously, the grains near the copper interface are coarser and elongated along the drawing direction; this will be developed further in Section 4. Fig. 7 is the inverse pole figures (IPFs) obtained from the steel region in the HCSW at each processing step. Note that these IPFs are obtained from normal and parallel planes to the drawing direction. The steel region does not show any obvious textured features in the As-Clad condition, Fig. 7(a). The D-1 drawing substantially develops the 1 1 0 fiber texture, Fig. 7(b). As shown in Fig. 7(c), the 1 1 0 fiber texture becomes more intense with the subsequent D-2 draw. During the HT-1 heat treatment, the drawn 1 1 0 fiber texture becomes more diffused, Fig. 7(d), although the 1 1 0 fiber texture is still observable. The strong 1 1 0 fiber texture is recovered with the Fine drawing as shown in Fig. 7(e). The 1 1 0 strong texture in the HT-2, Fig. 7(f), is still retained even after heat treatment. Similar to the LCSW, the large copper grains in the heat treated HCSW limited the counting statistics in the EBSD scans as seen in Fig. 8. Again, the IPFs for copper grains in Fig. 8 are obtained from normal and parallel planes to the drawing direction. In the D-1 condition, the majority of the grains are present with their 1 1 1 direction parallel to the drawing direction, Fig. 8(a). The copper grains in D-2 condition develop a relatively diffused texture as compared to those in the D-1 condition. A slightly higher density of the 1 1 w textures parallel to the drawing direction developed for D-2 condition, Fig. 8(b). Similar to the LCSW, a tendency for the 1 1 1 –1 1 2 (distributed along the 1 1 1 –1 0 0 symmetry boundary of the unit angle) is aligned with the drawing direction of the 1 1 1 fiber texture. In the Fine drawn condition, the copper grains form a diffuse texture for the major 1 1 1 plus the minor 1 0 0 fiber orientations, Fig. 8(c). For the Fine drawn condition, the copper grains in the HCSW show stronger 1 0 0 fiber texture compared to the ones in the D-1 and D-2 conditions similar to those in the LCSW.
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Fig. 5. Backscattered electron SEM images of the HCSWs in (a) HT-1, (b) fine drawn and (c) HT-2 conditions. Note that these images are taken from the parallel plane to the drawing direction, and the drawing direction is horizontal direction.
Fig. 6. Inverse pole figure (IPF) maps obtained from Fe/Cu interface in HCSW. (a) and (b) are those obtained from As-Clad wire taken from normal (⊥) and parallel () planes to drawing direction. (c) and (d) are from the sample in fine drawn wire taken from normal and parallel planes to drawing direction, respectively. (e) and (f) are obtained from HT-2 heat treated wire after fine drawing taken from normal and parallel planes to drawing direction.
Fig. 9(a) shows the variation in nanohardness across the interface for the LCSW in the As-Clad and Fine drawn conditions. In the As-Clad condition, the average nanohardness values are 2.4 and 1.4 GPa in the steel and copper regions, respectively. A discontinuous change at the steel–copper interface exists in the As-Clad wire. In the Fine drawn condition, the steel and copper regions have a
higher nanohardness value of 3.9 and 2.1 GPa, respectively, and the nanohardness smoothly varies across the interface. Fig. 9(b) shows variations in the nanohardness across the interface for the HCSW in the As-Clad, Fine drawn and HT-2 conditions. Similar to the LCSW, the As-Clad HCSW shows a discontinuous change at the steel–copper interface with average nanohardness values of 3.6
Fig. 7. Inverse pole figure obtained from the steel region in HCSW. (a) is obtained from normal (⊥) and parallel () planes to the drawing direction in As-Clad, (b) is obtained from normal and parallel planes to the drawing direction in D-1, and (c) is obtained from normal and parallel planes to the drawing direction in D-2, (d) is obtained from normal and parallel planes to the drawing direction in HT-1, and (e) is obtained from normal and parallel planes to the drawing direction in fine drawn, (f) is obtained from normal and parallel planes to the drawing direction in HT-2.
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Fig. 8. Inverse pole figure obtained from the Cu region in LCSW. (a) is obtained from normal (⊥) and parallel () planes to the drawing direction in D-1, (b) is obtained from normal and parallel planes to the drawing direction in D-2, and (c) is obtained from normal and parallel planes to the drawing direction in fine drawn.
and 1.6 GPa in the steel and copper regions, respectively. In the Fine drawn condition, the steel and copper region’s nanohardness values increased to 5.7 and 2.0 GPa, respectively. Unlike the Fine drawn LCSW, the nanohardness change at the interface is not smoothly varying in the HCSW. Only post HT-2 does the HCSW nanohardness value across the interface smoothly vary from 2.9 to 2.3 GPa between the steel and copper, respectively. Fig. 10 shows the magnified IPF maps obtained from copper–steel interface of the LCSW and HCSW in the Fine drawn condition normal to the drawing direction. In the LCSW, Fig. 10(a), the grain size of the copper near the interface is a few microns, which is finer than those away from the interface. In contrast, the steel grains near the interface have similar grain sizes at and away from the interface. The HCSW’s copper and steel grains near and far from the interface exhibit similar sizes, Fig. 10(b), which are 1.2 m. Fig. 11(a) and (b) shows the stress distribution in the axial direction (S22 component) analyzed by FEM for Fine drawn LCSW and HCSW, respectively. Note that each figure shows only one-half of the wire and is radially symmetric on the other side. LCSW and HCSW showed similar stress distribution along drawing direction after the drawing. While the stress is distributed rather uniformly in the copper strip region, a large stress concentration is observed at the perimeter region of the steel as compared to the middle section in the steel. The region between the center and perimeter regions is free from a significant stress concentration. As compared to the LCSW, HCSW showed higher and localized stress concentration at the perimeter region and a larger stress gradient is observed from the perimeter to middle section of the steel region in the HCSW as compared to LCSW, Fig. 11(a) and (b). On the other hand, the logarithmic shear strain (LE12 component) distribution at the cop-
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Fig. 9. Variation in hardness across the interface measured by nanoindentation in (a) LCSW and (b) HCSW.
per strip is different between the LCSW and HCSW. In the LCSW, the copper strip near the copper–steel interface is more strained compared to that away from the interface, Fig. 11(c). Since the perimeter region of the steel wire is strained as well as the copper strip in the LCSW, Fig. 11(c) the copper and low carbon steel are co-deformed during drawing. In HCSW, the copper strip is subjected to a higher magnitude of shear strain as compared to that in the LCSW, Fig. 11(d). The perimeter region of the steel is subjected to a large stress concentration along the drawing direction and experiences smaller shear strain than the middle section in the high carbon steel and copper strip.
Fig. 10. Magnified inverse pole figure (IPF) maps obtained from steel–copper interface in (a) LCSW and (b) HCSW in fine drawn condition. Note that the image was obtained from the plane normal to the drawing direction.
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Fig. 11. Stress distribution of along the drawing direction (SS22 component) for (a) LCSW and (b) HCSW analyzed by FEM. Shear strain distribution (LE12 component) for (c) LCSW and (d) HCSW analyzed by FEM.
4. Discussion In the present study, we have investigated the microstructure evolution in two different bimetallic wires of LCSW and HCSW. Although previous work has reported the drawing behavior of copper and steel single wires individually [16–29], this paper has investigated the simultaneous co-deformation behavior between the two wires. Upon drawing, the hardness of each material increased because of work hardening. As shown in Fig. 1(a), the steel region in the LCSW had almost the same Vickers hardness values of approximately 250 VHN in the D-1, D-2 and Fine drawn conditions, respectively. On the other hand, the Vickers hardness of the steel region in HCSW increased up to 466 HV with subsequent draws to the Fine drawn condition. The higher hardness in the drawn HCSW is attributed to the finer filamentary microstructure, Fig. 6(c) and (d), which is formed as a consequence of the drawing of the wires with a pearlitic grain structure. Upon heat treating, there was a loss of the bulk hardness because of recrystallization as evident by the formation of equiaxed, finer grains, Fig. 6. The hardness values of the HCSW wire in the As-Clad and HT-1 conditions were nearly twice as large as those of the LCSW’s As-Clad and HT-1 conditions. Although the HT-2 was equivalent in time and temperature between the LCSW and HCSW, the HCSW plotted in Fig. 1(b) exhibited a greater percentage loss in Vickers hardness by the HT-2 heat treatment. For the HCSW, the different heat treatment temperature between the HT-1 to HT-2 resulted in the microstructure change from pearlite grains to the ferrite grains with the spherical Fe3 C precipitates along the grain boundaries as shown in Fig. 6(a), (b), (e) and (f). Therefore, the significant decrease in the hardness is attributed to the formation from the pearlite to ferrite grain structure. In the steel region, the main texture observed in the as-drawn LCSW and HCSW is the 1 1 0 fiber texture, as shown in Figs. 3 and 7. The 1 1 0 fiber texture in the Fine drawn LCSW is slightly diffused as compared to those in D-1 and D-2 conditions, although this was not the case for the HCSW. Since the strength difference is smaller between the low carbon steel and copper than between the high carbon steel and copper, the perimeter of the low carbon steel was easier to co-deform with copper and experienced larger shear deformation originating from the friction between the wire and the die upon drawing as compared to that of the high
carbon steel. Thus, the diffused texture could be explained by the shear deformation. The strong 1 1 0 fiber texture in the drawn sample was generally weakened in severity because of recrystallization during the post heat treatments, HT-1 and HT-2, for both the LCSW and HCSW. The recrystallization was evident by the change to an equiaxed, refined grain structure post heat treatment, Figs. 2 and 6. However, the 1 1 0 fiber texture feature still remained rather stronger in the HT-2 condition as compared to the HT-1 condition, particularly for the HCSW. This is attributed to the slower recrystallization kinetics because of (1) the lower heat treatment temperature in HT-2 than HT-1 for both LCSW and HCSW and (2) the Fe3 C particles in the grain boundaries for the HCSW that can inhibit (abnormal) grain growth for the heat treated condition studied. In the Cu region, a 1 1 1 –1 1 2 deformation direction developed with decreasing wire diameter after the D-1 and D-2 draws for both the LCSW and HCSW. While the global texture of drawn copper wire consists of a major 1 1 1 plus minor 1 0 0 fiber texture [16,21,27,28], this 1 1 1 –1 1 2 behavior is similar to the textured feature seen in the perimeter and middle sections of wire drawn copper [21,27,28]. Since the clad copper strip would be subjected to shear deformation as shown in Fig. 11(c) and (d), this could explain the evolution of the 1 1 1 –1 1 2 texture in these wires. On the other hand, the HCSW and LCSW in the Fine drawn condition showed the 1 1 1 fiber with intense 1 0 0 fiber texture indicating the occurrence of recrystallization during the Fine draw [28]. The copper grains in the HCSW are much finer, and rather equiaxed compared to the ones in the LCSW, Figs. 5 and 6. Since the high carbon steel exhibits higher hardness than LCSW, higher flow stress would be required when the fine drawing is conducted under similar conditions for these wires as shown in Table 1. In addition, the copper strip in the HCSW is thinner than that in the LCSW. Thus, the copper strip in the HCSW likely experienced a higher temperature because of the larger frictional heat between the die and the wire. Since the frictional heat was reported to cause the recrystallization in single copper wire [30,31], the larger amount of recrystallization would increase the intensity of the 1 0 0 fiber in the Fine drawn HCSW, Fig. 7, as compared to the LCSW. Arguably the most intriguing observation in the co-deformation microstructures is the grain size discrepancies at the copper–steel interfaces. The nanoindentation measurement showed that there was a discontinuous nanohardness change across the copper–steel interface in both As-Clad LCSW and HCSW, Fig. 9(a) and (b). Interestingly, a smoothly varying hardness across the copper–steel interface is observed in the Fine drawn LCSW in contrast to an abrupt hardness variation in the Fine drawn HCSW. As shown in the magnified IPF maps in Fig. 10(a), the copper grains near the interface were slightly smaller than those away from the interface for the LCSW while the HCSW’s copper grains were approximately the same size and shape at and away from the interface, Fig. 10(b). The grain refinement in the LCSW could contribute to an increase in hardness for the copper because of a Hall–Petch effect [32,33]. This increase towards the steel value would result in the smooth variation in the hardness in the fine drawn LCSW. The slightly refined copper grain near the interface are attributed to the shear strain gradient within the copper strip as depicted in FEM image of Fig. 11(c). This grain size discrepancy was also observed for the copper–steel interface in HCSW after the HT-2 treatment. As previously noted in Fig. 6(f), the grains at the steel region near the interface after HT-2 are coarser than those further from the interface. The difference in the stress localization shown by the FEM in Fig. 11(a) and (b), could also explain this grain size discrepancy. The HCSW experienced high and localized stress concentration at the perimeter of steel region during Fine drawn prior to the HT2. The highly localized stress in the fine drawn HCSW facilitated
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faster recrystallization and grain growth for the steel grains near the interface as compared to those far from the interface which caused the smooth hardness variation in the HT-2 condition. 5. Summary A series of LCSW and HCSW were subjected to drawing and heat treatments. Their hardness measurement and microstructure characterization delivered following conclusions: (1) The Vickers hardness measurement showed the HCSW had a higher strength as compared to the LCSW. However, the strength of the HCSW significantly decreased by the subsequent heat treatment of HT-2 to values lower than those of the drawn LCSW. This is attributed to the dispersion of the spherical Fe3 C particles along the grain boundary rather than forming a pearlitic grain structure. (2) Both LCSW and HCSW showed strong textured feature after drawing, and the texture is weakened by the heat treatment process because of recrystallization and grain growth. In the steel region, the strong 1 1 0 fiber texture could be observed in the as-drawn wires. With decreasing the wire diameter, the 1 1 0 fiber texture is diffuse in the LCSW while this was not the case for HCSW. In the copper region, a 1 1 1 –1 1 2 deformation direction, which is distributed along the 1 1 1 –1 0 0 symmetry boundary of the unit angle, developed for both LCSW and HCSW. The diffused 1 1 0 texture and the 1 1 2 deformation direction in the copper region have been explained by the shear deformation of the wire using FEM. (3) There is a grain size discrepancy at the steel–copper interface. In the Fine drawn LCSW, copper grains are slightly finer near the copper–steel interface than those away from the interface. This is because of the shear strain gradient in the copper region upon Fine draw LCSW shown by FEM. Although there was no such grain size discrepancy in the Fine drawn HCSW, the subsequent HT-2 heat treatment resulted in the coarsening of the steel grains near the interface. This is attributed to a large localized stress concentration at the perimeter of the steel region during the Fine drawing process. These grain size discrepancies caused the smooth variation in nanohardness.
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Acknowledgement This research was performed through The University of Alabama’s Materials Metrology Research Consortium through a grant from Fushi Copperweld, Inc. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16]
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