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Oxidation and contact resistance of Sn–Ag coated superconducting strands for the Large Hadron Collider (LHC)
Applied Surface Science 253 (2006) 1393–1398 www.elsevier.com/locate/apsusc
Oxidation and contact resistance of Sn–Ag coated superconducting strands for the Large Hadron Collider (LHC) C. Scheuerlein a,*, M. Taborelli a, M. Cantoni b a
European Laboratory for Nuclear Research (CERN), CH-1211, Geneva 23, Switzerland b Ecole Polytechnique Fe´de´rale de Lausanne, Switzerland
Received 20 January 2006; received in revised form 7 February 2006; accepted 7 February 2006 Available online 4 April 2006
Abstract The oxides formed on the Sn–Ag coated Large Hadron Collider (LHC) superconducting cables during a 200 8C heat treatment in air are described and the oxide composition is compared with the interstrand contact resistance (RC). The analysis of more than 250 interstrand contact areas shows that the higher the average Cu content with respect to the Sn content in the oxide, the higher is RC. During the 200 8C heat treatment, Sn in the coating is transformed into a Cu3Sn layer, on which an oxide grows that consists essentially of a thin outermost layer of CuO on top of Cu2O, similar to the oxide structure formed on bare Cu. The underlying Cu3Sn layer acts as an O diffusion barrier that prevents O diffusion into the Cu bulk during the subsequent cable heat treatment under high pressure. On contact zones where the Cu3Sn layer is not formed during the 200 8C heat treatment mainly Sn oxide grows and RC is comparatively low. # 2006 Elsevier B.V. All rights reserved. Keywords: Contact resistance (73.40.Cg); Oxidation (81.65.Mq); Electron spectroscopy (82.80.Pv); Large Hadron Collider (LHC)
1. Introduction In order to reduce coupling currents between the crossing strands of the Large Hadron Collider (LHC) superconducting cables, the interstrand contact resistance (RC) must be controlled (for the LHC inner and outer dipole cables RC must exceed 15 and 40 mV, respectively) [1]. For this purpose a Sn–Ag coating with an average thickness of about 0.5 mm is applied onto the copper matrix of the strands [2,3] and the required RC values are obtained through a 200 8C cable heat treatment in air, lasting typically a few hours. During the 200 8C cable heat treatment Sn and Cu interdiffuse and form the intermetallic compounds Cu6Sn5 and Cu3Sn. In a previous article the influence of the CuSn intermetallics on RC has been discussed [4]. The goal of the present study is to better understand the influence of the surface oxides on RC of the heat treated LHC cables. The oxides present on the cable surface after the heat treatment in air can vary around the circumference of a single strand and even more for the different strands within one cable,
* Corresponding author. Tel.: +41 22 767 8829; fax: +41 22 767 6300. E-mail address: [email protected] (C. Scheuerlein). 0169-4332/$ – see front matter # 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2006.02.015
which can be easily seen from the varying strand colours. Therefore, a surface analysis measurement at a single interstrand contact zone is unlikely to be representative for the average oxide composition on the cable surface. On the other hand, RC values obtained for the LHC cable samples are influenced by several hundred interstrand contact points. In order to obtain an average oxide composition that is representative for a cable, in the present study at least 15 strands of a certain cable sample have been analysed by Auger electron spectroscopy (AES) and the average Sn and Cu content in the strand oxide has been determined. In addition, the outer oxide surface was analysed by X-ray photoelectron spectroscopy (XPS), and oxide cross sections have been characterised by energy dispersive X-ray spectroscopy (EDS) in a scanning transmission electron microscope (STEM), using electron transparent TEM lamellas prepared by focused ion beam (FIB). 2. Experimental 2.1. The samples All samples analysed were extracted from outer LHC dipole (type 2) cables. The cables are made of 36 strands with a
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diameter of 0.825 mm. The Cu strand matrix is coated by a Sn– Ag alloy with an average thickness of around 0.5 mm. During the continuous hot dip coating process a coating structure consisting of Cu3Sn, Cu6Sn5 and Sn–Ag alloy is obtained. All coated cable samples are oxidised in air at 200 8C during typically 8 h, prior to RC measurements. During this heat treatment Sn and Cu interdiffusion continues. If the 200 8C heat treatment is long enough the entire Sn coating is transformed into an approximately 2 mm thick Cu3Sn layer. After the 200 8C heat treatment the cable samples are mounted in a sample holder under a pressure of 50 MPa. This pressure simulates the actual stress to which the cable is subjected within a manufactured LHC magnet. For the RC measurement the sample holder is immersed in a liquid helium bath. After a first RC measurement of the as-received cable sample, the sample holder is removed from the cryostat for a 30 min heat treatment in air at a temperature of 190 8C, without releasing the applied pressure. This procedure intends to simulate the curing procedure in the LHC magnet coils. All RC values as well the surface analysis results presented in this article have been obtained subsequent to this 190 8C heat treatment. Typically, RC values for Sn–Ag coated LHC cables decrease during this heat treatment by a factor of 3–4.
Fig. 1. AES survey spectra acquired on a contact zone on which the oxide contains mainly Sn (upper spectrum) or Cu (lower spectrum). The respective CuLMM/SnMNN peak area ratios are 0.23 and 11.
After oxidation at 200 8C in air the strand oxide thickness is in the order of 20 nm [4] and it is assumed that the AES probing depth does not extend into the underlying metal. It is further assumed that the varying degree of contamination that covers the strand oxides does not strongly influence the relation between CuLMM/SnMNN and RC. Elements other than Cu and Sn are not considered for the determination of the Sn content.
2.2. Auger electron spectroscopy
2.3. Surface analysis by X-ray photoelectron spectroscopy
AES has been chosen for the analysis of the interstrand contact surfaces because of its better spatial resolution and shorter acquisition times as compared to X-ray photoelectron spectroscopy. Up to 18 strand samples, which are extracted from the same cable sample, are fixed by electrically conductive double-sided adhesive tape onto the sample holder that is inserted into the UHV system of the Auger electron spectrometer. In addition a Cu wire is brought into contact close the analysis area of each strand, in order to reduce sample charging during the AES measurements. AES measurements are performed with an Auger electron spectrometer consisting of a single-pass cylindrical mirror analyser (CMA) PHI 15-110B, from Physical Electronics, with a coaxial electron gun. The spectra are acquired with a relative energy resolution DE/E of 1.2% (full width at half-maximum). The primary electron (PE) energy is 5 keV and the primary beam current is typically 0.4 mA, incident normal to the sample surface. The lateral resolution is about 30 mm. The ratio of the Cu-LMM to Sn-MNN areas (CuLMM/ SnMNN) without any normalisation is used in order to compare the Sn content in the strand oxide. Peak areas are calculated from the direct spectra by subtracting, as a crude approximation, a linear background (between 410–440 and 875–950 eV for Sn-MNN and Cu-LMM peak areas, respectively). Fig. 1 shows two spectra acquired on strand oxide layers that are mainly composed of Sn oxide and of Cu oxide, respectively. The maximum CuLMM/SnMNN value has been set to 20. For higher ratios the Sn-MNN peak approaches the noise level and cannot be measured accurately. In the present study the CuLMM peak was clearly measurable on all samples.
The chemical state of the outermost few nanometers of the strand oxides was analysed by XPS, using a PHI ESCA 5400 instrument. The contact zones of the extracted wire samples were analysed in the as-received state, i.e. without in situ sputter cleaning. The analysed sample area has a diameter of about 1 mm2 (aperture #2) and the electron take-off angle is 458. Photoelectrons are excited by a non-monochromatic Mg Ka X-ray source operating at 350 W/15 kV. The PHI model 10-360 spherical sector electron spectrometer was operated with fixed pass energy of 89.45 eV for survey spectra and of 35.75 eV for high resolution spectra. The energy scale has been calibrated to the position of the Cu 2p3/2 and Au 4f7/2 peaks on sputter-cleaned samples. Published Cu 2p3/2 reference peak positions [5] are used for chemical state analysis (see Table 1). SnO and SnO2 cannot Table 1 Cu 2p3/2 binding energy and Cu-LMM kinetic energy for different Cu chemical states Chemical state
Binding energy (eV) Cu 2p3/2
Kinetic energy (eV) Cu-LMM
Cu metal CuO Cu2O Cu(OH)2 Cu3Sn [17]
932.6 933.7 932.5 935.1 933.2
918.6 918.1 916.2 916.2 918.1
Auger parameter (eV) 14.0 15.6 16.3 18.9 15.1
All peak energies are from Ref. [5], except the values for Cu3Sn. The Auger parameter is defined as the difference between the Auger electron kinetic energy and the photoelectron binding energy [16].
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easily be distinguished by XPS and AES and no attempt has been made to determine the Sn oxidation state. 2.4. Oxide bulk analysis by energy dispersive X-ray spectroscopy in a scanning transmission electron microscope Intermetallic compound layer and oxide bulk cross sections were prepared by focused ion beam. The electron transparent FIB lamellas have a thickness of about 100 nm. For more details of the FIB lamella preparation see Ref. [4]. An accurate quantification of the intermetallic composition and a semi-quantitative determination of the bulk oxide composition have been obtained by EDS in a STEM (Philips CM300UT).
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The RC value of 11 mV for the B sample is exceptionally low for cables produced by the same manufacturer and the chemical oxide composition of this sample differs strongly from that of other B samples in that it contains a large amount of Cl throughout the oxide layer. This Cl may be left from the flux that is used to activate the strand surface before the Sn–Ag deposition and may cause a reduction in RC. The RC value of 437 mV obtained for sample C is exceptionally high (typically RC values do not exceed 200 mV). The reason for this exceptionally high RC value is not understood. 3.2. Oxidation state of Cu oxides
3.1. Average Cu content in the oxide and interstrand contact resistance
Since the AES measurements at the energy resolution used in the present study do not provide information about the chemical state, the outermost oxide has been characterised by XPS. An indication about the bulk composition of the oxides present on different samples has been obtained using electron transparent oxide cross sections that were characterised by EDS in a STEM.
In Fig. 2 the average CuLMM/SnMNN values for a certain cable sample are plotted as a function of the corresponding RC values that were obtained for the same cable. It can be seen that there is a clear trend of the CuLMM/SnMNN ratio to increase with RC, with the exception of two out of the total of 16 cable samples, which are indicated by empty symbols. Excluding these two samples, which are considered to be outliers, a tentative linear correlation between the average Cu-LMM to Sn-MNN peak area ratio and the cable RC is fitted. The linear correlation coefficient is 0.77 when the outliers are not considered (when taking into account the two samples considered as outliers the correlation coefficient would be 0.22). The two samples considered as outliers are not taken into account for the following reasons.
3.2.1. Oxide thickness and bulk chemical composition as determined by EDS A STEM image of the oxide layer cross section on top of a continuous Cu3Sn coating after a 200 8C heat treatment in air is shown in Fig. 3. The oxide with a thickness of typically 20 nm is embedded between the Cu3Sn layer and a protective Pt layer, which is deposited onto the sample in order to protect the oxide during the preparation of the FIB lamella. The average composition of the intermetallic layer as determined by EDS is 75.1 at.% Cu and 24.9 at.% Sn. Because of the small oxide thickness an influence of the Cu3Sn layer on the oxide result cannot be excluded entirely and the obtained values are therefore considered as semi-quantitative results. For 10 EDS measurements, which gave a Sn content
Fig. 2. Average Cu-LMM to Sn-MNN peak area ratio as a function of interstrand contact resistance (RC) in LHC cables. In total the plot represents more than 250 Auger electron spectra that have been acquired on different interstrand contact zones. Cable samples of different manufacturers are indicated by different symbols. The two extreme points indicated by empty symbols are considered as outliers and not taken into account for the calculation of the correlation coefficient.
Fig. 3. STEM dark field image of a cross section through the Cu3Sn intermetalic layer and the strand oxide layer, which are formed during the 200 8C cable heat treatment in air. In the image the oxide layer with a thickness of about 20 nm is embedded between the Cu3Sn layer and a Pt layer that has been deposited onto the oxide prior the FIB lamella preparation.
3. Results
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Fig. 4. STEM dark field image of a fractured intermetallic layer and Sn La1, Cu Ka1 and O Ka1 chemical maps of the corresponding oxide cross section. It can be seen that Cu is absent in the oxide layers.
below 1 at.%, the average oxide composition is 30.3 at.% O and 69.5 at.% Cu. The Cu/O ratio of 2.3 indicates that the oxide bulk on top of the Cu3Sn layer is essentially composed of Cu2O rather than CuO. The fact that the ratio is higher than 2 is accounted to the proximity of Cu from the oxide substrate. In contrast to the oxide that it is observed on the continuous Cu3Sn layer, the oxide that is formed on a fractured intermetallic layer contains mainly Sn while almost no Cu is detected (see chemical maps in Fig. 4).
analysis shows that after room temperature oxidation the oxide consists of about 9% Sn oxide and 91% Cu oxide (weighted Cu2p/ Sn3d photoelectron ratio is 11.4). The Sn oxide content may be slightly underestimated because some of the detected Cu is metallic, which may indicate a contribution of the Cu substrate to the Cu signal. All Sn that is detected by XPS is oxidised. After the 200 8C air heating only Cu oxide can be detected by XPS, i.e. after the 200 8C heat treatment the Sn content in the oxide is below the XPS detection limit of about 1 at.%.
3.2.2. Outermost surface oxide composition as determined by XPS A typical photoelectron survey spectrum acquired on the oxide that is formed during the 200 8C heat treatment in air on top of a continuous Cu3Sn layer is shown in Fig. 5. The surface is mainly composed of Cu, O, some Sn and C from hydrocarbon surface contamination (the samples are not in situ sputtercleaned in order to avoid oxide reduction). The Cu chemical state at the outermost oxide surface is found from the high resolution photoelectron spectrum. From the Cu 2p3/2 energy position at 933.8 eV and the characteristic shape of the satellite peak at around 943 eV it can be concluded that the outermost Cu oxide is mainly composed of CuII as in CuO (see Table 1).
4. Discussion
3.2.3. Oxidation of Cu–8 at.% Sn bronze in air In order to understand what are the rate controlling steps in the oxidation of Cu–Sn alloys, a sputter-cleaned Cu–8 at.% Sn bronze sample was oxidised at 200 8C in air. XPS surface
4.1. The formation of Cu oxides on the LHC strands The AES analysis of more than 250 interstrand contact zones in LHC cables shows that Cu oxides are commonly present after the 200 8C heat treatment in air. In many cases only Cu oxide is detected, i.e. the Sn content in the oxide is below 1 at.%, which is the approximate detection limit of AES. In this case EDS analysis indicates that the oxide bulk is mainly composed of Cu2O and XPS reveals that the outer oxide surface is essentially composed of CuO. 4.1.1. Thermodynamics of oxide growth In Ref. [6] the formation of Cu oxides during heat treatments of Sn coated Cu has been ascribed to the formation of an intermetallic Cu3Sn layer, while on bronze or Cu6Sn5 mainly Sn oxide has been found. This observation is in agreement with the EDS/STEM results presented in this article, which detected only Cu oxide on a continuous Cu3Sn layer, while mainly Sn oxide is found when the coating is not fully reacted or when the intermetallic layer is absent. As shown in Table 2 the free energy of formation (DG0) for the stable Sn oxides (SnO and SnO2) is significantly higher than DG0 for the Cu oxides CuO and Cu2O [7]. Hence, thermodynamically the growth of Sn oxide should be favoured over the formation of Cu oxides, also on Cu3Sn. Indeed an XPS study [17] shows that the oxide formed on Cu3Sn bulk at 200 8C Table 2 Free energy of formation per mole of oxygen consumed for more stable tin and copper oxides [7] Reaction
Fig. 5. XPS survey spectrum and high resolution spectrum of the Cu 2p peak acquired on an interstrand contact after 200 8C cable heat treatment in air. The Cu 2p3/2 energy position at 933.8 and the shape of the satellite peak at around 943 eV show that the outermost Cu oxide is mainly in the form of CuO.
4Cu + O2 = 2Cu2O 2Cu + O2 = 2CuO 2Sn + O2 = 2SnO Sn + O2 = SnO2
DG0 (kJ/mol) 292.9 254.6 515.0 515.8
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in a pure O2 atmosphere at low pressure (4.3 mbar) is essentially composed of SnO and SnO2, as predicted by the calculated Cu–Sn–O phase diagram at 200 8C [17]. It can thus be concluded that under equilibrium conditions Cu2O and CuO are not formed on Cu3Sn and their presence must be due to kinetic effects.
the oxide–substrate interface and in the oxide layer, it is concluded that the presence of Ag in the strand coating does not have a strong influence on the strand oxidation behaviour at 200 8C in air.
4.1.2. Kinetics of oxide growth Both, copper(I) oxide Cu2O and copper(II) oxide CuO are semiconductors [8]. Cu2O is a cation deficient p-semiconductor, i.e. metal cations can diffuse along metal ion vacancies in the oxide towards the oxide/gas phase interface, where the oxidation process can proceed. In contrast to Cu2O, SnO2 is an anion deficient n-semiconducting oxide. In this case the oxide is oxygen deficient, which allows oxygen anions to diffuse through the oxide towards the metal/oxide interface, at which the oxide growth will proceed [9]. The reason for the different oxides that are formed on Cu3Sn at low O2 pressure and at atmospheric pressure in air may therefore be attributed to the different diffusion rates of Cu cations in Cu2O and oxygen anions in SnO2. At high O2 pressure (in air) the oxide growth is determined by the higher Cu cation diffusion rate with respect to the oxygen anion diffusion. This is confirmed by the fact that the Sn oxide vanishes upon 200 8C oxidation of Cu–8 at.% Sn bronze in air. At low O2 pressure the oxide growth is determined by the thermodynamics. In this case the oxygen absorption and diffusion towards the interface is oxidation rate controlling.
The formation of a continuous Cu3Sn layer during the cable heat treatment is crucial to obtain the desired RC values [4]. EDS and XPS analysis reveal that the oxide formed during the 200 8C heat treatment at atmospheric pressure in air on Cu3Sn consists essentially of Cu2O and a thin outer layer of CuO. This oxide structure is similar to that formed on bare Cu that is oxidised at elevated temperatures in air [13]. Similarly to oxidised Cu3Sn, RC of oxidised Cu is comparatively high, but in this case of bare Cu strands RC drops dramatically during heat treatments that are performed under pressure, simulating the curing process in LHC magnets [14]. Therefore, the Cu3Sn layer may act as a diffusion barrier for oxygen ions, stabilising the Cu oxide during the curing treatment [15]. The data presented in Fig. 2 shows a clear relation between interstrand contact resistance and the oxide type, i.e. the more Cu oxide is present on the strand surfaces the higher is the interstrand contact resistance. The reason for this might be a higher electrical resistance of the p-semiconducting Cu oxide with respect to the n-semiconducting Sn oxide, and/or differences in oxide thickness. Another explanation may be that the formation of Cu oxide during the cable heat treatment indicates that a continuous Cu3Sn layer has been formed, which increases the oxide stability.
4.1.3. Reasons for the oxide inhomogeneity The main reasons why different oxides are formed on the strands during identical heat treatments are irregularities of the underlying strand coatings. These irregularities can for instance be the result of insufficient wetting of the Cu substrate by the Sn–Ag alloy, remaining flux that is applied prior to the hot dip coating process or because of differences in the coating temperature and/or cooling rates, which influence the coating microstructure [10]. Other elements in the coating bath, apart from Cu and Ag, can also alter the coating chemistry and microstructure [11], and, as a result, influence the surface oxide structure and thickness that grows during the 200 8C air heating treatment. 4.1.4. The influence of Ag in the coating on the surface oxidation The Ag in the Sn–3.5 wt.% Ag solder is present in the form of intermetallic Ag3Sn particles, which increase the mechanical strength of the non-reacted coating. EDS analysis of reacted coating cross sections shows that also in this case Ag is in the form of Ag3Sn. At lower temperatures (up to 170 8C) the Ag content can influence the Cu3Sn growth rate while at 205 8C no significant difference in the Cu3Sn growth rate with respect to pure Sn–Cu diffusion couples has been observed [12]. In the present study Ag was never detected at the coating surface or in the oxide layer. Since the presence of Ag does not significantly influence the intermetallics growth rate and because no Ag can be detected at
4.2. Strand oxide composition and RC
5. Conclusion The interstrand contact resistance in the LHC superconducting cables is related to the oxide composition, i.e. the more Cu oxide relative to the amount of Sn oxide the higher is the contact resistance. Despite the fact that for stable Sn oxides the free energy of formation is higher than for the stable Cu oxides, on many strand samples only Cu oxide is detected. Cu oxide formation depends critically on the formation of a continuous Cu3Sn layer. The Cu oxide formed during the 200 8C heat treatment in air has the same structure as the oxide formed on bare Cu (thin CuO layer on top of Cu2O bulk). The fact that RC on the Sn coated strand Cu matrix remains sufficiently high during curing, while it is drastically reduced on the un-coated Cu indicates that the Cu3Sn layer is stabilising the Cu oxide during the heat treatment under high pressure. Acknowledgements We are grateful to D. Richter and A. Bastos Marzal for RC results and advice about interstrand contact resistance measurements.
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References [1] O. Bru¨ning, P. Collier, P. Lebrun, S. Myers, R. Ostojic, J. Poole, P. Proudlock (Eds.), LHC Design Report, vol. 1, The LHC Main Ring, CERN-2004-003, 2004. [2] J.-M. Depond, D. Leroy, L.R. Oberli, D. Richter, IEEE Trans. Appl. Supercond. 7 (2) (1997) 793. [3] D. Richter, J.D. Adam, J.-M. Depond, D. Leroy, L.R. Oberli, IEEE Trans. Appl. Supercond. 7 (2) (1997) 786. [4] C. Scheuerlein, Ph. Gasser, P. Jacob, D. Leroy, L. Oberli, M. Taborelli, J. Appl. Phys. 97 (3) (2005). [5] J.F. Moulder, W.F. Stickle, P.E. Sobol, K.D. Bomben, Handbook of X-ray .Photoelectron Spectroscopy, Physical Electronics Inc., Eden Prairie, 1995 [6] F. Bartels, J.W. Morris, G. Dalke, W. Gust, Intermetallic phase formation in thin solid-liquid diffusion couples, J. Electron. Mater. 23 (8) (1994) 787. [7] D.M. Tench, Oxidation of solder coatings, in: F.G. Yost, F.M. Hosking, D.R. Frear (Eds.), The Mechanics of Solder Alloy Wetting and Spreading, Van Nostrand Reinhold, New York, 1993. [8] T.J. Richardson, J.L. Slack, M.D. Rubin, Electrochromism in copper oxide thin films, Electrochim. Acta 46 (13–14) (2001) 2281.
[9] S.A. Bradford, Fundamentals of corrosion in gases, American Society for Metals Handbook, vol. 13, Corrosion, ninth ed., 1987, pp. 62–76. [10] F. Ochoa, J.J. Williams, N. Chawla, J. Electron. Mater. 32 (12) (2003) 1414. [11] I.E. Anderson, B.A. Cook, J. Harringa, R.L. Terpstra, J. Electron. Mater. 31 (11) (2002) 1166. [12] P.T. Vianco, K.L. Erickson, P.L. Hopkins, J. Electron. Mater. 23 (8) (1994) 721. [13] H. Leidheiser, The Corrosion of Copper, Tin and their Alloys, Wiley, New York, 1971. [14] A. Ghosh, Interstrand resistance studies of oxidised superconducting wires, SSC Technical Note No. 103, SSCL-N-808, 1992. [15] D. Richter, J.D. Adam, D. Leroy, L.R. Oberli, IEEE Trans. Appl. Supercond. 9 (2) (1999) 735. [16] S.D. Waddington, Auger chemical shifts and the Auger parameter, in: D. Briggs, M.P. Seah (Eds.), Practical Surface Analysis, vol. 1, Auger- and Xray Photoelectron Spectroscopy, John Wiley & Sons, New York, 1995, p. 588. [17] D.N. Wang, A.C. Miller, M.R. Notis, XPS study of the oxidation behaviour of the Cu3Sn intermetallic compound at low temperatures, Surf. Interface Anal. 24 (2) (1998) 127–132.
Oxidation and contact resistance of Sn–Ag coated superconducting strands for the Large Hadron Collider (LHC) C. Scheuerlein a,*, M. Taborelli a, M. Cantoni b a
European Laboratory for Nuclear Research (CERN), CH-1211, Geneva 23, Switzerland b Ecole Polytechnique Fe´de´rale de Lausanne, Switzerland
Received 20 January 2006; received in revised form 7 February 2006; accepted 7 February 2006 Available online 4 April 2006
Abstract The oxides formed on the Sn–Ag coated Large Hadron Collider (LHC) superconducting cables during a 200 8C heat treatment in air are described and the oxide composition is compared with the interstrand contact resistance (RC). The analysis of more than 250 interstrand contact areas shows that the higher the average Cu content with respect to the Sn content in the oxide, the higher is RC. During the 200 8C heat treatment, Sn in the coating is transformed into a Cu3Sn layer, on which an oxide grows that consists essentially of a thin outermost layer of CuO on top of Cu2O, similar to the oxide structure formed on bare Cu. The underlying Cu3Sn layer acts as an O diffusion barrier that prevents O diffusion into the Cu bulk during the subsequent cable heat treatment under high pressure. On contact zones where the Cu3Sn layer is not formed during the 200 8C heat treatment mainly Sn oxide grows and RC is comparatively low. # 2006 Elsevier B.V. All rights reserved. Keywords: Contact resistance (73.40.Cg); Oxidation (81.65.Mq); Electron spectroscopy (82.80.Pv); Large Hadron Collider (LHC)
1. Introduction In order to reduce coupling currents between the crossing strands of the Large Hadron Collider (LHC) superconducting cables, the interstrand contact resistance (RC) must be controlled (for the LHC inner and outer dipole cables RC must exceed 15 and 40 mV, respectively) [1]. For this purpose a Sn–Ag coating with an average thickness of about 0.5 mm is applied onto the copper matrix of the strands [2,3] and the required RC values are obtained through a 200 8C cable heat treatment in air, lasting typically a few hours. During the 200 8C cable heat treatment Sn and Cu interdiffuse and form the intermetallic compounds Cu6Sn5 and Cu3Sn. In a previous article the influence of the CuSn intermetallics on RC has been discussed [4]. The goal of the present study is to better understand the influence of the surface oxides on RC of the heat treated LHC cables. The oxides present on the cable surface after the heat treatment in air can vary around the circumference of a single strand and even more for the different strands within one cable,
* Corresponding author. Tel.: +41 22 767 8829; fax: +41 22 767 6300. E-mail address: [email protected] (C. Scheuerlein). 0169-4332/$ – see front matter # 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2006.02.015
which can be easily seen from the varying strand colours. Therefore, a surface analysis measurement at a single interstrand contact zone is unlikely to be representative for the average oxide composition on the cable surface. On the other hand, RC values obtained for the LHC cable samples are influenced by several hundred interstrand contact points. In order to obtain an average oxide composition that is representative for a cable, in the present study at least 15 strands of a certain cable sample have been analysed by Auger electron spectroscopy (AES) and the average Sn and Cu content in the strand oxide has been determined. In addition, the outer oxide surface was analysed by X-ray photoelectron spectroscopy (XPS), and oxide cross sections have been characterised by energy dispersive X-ray spectroscopy (EDS) in a scanning transmission electron microscope (STEM), using electron transparent TEM lamellas prepared by focused ion beam (FIB). 2. Experimental 2.1. The samples All samples analysed were extracted from outer LHC dipole (type 2) cables. The cables are made of 36 strands with a
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diameter of 0.825 mm. The Cu strand matrix is coated by a Sn– Ag alloy with an average thickness of around 0.5 mm. During the continuous hot dip coating process a coating structure consisting of Cu3Sn, Cu6Sn5 and Sn–Ag alloy is obtained. All coated cable samples are oxidised in air at 200 8C during typically 8 h, prior to RC measurements. During this heat treatment Sn and Cu interdiffusion continues. If the 200 8C heat treatment is long enough the entire Sn coating is transformed into an approximately 2 mm thick Cu3Sn layer. After the 200 8C heat treatment the cable samples are mounted in a sample holder under a pressure of 50 MPa. This pressure simulates the actual stress to which the cable is subjected within a manufactured LHC magnet. For the RC measurement the sample holder is immersed in a liquid helium bath. After a first RC measurement of the as-received cable sample, the sample holder is removed from the cryostat for a 30 min heat treatment in air at a temperature of 190 8C, without releasing the applied pressure. This procedure intends to simulate the curing procedure in the LHC magnet coils. All RC values as well the surface analysis results presented in this article have been obtained subsequent to this 190 8C heat treatment. Typically, RC values for Sn–Ag coated LHC cables decrease during this heat treatment by a factor of 3–4.
Fig. 1. AES survey spectra acquired on a contact zone on which the oxide contains mainly Sn (upper spectrum) or Cu (lower spectrum). The respective CuLMM/SnMNN peak area ratios are 0.23 and 11.
After oxidation at 200 8C in air the strand oxide thickness is in the order of 20 nm [4] and it is assumed that the AES probing depth does not extend into the underlying metal. It is further assumed that the varying degree of contamination that covers the strand oxides does not strongly influence the relation between CuLMM/SnMNN and RC. Elements other than Cu and Sn are not considered for the determination of the Sn content.
2.2. Auger electron spectroscopy
2.3. Surface analysis by X-ray photoelectron spectroscopy
AES has been chosen for the analysis of the interstrand contact surfaces because of its better spatial resolution and shorter acquisition times as compared to X-ray photoelectron spectroscopy. Up to 18 strand samples, which are extracted from the same cable sample, are fixed by electrically conductive double-sided adhesive tape onto the sample holder that is inserted into the UHV system of the Auger electron spectrometer. In addition a Cu wire is brought into contact close the analysis area of each strand, in order to reduce sample charging during the AES measurements. AES measurements are performed with an Auger electron spectrometer consisting of a single-pass cylindrical mirror analyser (CMA) PHI 15-110B, from Physical Electronics, with a coaxial electron gun. The spectra are acquired with a relative energy resolution DE/E of 1.2% (full width at half-maximum). The primary electron (PE) energy is 5 keV and the primary beam current is typically 0.4 mA, incident normal to the sample surface. The lateral resolution is about 30 mm. The ratio of the Cu-LMM to Sn-MNN areas (CuLMM/ SnMNN) without any normalisation is used in order to compare the Sn content in the strand oxide. Peak areas are calculated from the direct spectra by subtracting, as a crude approximation, a linear background (between 410–440 and 875–950 eV for Sn-MNN and Cu-LMM peak areas, respectively). Fig. 1 shows two spectra acquired on strand oxide layers that are mainly composed of Sn oxide and of Cu oxide, respectively. The maximum CuLMM/SnMNN value has been set to 20. For higher ratios the Sn-MNN peak approaches the noise level and cannot be measured accurately. In the present study the CuLMM peak was clearly measurable on all samples.
The chemical state of the outermost few nanometers of the strand oxides was analysed by XPS, using a PHI ESCA 5400 instrument. The contact zones of the extracted wire samples were analysed in the as-received state, i.e. without in situ sputter cleaning. The analysed sample area has a diameter of about 1 mm2 (aperture #2) and the electron take-off angle is 458. Photoelectrons are excited by a non-monochromatic Mg Ka X-ray source operating at 350 W/15 kV. The PHI model 10-360 spherical sector electron spectrometer was operated with fixed pass energy of 89.45 eV for survey spectra and of 35.75 eV for high resolution spectra. The energy scale has been calibrated to the position of the Cu 2p3/2 and Au 4f7/2 peaks on sputter-cleaned samples. Published Cu 2p3/2 reference peak positions [5] are used for chemical state analysis (see Table 1). SnO and SnO2 cannot Table 1 Cu 2p3/2 binding energy and Cu-LMM kinetic energy for different Cu chemical states Chemical state
Binding energy (eV) Cu 2p3/2
Kinetic energy (eV) Cu-LMM
Cu metal CuO Cu2O Cu(OH)2 Cu3Sn [17]
932.6 933.7 932.5 935.1 933.2
918.6 918.1 916.2 916.2 918.1
Auger parameter (eV) 14.0 15.6 16.3 18.9 15.1
All peak energies are from Ref. [5], except the values for Cu3Sn. The Auger parameter is defined as the difference between the Auger electron kinetic energy and the photoelectron binding energy [16].
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easily be distinguished by XPS and AES and no attempt has been made to determine the Sn oxidation state. 2.4. Oxide bulk analysis by energy dispersive X-ray spectroscopy in a scanning transmission electron microscope Intermetallic compound layer and oxide bulk cross sections were prepared by focused ion beam. The electron transparent FIB lamellas have a thickness of about 100 nm. For more details of the FIB lamella preparation see Ref. [4]. An accurate quantification of the intermetallic composition and a semi-quantitative determination of the bulk oxide composition have been obtained by EDS in a STEM (Philips CM300UT).
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The RC value of 11 mV for the B sample is exceptionally low for cables produced by the same manufacturer and the chemical oxide composition of this sample differs strongly from that of other B samples in that it contains a large amount of Cl throughout the oxide layer. This Cl may be left from the flux that is used to activate the strand surface before the Sn–Ag deposition and may cause a reduction in RC. The RC value of 437 mV obtained for sample C is exceptionally high (typically RC values do not exceed 200 mV). The reason for this exceptionally high RC value is not understood. 3.2. Oxidation state of Cu oxides
3.1. Average Cu content in the oxide and interstrand contact resistance
Since the AES measurements at the energy resolution used in the present study do not provide information about the chemical state, the outermost oxide has been characterised by XPS. An indication about the bulk composition of the oxides present on different samples has been obtained using electron transparent oxide cross sections that were characterised by EDS in a STEM.
In Fig. 2 the average CuLMM/SnMNN values for a certain cable sample are plotted as a function of the corresponding RC values that were obtained for the same cable. It can be seen that there is a clear trend of the CuLMM/SnMNN ratio to increase with RC, with the exception of two out of the total of 16 cable samples, which are indicated by empty symbols. Excluding these two samples, which are considered to be outliers, a tentative linear correlation between the average Cu-LMM to Sn-MNN peak area ratio and the cable RC is fitted. The linear correlation coefficient is 0.77 when the outliers are not considered (when taking into account the two samples considered as outliers the correlation coefficient would be 0.22). The two samples considered as outliers are not taken into account for the following reasons.
3.2.1. Oxide thickness and bulk chemical composition as determined by EDS A STEM image of the oxide layer cross section on top of a continuous Cu3Sn coating after a 200 8C heat treatment in air is shown in Fig. 3. The oxide with a thickness of typically 20 nm is embedded between the Cu3Sn layer and a protective Pt layer, which is deposited onto the sample in order to protect the oxide during the preparation of the FIB lamella. The average composition of the intermetallic layer as determined by EDS is 75.1 at.% Cu and 24.9 at.% Sn. Because of the small oxide thickness an influence of the Cu3Sn layer on the oxide result cannot be excluded entirely and the obtained values are therefore considered as semi-quantitative results. For 10 EDS measurements, which gave a Sn content
Fig. 2. Average Cu-LMM to Sn-MNN peak area ratio as a function of interstrand contact resistance (RC) in LHC cables. In total the plot represents more than 250 Auger electron spectra that have been acquired on different interstrand contact zones. Cable samples of different manufacturers are indicated by different symbols. The two extreme points indicated by empty symbols are considered as outliers and not taken into account for the calculation of the correlation coefficient.
Fig. 3. STEM dark field image of a cross section through the Cu3Sn intermetalic layer and the strand oxide layer, which are formed during the 200 8C cable heat treatment in air. In the image the oxide layer with a thickness of about 20 nm is embedded between the Cu3Sn layer and a Pt layer that has been deposited onto the oxide prior the FIB lamella preparation.
3. Results
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Fig. 4. STEM dark field image of a fractured intermetallic layer and Sn La1, Cu Ka1 and O Ka1 chemical maps of the corresponding oxide cross section. It can be seen that Cu is absent in the oxide layers.
below 1 at.%, the average oxide composition is 30.3 at.% O and 69.5 at.% Cu. The Cu/O ratio of 2.3 indicates that the oxide bulk on top of the Cu3Sn layer is essentially composed of Cu2O rather than CuO. The fact that the ratio is higher than 2 is accounted to the proximity of Cu from the oxide substrate. In contrast to the oxide that it is observed on the continuous Cu3Sn layer, the oxide that is formed on a fractured intermetallic layer contains mainly Sn while almost no Cu is detected (see chemical maps in Fig. 4).
analysis shows that after room temperature oxidation the oxide consists of about 9% Sn oxide and 91% Cu oxide (weighted Cu2p/ Sn3d photoelectron ratio is 11.4). The Sn oxide content may be slightly underestimated because some of the detected Cu is metallic, which may indicate a contribution of the Cu substrate to the Cu signal. All Sn that is detected by XPS is oxidised. After the 200 8C air heating only Cu oxide can be detected by XPS, i.e. after the 200 8C heat treatment the Sn content in the oxide is below the XPS detection limit of about 1 at.%.
3.2.2. Outermost surface oxide composition as determined by XPS A typical photoelectron survey spectrum acquired on the oxide that is formed during the 200 8C heat treatment in air on top of a continuous Cu3Sn layer is shown in Fig. 5. The surface is mainly composed of Cu, O, some Sn and C from hydrocarbon surface contamination (the samples are not in situ sputtercleaned in order to avoid oxide reduction). The Cu chemical state at the outermost oxide surface is found from the high resolution photoelectron spectrum. From the Cu 2p3/2 energy position at 933.8 eV and the characteristic shape of the satellite peak at around 943 eV it can be concluded that the outermost Cu oxide is mainly composed of CuII as in CuO (see Table 1).
4. Discussion
3.2.3. Oxidation of Cu–8 at.% Sn bronze in air In order to understand what are the rate controlling steps in the oxidation of Cu–Sn alloys, a sputter-cleaned Cu–8 at.% Sn bronze sample was oxidised at 200 8C in air. XPS surface
4.1. The formation of Cu oxides on the LHC strands The AES analysis of more than 250 interstrand contact zones in LHC cables shows that Cu oxides are commonly present after the 200 8C heat treatment in air. In many cases only Cu oxide is detected, i.e. the Sn content in the oxide is below 1 at.%, which is the approximate detection limit of AES. In this case EDS analysis indicates that the oxide bulk is mainly composed of Cu2O and XPS reveals that the outer oxide surface is essentially composed of CuO. 4.1.1. Thermodynamics of oxide growth In Ref. [6] the formation of Cu oxides during heat treatments of Sn coated Cu has been ascribed to the formation of an intermetallic Cu3Sn layer, while on bronze or Cu6Sn5 mainly Sn oxide has been found. This observation is in agreement with the EDS/STEM results presented in this article, which detected only Cu oxide on a continuous Cu3Sn layer, while mainly Sn oxide is found when the coating is not fully reacted or when the intermetallic layer is absent. As shown in Table 2 the free energy of formation (DG0) for the stable Sn oxides (SnO and SnO2) is significantly higher than DG0 for the Cu oxides CuO and Cu2O [7]. Hence, thermodynamically the growth of Sn oxide should be favoured over the formation of Cu oxides, also on Cu3Sn. Indeed an XPS study [17] shows that the oxide formed on Cu3Sn bulk at 200 8C Table 2 Free energy of formation per mole of oxygen consumed for more stable tin and copper oxides [7] Reaction
Fig. 5. XPS survey spectrum and high resolution spectrum of the Cu 2p peak acquired on an interstrand contact after 200 8C cable heat treatment in air. The Cu 2p3/2 energy position at 933.8 and the shape of the satellite peak at around 943 eV show that the outermost Cu oxide is mainly in the form of CuO.
4Cu + O2 = 2Cu2O 2Cu + O2 = 2CuO 2Sn + O2 = 2SnO Sn + O2 = SnO2
DG0 (kJ/mol) 292.9 254.6 515.0 515.8
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in a pure O2 atmosphere at low pressure (4.3 mbar) is essentially composed of SnO and SnO2, as predicted by the calculated Cu–Sn–O phase diagram at 200 8C [17]. It can thus be concluded that under equilibrium conditions Cu2O and CuO are not formed on Cu3Sn and their presence must be due to kinetic effects.
the oxide–substrate interface and in the oxide layer, it is concluded that the presence of Ag in the strand coating does not have a strong influence on the strand oxidation behaviour at 200 8C in air.
4.1.2. Kinetics of oxide growth Both, copper(I) oxide Cu2O and copper(II) oxide CuO are semiconductors [8]. Cu2O is a cation deficient p-semiconductor, i.e. metal cations can diffuse along metal ion vacancies in the oxide towards the oxide/gas phase interface, where the oxidation process can proceed. In contrast to Cu2O, SnO2 is an anion deficient n-semiconducting oxide. In this case the oxide is oxygen deficient, which allows oxygen anions to diffuse through the oxide towards the metal/oxide interface, at which the oxide growth will proceed [9]. The reason for the different oxides that are formed on Cu3Sn at low O2 pressure and at atmospheric pressure in air may therefore be attributed to the different diffusion rates of Cu cations in Cu2O and oxygen anions in SnO2. At high O2 pressure (in air) the oxide growth is determined by the higher Cu cation diffusion rate with respect to the oxygen anion diffusion. This is confirmed by the fact that the Sn oxide vanishes upon 200 8C oxidation of Cu–8 at.% Sn bronze in air. At low O2 pressure the oxide growth is determined by the thermodynamics. In this case the oxygen absorption and diffusion towards the interface is oxidation rate controlling.
The formation of a continuous Cu3Sn layer during the cable heat treatment is crucial to obtain the desired RC values [4]. EDS and XPS analysis reveal that the oxide formed during the 200 8C heat treatment at atmospheric pressure in air on Cu3Sn consists essentially of Cu2O and a thin outer layer of CuO. This oxide structure is similar to that formed on bare Cu that is oxidised at elevated temperatures in air [13]. Similarly to oxidised Cu3Sn, RC of oxidised Cu is comparatively high, but in this case of bare Cu strands RC drops dramatically during heat treatments that are performed under pressure, simulating the curing process in LHC magnets [14]. Therefore, the Cu3Sn layer may act as a diffusion barrier for oxygen ions, stabilising the Cu oxide during the curing treatment [15]. The data presented in Fig. 2 shows a clear relation between interstrand contact resistance and the oxide type, i.e. the more Cu oxide is present on the strand surfaces the higher is the interstrand contact resistance. The reason for this might be a higher electrical resistance of the p-semiconducting Cu oxide with respect to the n-semiconducting Sn oxide, and/or differences in oxide thickness. Another explanation may be that the formation of Cu oxide during the cable heat treatment indicates that a continuous Cu3Sn layer has been formed, which increases the oxide stability.
4.1.3. Reasons for the oxide inhomogeneity The main reasons why different oxides are formed on the strands during identical heat treatments are irregularities of the underlying strand coatings. These irregularities can for instance be the result of insufficient wetting of the Cu substrate by the Sn–Ag alloy, remaining flux that is applied prior to the hot dip coating process or because of differences in the coating temperature and/or cooling rates, which influence the coating microstructure [10]. Other elements in the coating bath, apart from Cu and Ag, can also alter the coating chemistry and microstructure [11], and, as a result, influence the surface oxide structure and thickness that grows during the 200 8C air heating treatment. 4.1.4. The influence of Ag in the coating on the surface oxidation The Ag in the Sn–3.5 wt.% Ag solder is present in the form of intermetallic Ag3Sn particles, which increase the mechanical strength of the non-reacted coating. EDS analysis of reacted coating cross sections shows that also in this case Ag is in the form of Ag3Sn. At lower temperatures (up to 170 8C) the Ag content can influence the Cu3Sn growth rate while at 205 8C no significant difference in the Cu3Sn growth rate with respect to pure Sn–Cu diffusion couples has been observed [12]. In the present study Ag was never detected at the coating surface or in the oxide layer. Since the presence of Ag does not significantly influence the intermetallics growth rate and because no Ag can be detected at
4.2. Strand oxide composition and RC
5. Conclusion The interstrand contact resistance in the LHC superconducting cables is related to the oxide composition, i.e. the more Cu oxide relative to the amount of Sn oxide the higher is the contact resistance. Despite the fact that for stable Sn oxides the free energy of formation is higher than for the stable Cu oxides, on many strand samples only Cu oxide is detected. Cu oxide formation depends critically on the formation of a continuous Cu3Sn layer. The Cu oxide formed during the 200 8C heat treatment in air has the same structure as the oxide formed on bare Cu (thin CuO layer on top of Cu2O bulk). The fact that RC on the Sn coated strand Cu matrix remains sufficiently high during curing, while it is drastically reduced on the un-coated Cu indicates that the Cu3Sn layer is stabilising the Cu oxide during the heat treatment under high pressure. Acknowledgements We are grateful to D. Richter and A. Bastos Marzal for RC results and advice about interstrand contact resistance measurements.
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