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Microstructure and properties of a silver–erbium oxide alloy
Journal of Alloys and Compounds 454 (2008) 292–296
Microstructure and properties of a silver–erbium oxide alloy D.M. Herman a , G.H. Cao b , A.T. Becker a,b , A.M. Russell a,b,∗ , A.P. Constant a a
b
Materials Science and Engineering Department, 2220 Hoover Hall, Iowa State University, Ames, IA 50011, USA Materials and Engineering Physics Program, 126 Metals Development Building, Ames Laboratory of the USDOE, Ames, IA 50011, USA Received 7 July 2006; accepted 10 December 2006 Available online 16 January 2007
Abstract A Ag–4 wt.% Er solid solution alloy was produced by arc melting and quenching to room temperature. When this alloy was annealed in air at various temperatures, the Er was converted to ∼1-m thick Er2 O3 lamellae in the Ag matrix. Subsequent swaging and stress-relief annealing converted the lamellar microstructure to predominantly spheroidal Er2 O3 particles. The alloys were examined by optical microscopy, scanning electron microscopy (SEM), and transmission electron microscopy (TEM). Electrical resistivities of 2.0–3.3 cm and hardness values of 83–162 HV were measured. © 2007 Elsevier B.V. All rights reserved. Keywords: Metal matrix composites; Rare earth alloys and compounds; Electrical transport; Microstructure; Transmission electron microscopy, TEM
1. Introduction Several composite materials with a silver (Ag) matrix and a dispersed oxide phase have been developed for electrical contact applications. Ag–CdO composites were among the first of these materials, but they suffer from both physical properties limitations and Cd toxicity concerns, which motivated further research into similar materials, such as Ag–SnO2 and Ag–ZnO with ternary additions [1,2]. Much of the current research into silver–metal oxide materials involves optimizing the dispersion of the oxide phase by reducing the particle size, achieving a more uniform distribution [1], and improving the adhesion between the Ag and ceramic interface [3] within the composite. More recently, rare earth materials have been used as a ternary addition [4] to optimize the properties of another oxide [5]. Research into the effectiveness of Ag-rare earth oxide contact materials has been limited [6]. The Ag–Er2 O3 specimens in this study were prepared by arc melting Ag and Er, solutionizing the alloy at high temperature in an inert gas environment, quenching in water, and then annealing in air to internally oxidize the Er to form Er2 O3 second phase
particles in a Ag matrix. This process offers a simpler processing alternative to P/M with the potential to achieve more uniform dispersion of the oxide phase. Internal oxidation has been performed on Ag–Sn alloys [1]. Due to the rather noble nature of Sn, a secondary addition (e.g., In, Te, or Bi) is required to promote the internal oxidation reaction [7,8]. The Ag–Er alloy used in this study did not require ternary alloying elements to achieve oxidation of the Er when the metal was annealed in air. Other investigators have studied Ag-rare earth oxide materials produced by internal oxidation of the rare earth metal [4–6,9]. Reports from these studies have described the resulting microstructure and hardness but not the electrical properties. Elssner and Gebhardt [6] reported hardness values in Ag–0.28 at.% Er that were similar to those measured in this study. He also found that this composition and other dilute rare earth-Ag alloys maintained their high room temperature hardness values at temperatures well above room temperature. For example, the 110 HV hardness of Elssner’s Ag–0.28 at.% Er alloy changed little from room temperature to 800 ◦ C. Above 800 ◦ C the hardness decreased exponentially as the temperature approached the melting temperature. 2. Experimental procedure
∗
Corresponding author at: Materials Science and Engineering Department, 2220 Hoover Hall, Iowa State University, Ames, IA 50011, USA. Tel.: +1 515 294 3204; fax: +1 515 294 5444. E-mail address: [email protected] (A.M. Russell). 0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2006.12.036
To make a 96 wt.% Ag–4 wt.% Er button, 99.99% pure Ag was arc melted in an Ar atmosphere with 99.9% pure Er obtained from the Ames Laboratory Materials Preparation Center [10]. The sample was melted six times to ensure
D.M. Herman et al. / Journal of Alloys and Compounds 454 (2008) 292–296 uniformity. The Ag–Er equilibrium phase diagram indicates that a Ag–4 wt.% Er alloy would form a single-phase solid solution at 750 ◦ C. If cooled slowly enough to room temperature to preserve equilibrium, such an alloy would transform to a more dilute Ag solid solution plus Ag51 Er14 intermetallic compound. To avoid forming the Ag51 Er14 phase, the Ag–Er arc melted button was sealed in a Hefilled, Pyrex ampoule, annealed for 25 h at 750 ◦ C, and submerged in water as the ampoule was broken to quench the metal, forming a metastable, single-phase solid solution. The Ag–4% Er button was cut into four pieces that were internally oxidized at 400, 500, 600, or 700 ◦ C by annealing for 24 h in air, followed by air cooling. The four samples were mounted, polished, and etched with dilute K2 Cr2 O7 –NaCl–H2 SO4 acid solution to reveal the microstructure for optical metallography. SEM was performed on all samples before and after etching. The TEM specimen was prepared from the sample that had been internally oxidized at 500 ◦ C by cutting, grinding, and ion-milling. TEM bright field imaging and selected area electron diffraction (SAED) were performed using a Philips CM30 electron microscope equipped with a LaB6 electron source, operated at an accelerating voltage of 300 kV. Hardness measurements were taken using a LECO microhardness tester with a Vickers indenter. The load was 100 gf with a dwell time of 13 s for all hardness tests. X-ray diffraction was performed on the sample annealed at 700 ◦ C. The bulk sample was used, not powder. A scan of 25 to 75◦ 2θ was performed with a scan rate of 0.004◦ 2θ s−1 . The specimen annealed at 700 ◦ C was placed into a 1.27 cm diameter, thinwalled stainless steel sleeve and swaged at room temperature to a diameter of 2.54 mm. At this point, the stainless steel was removed, and the bare metal was swaged to a final diameter of 0.81 mm. A 30- min stress-relief anneal was performed in air at 300 ◦ C when the specimen diameter was 2.54 mm. It is estimated that this swaging produced a true strain of approximately 5.0. Exact determination of the true strain is not possible, since the starting specimen geometry was not cylindrical. A four-point conductivity test was performed at room temperature on the swaged 700 ◦ C sample. Current was passed between two outer contact points along the cylindrical wire length with a voltage drop being measured between two inner contact points.
3. Experimental results Upon removal from the furnace used for the solutionizing anneal, the samples and the Pyrex ampoule were discolored. The surface of the samples had a pinkish hue (Er2 O3 is pink), leading to the conclusion that the Er in the samples had either reduced some of the oxides in the glass during the anneal or had reacted with the water during the quench, forming a layer of Er2 O3 on the samples. This oxide layer was found to be only superficial, and it was removed before further testing. The microstructure shown in Fig. 1 was observed in the optical micrographs taken on all four samples. This microstructure was dominant and appeared to be a lamellar, rather than acicular, second phase. Individual variations of the spacing between the lamellar microstructure, thickness of the darker phase, and the orientation of the lamellae of this second phase were seen both within individual samples and amongst the four specimens. The microstructure shown in Fig. 1 could be distinguished in both back-scattered electron imaging and secondary electron imaging modes in the SEM. The hardness of all samples measured in this study was much higher than the 25 HV microhardness of pure, annealed Ag. The 400, 500, and 700 ◦ C samples showed similar hardness values (Table 1), but the 600 ◦ C sample showed a lower average hardness with a larger standard deviation than the
293
Fig. 1. Optical micrograph of Ag–Er alloy annealed in air at 400 ◦ C. EDS–SEM examination of this specimen showed that the darker phase contained higher Er content. Table 1 Vickers microhardness data (HV = Vickers microhardness number) Annealing temperature (◦ C)
Average hardness (HV)
Standard deviation (HV)
400 500 600 700
103.6 101.4 83.0 101.6
2.1 5.9 26.9 4.2
other three specimens. Hardness measurements were taken from both transverse and longitudinal sections of the 700 ◦ C sample that had been swaged to a diameter of 2.54 mm before being given a stress-relief anneal. Both types of sections displayed a mixed microstructure with some regions possessing a lamellar microstructure and others spheroidal particles (Fig. 2). The two different microstructures were found to have differing microhardness values (Table 2).
Fig. 2. A 700 ◦ C optical micrograph of the specimen oxidized at 700 ◦ C, swaged, before being annealed at 300 ◦ C showing both the spheroidal phase (centered in the micrograph) and the lamellar phase (near the upper right and lower left regions of the micrograph; lamellae are difficult to resolve at this low magnification).
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Table 2 Vickers microhardness measurements in specimen internally oxidized at 700 ◦ C and swaged to 2.54 mm diameter
Transverse section–spheroidal regions Transverse section–lamellar regions
Average hardness (HV)
Standard deviation (HV)
161.6 103.9
11.2 4.8
Fig. 5. EDS line scan showing higher Er and O content as several lamellae similar to those shown in Fig. 1 are crossed in the specimen annealed at 400 ◦ C.
Fig. 3. Higher magnification optical micrograph of the spheroidal phase structure in Fig. 2.
size larger than about 10 m. It may be that grain growth during the 25 h 750◦ C solutionizing anneal produced a coarse grain structure in the single-phase solid solution Ag–Er alloy that was retained at room temperature.
In the transverse cross-section of the sample oxidized at 700 ◦ C and subsequently swaged before being annealed at 300◦ C, the spheroidal microstructure dominated, although the lamellar microstructure was visible near the rod’s surface. Energy dispersive spectroscopy (EDS) showed that lamellae (Fig. 1) and spheroidal particles (Figs. 2 and 3) correspond to lower concentrations of Ag and higher concentrations of Er and O (Figs. 4 and 5). TEM of the sample annealed at 500 ◦ C showed the size, crystal structure, and morphology of the Er2 O3 particles (Figs. 6–8). The X-ray diffraction data from the sample annealed at 700 ◦ C were compared to published powder diffraction files of Ag, Er2 O3 , and Ag51 Er14 . The XRD data matched the peaks of Ag and Er2 O3 ; no peaks from the Ag51 Er14 intermetallic compound were observed. The complex microstructure in these specimens made grain size determination difficult. Bright field TEM showed no grain boundaries in the thinned region of the foil, suggesting a grain
Fig. 4. EDS line scan of spheroidal microstructure in the specimen annealed at 700 ◦ C (Fig. 3). The scan crossed a spheroid from 2 to 8 m on the abscissa.
Fig. 6. TEM images of (a) two Er2 O3 particles (A and B) and the matrix phase, (b) lamellar Er2 O3 phase.
D.M. Herman et al. / Journal of Alloys and Compounds 454 (2008) 292–296
Fig. 7. SAED patterns of particle A in Fig. 6 along (a) [0 0 1], (b) [0 1 1], and (c) [1 0 3] zone axes. Indexing shows that particle A has cubic crystal structure ˚ which matches that of Er2 O3 . (d) SAED pattern of the particle with a0 = 10.5 A, B along [1 1 2] zone axis, which also indexes as Er2 O3 .
295
defects in the crystal. At lower temperatures, O diffuses in the bulk Ag via an interstitial mechanism through low resistance diffusion paths such as grain boundaries and the free surface. At higher temperatures atomic O begins to segregate onto low indexed planes. It is possible that this process occurs due to an interstitial diffusion mechanism where O substitutes for Ag atoms in the lattice. Such a planar distribution of O atoms could be the precursor to Er2 O3 formation on those same planes, yielding the lamellar microstructure shown in Fig. 1. The change to a predominantly spheroidal microstructure after swaging and annealing is apparently the result of surface energy minimization in the previously existing lamellar structure that accompanied recrystallization in the heavily cold-worked Ag matrix. It was somewhat surprising that the microstructures and hardness values of the four specimens annealed in air at 400, 500, 600, and 700 ◦ C were all so similar. In all four of these specimens, there appeared to be sufficient diffusion to segregate the oxide phase into lamellae with micron-scale dimensions. This seems to suggest that the annealing temperature that would produce the finest possible oxide phase size (and hence, the highest hardness) may lie below 300 ◦ C, a processing option that was not explored in this study. 5. Conclusions
Fig. 8. SAED patterns of the matrix Ag along (a) [0 0 1], (b) [0 1 1], and (c) [−1 1 2] zone axes. Indexing shows that the matrix has a lattice parameter ˚ which matches that of Ag. a0 = 4.2 A,
A four-point conductivity measurement was taken on the wire, which showed a resistivity of ∼3.3 cm, which is substantially higher than the 1.59 cm resistivity of pure, annealed Ag. After the wire had been annealed at 400 ◦ C for 30 min, the same test showed a resistivity of ∼2.0 cm (∼86% IACS). The resistivity values are approximate since some exfoliation occurred on the wire surface during swaging, giving a wire geometry that was not perfectly cylindrical. Since the spalling slightly reduces the cross-sectional area for current flow, it is likely that the specimen’s actual resistivity is slightly lower than the values measured in the four-point conductivity test.
The Ag–Er2 O3 material of this study possesses physical properties that would make it a useful material for electrical contacts. It was seen in previous studies [6] that rare earth concentrations under 1 wt.% could still create a significant improvement to the hardness of Ag. Since lower Er concentrations would be expected to give higher electrical and thermal conductivity, a lower concentration would be preferable. Although the highpurity Er used in this project is more costly than Ag, the costs of commercial-purity rare earth metals are roughly similar to (and often lower than) that of Ag. Additionally internal oxidation is a simpler processing method than powder metallurgy. Current methods of internal oxidation require ternary alloy additions to promote diffusion and internal oxidation of the primary oxide formers, and high O partial pressures are usually used to promote oxidation. These two difficulties to the use of internal oxidation could be eliminated if a more reactive second metal such as a rare earth metal were used. Although a comparative test of actual switching performance would be necessary to determine how well existing Ag–metal oxide materials compare to Ag-rare earth oxides, this study shows that rare earth oxides have the potential to improve the switch and contact performance of Ag-based contact materials.
4. Discussion Acknowledgements The lamellar microstructure (Fig. 1) could be the result of preferential diffusion mechanisms for oxygen (O) in Ag. It was observed by Nagy et al. [11] that preferential diffusion directions exist for O movement in the Ag lattice and among
The authors gratefully acknowledge the assistance of J. Peters in performing XRD, S. Williams for TEM sample preparation, and B. Stumphy for solutionizing the samples.
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D.M. Herman et al. / Journal of Alloys and Compounds 454 (2008) 292–296
References [1] F. Heringhaus, P. Braumann, D. R¨uhlicke, E. Susnik, R. Wolmer, Proc. ICEC 2000, Stockholm, Sweden, 2000, pp. 199–204. [2] S. Mogck, B.J. Kooi, J.T.M. De Hosson, Acta Mater. 52 (2004) 5845–5851. [3] W. Vellinga, Acta Mater. 45 (1997) 933–950. [4] Y. Du, W. Zhang, N. Wang, J. Hu, Faming Zhuanli Shenqing Gongkai Shuomingshu (2001) 7, CN99115496.7. [5] J. Wang, Proceedings of 47th IEEE Holm Conference on Electrical Contacts, 2001, pp. 94–97. [6] G. Elssner, E. Gebhardt, Z. Metallkd. 60 (1969) 922–929.
[7] E.R. Eleite, A.P. Maciel, I.T. Weber, P.N. Lisboa-Filho, E. Longo, C.O. Paiva-Santos, A.V.C. Andrade, C.A. Pakoscimas, Y. Maniette, W.H. Schreiner, Adv. Mater. 14 (2002) 905–908. [8] J.W. Lee, H.C. Lee, Scripta Mater. 42 (1999) 169–173. [9] Y. Niu, F. Gesmundo, M. Al-Omary, J. Song, J. Alloys Compd. 317-318 (2001) 573–577. [10] Materials Preparation Center, Ames Laboratory US-DOE, Ames, IA, USA, http://www.mpc.ameslab.gov. [11] A.J. Nagy, G. Mestl, D. Herein, G. Weinberg, E. Kitzelmann, R. Schl¨ogl, J. Catal. 182 (1999) 417–429.
Microstructure and properties of a silver–erbium oxide alloy D.M. Herman a , G.H. Cao b , A.T. Becker a,b , A.M. Russell a,b,∗ , A.P. Constant a a
b
Materials Science and Engineering Department, 2220 Hoover Hall, Iowa State University, Ames, IA 50011, USA Materials and Engineering Physics Program, 126 Metals Development Building, Ames Laboratory of the USDOE, Ames, IA 50011, USA Received 7 July 2006; accepted 10 December 2006 Available online 16 January 2007
Abstract A Ag–4 wt.% Er solid solution alloy was produced by arc melting and quenching to room temperature. When this alloy was annealed in air at various temperatures, the Er was converted to ∼1-m thick Er2 O3 lamellae in the Ag matrix. Subsequent swaging and stress-relief annealing converted the lamellar microstructure to predominantly spheroidal Er2 O3 particles. The alloys were examined by optical microscopy, scanning electron microscopy (SEM), and transmission electron microscopy (TEM). Electrical resistivities of 2.0–3.3 cm and hardness values of 83–162 HV were measured. © 2007 Elsevier B.V. All rights reserved. Keywords: Metal matrix composites; Rare earth alloys and compounds; Electrical transport; Microstructure; Transmission electron microscopy, TEM
1. Introduction Several composite materials with a silver (Ag) matrix and a dispersed oxide phase have been developed for electrical contact applications. Ag–CdO composites were among the first of these materials, but they suffer from both physical properties limitations and Cd toxicity concerns, which motivated further research into similar materials, such as Ag–SnO2 and Ag–ZnO with ternary additions [1,2]. Much of the current research into silver–metal oxide materials involves optimizing the dispersion of the oxide phase by reducing the particle size, achieving a more uniform distribution [1], and improving the adhesion between the Ag and ceramic interface [3] within the composite. More recently, rare earth materials have been used as a ternary addition [4] to optimize the properties of another oxide [5]. Research into the effectiveness of Ag-rare earth oxide contact materials has been limited [6]. The Ag–Er2 O3 specimens in this study were prepared by arc melting Ag and Er, solutionizing the alloy at high temperature in an inert gas environment, quenching in water, and then annealing in air to internally oxidize the Er to form Er2 O3 second phase
particles in a Ag matrix. This process offers a simpler processing alternative to P/M with the potential to achieve more uniform dispersion of the oxide phase. Internal oxidation has been performed on Ag–Sn alloys [1]. Due to the rather noble nature of Sn, a secondary addition (e.g., In, Te, or Bi) is required to promote the internal oxidation reaction [7,8]. The Ag–Er alloy used in this study did not require ternary alloying elements to achieve oxidation of the Er when the metal was annealed in air. Other investigators have studied Ag-rare earth oxide materials produced by internal oxidation of the rare earth metal [4–6,9]. Reports from these studies have described the resulting microstructure and hardness but not the electrical properties. Elssner and Gebhardt [6] reported hardness values in Ag–0.28 at.% Er that were similar to those measured in this study. He also found that this composition and other dilute rare earth-Ag alloys maintained their high room temperature hardness values at temperatures well above room temperature. For example, the 110 HV hardness of Elssner’s Ag–0.28 at.% Er alloy changed little from room temperature to 800 ◦ C. Above 800 ◦ C the hardness decreased exponentially as the temperature approached the melting temperature. 2. Experimental procedure
∗
Corresponding author at: Materials Science and Engineering Department, 2220 Hoover Hall, Iowa State University, Ames, IA 50011, USA. Tel.: +1 515 294 3204; fax: +1 515 294 5444. E-mail address: [email protected] (A.M. Russell). 0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2006.12.036
To make a 96 wt.% Ag–4 wt.% Er button, 99.99% pure Ag was arc melted in an Ar atmosphere with 99.9% pure Er obtained from the Ames Laboratory Materials Preparation Center [10]. The sample was melted six times to ensure
D.M. Herman et al. / Journal of Alloys and Compounds 454 (2008) 292–296 uniformity. The Ag–Er equilibrium phase diagram indicates that a Ag–4 wt.% Er alloy would form a single-phase solid solution at 750 ◦ C. If cooled slowly enough to room temperature to preserve equilibrium, such an alloy would transform to a more dilute Ag solid solution plus Ag51 Er14 intermetallic compound. To avoid forming the Ag51 Er14 phase, the Ag–Er arc melted button was sealed in a Hefilled, Pyrex ampoule, annealed for 25 h at 750 ◦ C, and submerged in water as the ampoule was broken to quench the metal, forming a metastable, single-phase solid solution. The Ag–4% Er button was cut into four pieces that were internally oxidized at 400, 500, 600, or 700 ◦ C by annealing for 24 h in air, followed by air cooling. The four samples were mounted, polished, and etched with dilute K2 Cr2 O7 –NaCl–H2 SO4 acid solution to reveal the microstructure for optical metallography. SEM was performed on all samples before and after etching. The TEM specimen was prepared from the sample that had been internally oxidized at 500 ◦ C by cutting, grinding, and ion-milling. TEM bright field imaging and selected area electron diffraction (SAED) were performed using a Philips CM30 electron microscope equipped with a LaB6 electron source, operated at an accelerating voltage of 300 kV. Hardness measurements were taken using a LECO microhardness tester with a Vickers indenter. The load was 100 gf with a dwell time of 13 s for all hardness tests. X-ray diffraction was performed on the sample annealed at 700 ◦ C. The bulk sample was used, not powder. A scan of 25 to 75◦ 2θ was performed with a scan rate of 0.004◦ 2θ s−1 . The specimen annealed at 700 ◦ C was placed into a 1.27 cm diameter, thinwalled stainless steel sleeve and swaged at room temperature to a diameter of 2.54 mm. At this point, the stainless steel was removed, and the bare metal was swaged to a final diameter of 0.81 mm. A 30- min stress-relief anneal was performed in air at 300 ◦ C when the specimen diameter was 2.54 mm. It is estimated that this swaging produced a true strain of approximately 5.0. Exact determination of the true strain is not possible, since the starting specimen geometry was not cylindrical. A four-point conductivity test was performed at room temperature on the swaged 700 ◦ C sample. Current was passed between two outer contact points along the cylindrical wire length with a voltage drop being measured between two inner contact points.
3. Experimental results Upon removal from the furnace used for the solutionizing anneal, the samples and the Pyrex ampoule were discolored. The surface of the samples had a pinkish hue (Er2 O3 is pink), leading to the conclusion that the Er in the samples had either reduced some of the oxides in the glass during the anneal or had reacted with the water during the quench, forming a layer of Er2 O3 on the samples. This oxide layer was found to be only superficial, and it was removed before further testing. The microstructure shown in Fig. 1 was observed in the optical micrographs taken on all four samples. This microstructure was dominant and appeared to be a lamellar, rather than acicular, second phase. Individual variations of the spacing between the lamellar microstructure, thickness of the darker phase, and the orientation of the lamellae of this second phase were seen both within individual samples and amongst the four specimens. The microstructure shown in Fig. 1 could be distinguished in both back-scattered electron imaging and secondary electron imaging modes in the SEM. The hardness of all samples measured in this study was much higher than the 25 HV microhardness of pure, annealed Ag. The 400, 500, and 700 ◦ C samples showed similar hardness values (Table 1), but the 600 ◦ C sample showed a lower average hardness with a larger standard deviation than the
293
Fig. 1. Optical micrograph of Ag–Er alloy annealed in air at 400 ◦ C. EDS–SEM examination of this specimen showed that the darker phase contained higher Er content. Table 1 Vickers microhardness data (HV = Vickers microhardness number) Annealing temperature (◦ C)
Average hardness (HV)
Standard deviation (HV)
400 500 600 700
103.6 101.4 83.0 101.6
2.1 5.9 26.9 4.2
other three specimens. Hardness measurements were taken from both transverse and longitudinal sections of the 700 ◦ C sample that had been swaged to a diameter of 2.54 mm before being given a stress-relief anneal. Both types of sections displayed a mixed microstructure with some regions possessing a lamellar microstructure and others spheroidal particles (Fig. 2). The two different microstructures were found to have differing microhardness values (Table 2).
Fig. 2. A 700 ◦ C optical micrograph of the specimen oxidized at 700 ◦ C, swaged, before being annealed at 300 ◦ C showing both the spheroidal phase (centered in the micrograph) and the lamellar phase (near the upper right and lower left regions of the micrograph; lamellae are difficult to resolve at this low magnification).
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Table 2 Vickers microhardness measurements in specimen internally oxidized at 700 ◦ C and swaged to 2.54 mm diameter
Transverse section–spheroidal regions Transverse section–lamellar regions
Average hardness (HV)
Standard deviation (HV)
161.6 103.9
11.2 4.8
Fig. 5. EDS line scan showing higher Er and O content as several lamellae similar to those shown in Fig. 1 are crossed in the specimen annealed at 400 ◦ C.
Fig. 3. Higher magnification optical micrograph of the spheroidal phase structure in Fig. 2.
size larger than about 10 m. It may be that grain growth during the 25 h 750◦ C solutionizing anneal produced a coarse grain structure in the single-phase solid solution Ag–Er alloy that was retained at room temperature.
In the transverse cross-section of the sample oxidized at 700 ◦ C and subsequently swaged before being annealed at 300◦ C, the spheroidal microstructure dominated, although the lamellar microstructure was visible near the rod’s surface. Energy dispersive spectroscopy (EDS) showed that lamellae (Fig. 1) and spheroidal particles (Figs. 2 and 3) correspond to lower concentrations of Ag and higher concentrations of Er and O (Figs. 4 and 5). TEM of the sample annealed at 500 ◦ C showed the size, crystal structure, and morphology of the Er2 O3 particles (Figs. 6–8). The X-ray diffraction data from the sample annealed at 700 ◦ C were compared to published powder diffraction files of Ag, Er2 O3 , and Ag51 Er14 . The XRD data matched the peaks of Ag and Er2 O3 ; no peaks from the Ag51 Er14 intermetallic compound were observed. The complex microstructure in these specimens made grain size determination difficult. Bright field TEM showed no grain boundaries in the thinned region of the foil, suggesting a grain
Fig. 4. EDS line scan of spheroidal microstructure in the specimen annealed at 700 ◦ C (Fig. 3). The scan crossed a spheroid from 2 to 8 m on the abscissa.
Fig. 6. TEM images of (a) two Er2 O3 particles (A and B) and the matrix phase, (b) lamellar Er2 O3 phase.
D.M. Herman et al. / Journal of Alloys and Compounds 454 (2008) 292–296
Fig. 7. SAED patterns of particle A in Fig. 6 along (a) [0 0 1], (b) [0 1 1], and (c) [1 0 3] zone axes. Indexing shows that particle A has cubic crystal structure ˚ which matches that of Er2 O3 . (d) SAED pattern of the particle with a0 = 10.5 A, B along [1 1 2] zone axis, which also indexes as Er2 O3 .
295
defects in the crystal. At lower temperatures, O diffuses in the bulk Ag via an interstitial mechanism through low resistance diffusion paths such as grain boundaries and the free surface. At higher temperatures atomic O begins to segregate onto low indexed planes. It is possible that this process occurs due to an interstitial diffusion mechanism where O substitutes for Ag atoms in the lattice. Such a planar distribution of O atoms could be the precursor to Er2 O3 formation on those same planes, yielding the lamellar microstructure shown in Fig. 1. The change to a predominantly spheroidal microstructure after swaging and annealing is apparently the result of surface energy minimization in the previously existing lamellar structure that accompanied recrystallization in the heavily cold-worked Ag matrix. It was somewhat surprising that the microstructures and hardness values of the four specimens annealed in air at 400, 500, 600, and 700 ◦ C were all so similar. In all four of these specimens, there appeared to be sufficient diffusion to segregate the oxide phase into lamellae with micron-scale dimensions. This seems to suggest that the annealing temperature that would produce the finest possible oxide phase size (and hence, the highest hardness) may lie below 300 ◦ C, a processing option that was not explored in this study. 5. Conclusions
Fig. 8. SAED patterns of the matrix Ag along (a) [0 0 1], (b) [0 1 1], and (c) [−1 1 2] zone axes. Indexing shows that the matrix has a lattice parameter ˚ which matches that of Ag. a0 = 4.2 A,
A four-point conductivity measurement was taken on the wire, which showed a resistivity of ∼3.3 cm, which is substantially higher than the 1.59 cm resistivity of pure, annealed Ag. After the wire had been annealed at 400 ◦ C for 30 min, the same test showed a resistivity of ∼2.0 cm (∼86% IACS). The resistivity values are approximate since some exfoliation occurred on the wire surface during swaging, giving a wire geometry that was not perfectly cylindrical. Since the spalling slightly reduces the cross-sectional area for current flow, it is likely that the specimen’s actual resistivity is slightly lower than the values measured in the four-point conductivity test.
The Ag–Er2 O3 material of this study possesses physical properties that would make it a useful material for electrical contacts. It was seen in previous studies [6] that rare earth concentrations under 1 wt.% could still create a significant improvement to the hardness of Ag. Since lower Er concentrations would be expected to give higher electrical and thermal conductivity, a lower concentration would be preferable. Although the highpurity Er used in this project is more costly than Ag, the costs of commercial-purity rare earth metals are roughly similar to (and often lower than) that of Ag. Additionally internal oxidation is a simpler processing method than powder metallurgy. Current methods of internal oxidation require ternary alloy additions to promote diffusion and internal oxidation of the primary oxide formers, and high O partial pressures are usually used to promote oxidation. These two difficulties to the use of internal oxidation could be eliminated if a more reactive second metal such as a rare earth metal were used. Although a comparative test of actual switching performance would be necessary to determine how well existing Ag–metal oxide materials compare to Ag-rare earth oxides, this study shows that rare earth oxides have the potential to improve the switch and contact performance of Ag-based contact materials.
4. Discussion Acknowledgements The lamellar microstructure (Fig. 1) could be the result of preferential diffusion mechanisms for oxygen (O) in Ag. It was observed by Nagy et al. [11] that preferential diffusion directions exist for O movement in the Ag lattice and among
The authors gratefully acknowledge the assistance of J. Peters in performing XRD, S. Williams for TEM sample preparation, and B. Stumphy for solutionizing the samples.
296
D.M. Herman et al. / Journal of Alloys and Compounds 454 (2008) 292–296
References [1] F. Heringhaus, P. Braumann, D. R¨uhlicke, E. Susnik, R. Wolmer, Proc. ICEC 2000, Stockholm, Sweden, 2000, pp. 199–204. [2] S. Mogck, B.J. Kooi, J.T.M. De Hosson, Acta Mater. 52 (2004) 5845–5851. [3] W. Vellinga, Acta Mater. 45 (1997) 933–950. [4] Y. Du, W. Zhang, N. Wang, J. Hu, Faming Zhuanli Shenqing Gongkai Shuomingshu (2001) 7, CN99115496.7. [5] J. Wang, Proceedings of 47th IEEE Holm Conference on Electrical Contacts, 2001, pp. 94–97. [6] G. Elssner, E. Gebhardt, Z. Metallkd. 60 (1969) 922–929.
[7] E.R. Eleite, A.P. Maciel, I.T. Weber, P.N. Lisboa-Filho, E. Longo, C.O. Paiva-Santos, A.V.C. Andrade, C.A. Pakoscimas, Y. Maniette, W.H. Schreiner, Adv. Mater. 14 (2002) 905–908. [8] J.W. Lee, H.C. Lee, Scripta Mater. 42 (1999) 169–173. [9] Y. Niu, F. Gesmundo, M. Al-Omary, J. Song, J. Alloys Compd. 317-318 (2001) 573–577. [10] Materials Preparation Center, Ames Laboratory US-DOE, Ames, IA, USA, http://www.mpc.ameslab.gov. [11] A.J. Nagy, G. Mestl, D. Herein, G. Weinberg, E. Kitzelmann, R. Schl¨ogl, J. Catal. 182 (1999) 417–429.