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Journal of Alloys and Compounds 497 (2010) L13–L16

Contents lists available at ScienceDirect

Journal of Alloys and Compounds

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etter

echanical stability of Si thin film deposited on a Ti–50.3Ni(at%) alloy

o-min Kima, Gyu-bong Choa, Jung-pil Noha, Hyo-jun Ahna, Eun-soo Choic,huichi Miyazakia,b, Tae-hyun Nama,∗

School of Materials Science and Engineering & ERI, Gyeongsang National University, 900 Gazwadong, Jinju, Gyeongnam 660-701, Republic of KoreaInstitute of Materials Science, University of Tsukuba, Tsukuba, Ibaraki 305-8573, JapanDepartment of Civil Engineering, Hongik University, Seoul 121-791, Republic of Korea

r t i c l e i n f o

rticle history:eceived 18 January 2010eceived in revised form 24 February 2010ccepted 1 March 2010vailable online 7 March 2010

a b s t r a c t

Silicon thin film annealed at 973 K for 7.2 ks after being deposited on a Ti–50.3Ni(at%) substrate was notdetached from the substrate after 2.2% tensile deformation, which was ascribed to a diffusion bondingbetween the Si film and substrate. The B2–B19′ transformation start temperature (Ms) of the Ti–Ni sub-strate with Si thin film increased by annealing, which was ascribed to a tensile stress developed by thedifference in thermal expansion coefficient between the Si film and substrate.

eywords:hape memory alloys (SMA)oatingilicon thin filmension testartensitic phase transformation

© 2010 Elsevier B.V. All rights reserved.

node

. Introduction

There have been some attempts to apply Ti–Ni alloys for Li/metalulfide and Li ion secondary batteries with high flexibility requiredor mobile electrical appliances [1,2]. In commercial Li ion battery,raphitic carbon has been used as an anode material because ofigh discharge voltage, long cycle life and low cost. The availableapacity of graphitic carbon is approaching near the theoreticalalue (∼370 mAh g−1). In order to improve capacity of Li ion sec-ndary battery, silicon (Si) instead of graphitic carbon has beenxamined as an anode material because of its high theoreticalapacity of 4200 mAh g−1 on the basis of Li22Si5 formation [3,4],hich is 11 times higher than the capacity of the graphitic car-

on.Unfortunately, Si electrode has a serious problem which orig-

nates from the large volume change of Si (∼310%) duringharge–discharge (lithiation–delithiation) process. The large vol-me change causes the surface cracking and pulverization of Si

lm by the repetitive mechanical stress associated with the vol-me change and leads to poor cycle life [5,6]. Many researchesave been made to solve the problem of poor cycle life of Si elec-rode by enhancing the adhesion between Si and a current collector

∗ Corresponding author. Tel.: +82 55 751 5307; fax: +82 55 751 1749.E-mail address: tahynam@gsnu.ac.kr (T.-h. Nam).

925-8388/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.jallcom.2010.03.023

(substrate) since the generation of the stress is restrained by theenhanced adhesion [7–11].

Ti–Ni alloys are good candidates for substrate materials toaccommodate the stress generated during charge–discharge ofSi deposited on Ti–Ni substrates. The stress generated duringcharge–discharge of Si electrode is known to be about 3.0 GPa[12] which is very large comparing to the stress (0.1–0.6 GPa) fordeforming Ti–Ni substrates through rearrangement of the B19′

martensitic variants or the B2–B19′ stress-induced martensitictransformation in Ti–Ni alloys [13]. Therefore, Ti–Ni substratesare expected to accommodate some of the stress generatedduring charge–discharge of Si deposited on them by rearrange-ment of the B19′ martensitic variants or the stress-inducedmartensitic transformation and thus improve cycle life of Si elec-trode.

In order to improve cycle life of Si electrode, Si should not bedetached from a current collector (substrate) during the defor-mation of Ti–Ni substrates accompanied by charge–discharge.Therefore, it is essential to investigate a mechanical stability ofSi electrode during the deformation of Ti–Ni substrates associatedwith rearrangement of the B19′ martensitic variants or the stress-

induced martensitic transformation.

In this study, Si thin film was deposited on the surface ofTi–Ni substrates and then the mechanical stability of the Si thinfilm during deformation of the substrate was investigated. Effectsof Si thin film on martensitic transformation temperatures and

L14 B.-m. Kim et al. / Journal of Alloys and Compounds 497 (2010) L13–L16

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ig. 1. (a) and (b) are FE-SEM photographs of the Si thin film deposited on the surfaceespectively. (c) and (d) are XRD patterns obtained from as-deposited and annealed

hape memory characteristics of Ti–Ni substrates were investigatedlso.

. Experimental procedure

A Ti–50.3Ni(at%) alloy ingot was prepared by vacuum induction melting. Thelloy ingot was hot rolled into a sheet with a thickness of 0.5 mm. Ti–Ni substrates fori deposition were cut from the sheet and then solution treated at 1123 K for 3.6 ks inacuum. Dimensions of the substrates for samples of differential scanning calorime-ry (DSC), X-ray diffraction (XRD), field emission scanning electron microscopyFE-SEM) and thermal cycling tests under constant load were 3 mm (ϕ) × 0.2 mmt), 15 mm (L) × 15 mm (W) × 0.5 mm (t), 10 mm (L) × 10 mm (W) × 0.5 mm (t) and0 mm (L) × 3 mm (W) × 0.5 mm (t), respectively. Prior to Si film deposition, oneide of Ti–Ni substrates was electrochemically etched in 1 M H2SO4 + 2% HF solu-ion to enhance the adhesion with Si film during deposition process. Si thin filmere prepared by DC magnetron sputtering under the pressure of ∼5 × 10−3 Torr of

rgon (99.995%). Thickness of Si film was 560 nm, measured by an alpha step profilersing Alphastep, KLA Tencor. Si films deposited on Ti–Ni substrates were annealed

t 973 K for 7.2 ks in a vacuum to induce the inter-diffusion between Si film andi–Ni substrates.

The crystal structure of Si films and Ti–Ni substrates was investigated by XRDith Cu K� radiation at room temperature using Miniflex, Rigaku. The surface fea-

ures of the etched Ti–Ni substrate and the Si films were characterized by FE-SEMsing XL30S, Philips. Martensitic transformation temperatures were investigated

Fig. 2. DSC curves obtained from the etched Ti–50.3Ni(at%) alloy su

chemically etched Ti–50.3Ni(at%) alloy substrate showing surface and cross-section,les, respectively.

by DSC with a cooling and heating rate of 0.17 K/s. Thermal cycling tests under theapplied stress of 40 MPa with a cooling and heating rate of 0.017 K/s were made forinvestigating the shape memory effect. Si films were investigated by FE-SEM afterthe thermal cycling tests.

CR2032 coin-type cells were assembled in an argon-filled glove box with alithium metal anode (Cyprus Foote Mineral, 99.98%, USA) as the count electrode.The electrolyte was 1 M LiPF6 in a 1:1 mixture of EC and DMC. Galvanostaticcharge–discharge tests were performed in a range from 0.01 V to 1.5 V vs. Li/Li+,at a current density of 200 �A cm−2 corresponding to 1.5 Ah/g (0.42C)-as-depositedSi.

3. Results and discussion

Fig. 1(a) is a FE-SEM image showing the surface of Si thin filmdeposited on the surface of the Ti–Ni substrate prepared by electro-chemical etching. Si thin film is transparent and thus small poreswith a mean diameter of ∼3 �m fabricated by the electrochemi-

cal etching on the surface of the substrate are observed through Sifilms. Fig. 1(b) is a cross-sectional FE-SEM image of the Si depositedTi–Ni substrate in which Si film with a thickness of 560 nm is formedon the Ti–Ni substrate. It is evident that any macroscopic defects arenot observed at the interface between the Si film and Ti–Ni surface.

bstrate (a), as-deposited sample (b), and annealed sample (c).

B.-m. Kim et al. / Journal of Alloys and Compounds 497 (2010) L13–L16 L15

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Fig. 3. Temperature vs. elongation curves obtained

XRD pattern obtained from the Si deposited Ti–Ni substrateas-deposited sample) is shown in Fig. 1(c). The diffraction peakorresponding to the B2 parent phase is observed. Any diffractioneaks of Si are not found, which means that the Si film deposited onhe Ti–Ni substrate is amorphous. XRD pattern obtained from thei deposited Ti–Ni substrate annealed at 973 K for 7.2 ks (annealedample) is shown in Fig. 2(d). Diffraction peaks corresponding tohe B2 parent phase, Ti5Si4 and Ni2Si are observed simultaneously.his means that annealing of the as-deposited Si film at 973 Korms Ti5Si4 and Ni2Si by an inter-diffusion between the Si film andi–Ni substrate, which plays an important role for improving bondtrength between the Si film and Ti–Ni substrate as will be men-ioned later. It is also found that any diffraction peaks of Si do notppear in Fig. 1(d), suggesting that amorphous Si is not crystallizedfter annealing at 973 K.

In order to investigate an effect of Si thin film on martensiticransformation of Ti–Ni substrate, DSC measurements were madend then results obtained are shown in Fig. 2. Fig. 2(a) showsSC curves of a Ti–50.3Ni(at%) substrate whose surface was elec-

rochemically etched after solution treated at 1123 K. One clearSC peak is observed on each cooling and heating curve, whichriginates from the B2–B19′ martensitic transformation, which isimilar to the result obtained from a solution treated Ti–50.3Ni(at%)lloy [14]. This means that the electrochemical etching does notave a substantial effect on martensitic transformation of a solu-ion treated Ti–50.3Ni(at%) alloy, which is ascribed to the fact thathe thickness of the etched region (<1 �m) is very thin compar-

ng to the entire thickness of the sample (∼0.5 mm). Fig. 2(b) andc) show DSC curves obtained from the as-deposited and annealedamples, respectively. Comparing Fig. 2(a) and (b), it is found thathe B2–B19′ transformation start temperature (Ms) of the sam-les before and after Si deposition is very similar, which means

Fig. 4. FE-SEM photographs taken from the surface of as-deposited sam

as-deposited sample (a) and annealed sample (b).

that the Si film deposited on the porous Ti–Ni substrate does nothave a substantial effect on Ms. On the other hand, Ms of the sam-ple annealed at 973 K after Si deposition is 6 K higher than that ofthe Ti–Ni substrate. This is considered to be related with an inter-diffusion between the Si film and Ti–50.3Ni substrate, leading to anenhancement of bonding strength between them.

Tensile stress applied on Ti–Ni alloys is known to raise Ms

according to the Clausius–Clapeyron equation [15]. Thermal expan-sion coefficient of Si (∼4 × 10−6/K) is smaller than that of Ti–Nisubstrate (∼1 × 10−5/K). When the Si film is tightly boned with thesubstrate due to the inter-diffusion, a tensile stress may be devel-oped on the substrate. Therefore the increase in Ms in Fig. 2(c) isascribed to the tensile stress which originates from the differencein thermal expansion coefficient. A contribution of compositionalchange associated with the formations of Ti5Si4 and Ni2Si to thechange in Ms is also taken into consideration. Fig. 2, it is alsofound that the transformation heat (�H) associated with theB2–B19′ transformation of the substrate is found to be 30.3 J/gand it decreases slightly to 29.4 J/g by depositing Si film on thesubstrate. Significant decrease in �H is observed by annealing at973 K (27.2 J/g), which is ascribed to the formation of Ti5Si4 andNi2Si which do not contribute to �H associated with the B2–B19′

martensitic transformation. Large deviation from an equiatomiccomposition (Ti–50.0Ni) in the matrix near Ti5Si4 and Ni2Si alsodecreases �H since the B2–B19′ martensitic transformation wouldnot occur in the area.

In order to investigate an effect of Si thin film on the shape mem-

ory effect of the Ti–Ni substrate, thermal cycling tests under theapplied stress of 40 MPa were made and then temperature (T) vs.elongation (ε) curves obtained are shown in Fig. 3. Fig. 3(a) showsthe T–ε curve of the as-deposited sample and transformation elon-gation (εR) associated with the B2–B19′ transformation is found to

ple (a) and annealed sample (b) after 2.2% tensile deformation.

L16 B.-m. Kim et al. / Journal of Alloys and C

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ig. 5. Charge–discharge cycle performances of cells composed of as-depositedample and annealed sample. Capacities were normalized with respect to the max-mum capacity of each cell.

e 2.7%. Fig. 3(b) is the T–ε curve of the annealed sample and εRs found to be 2.2%. This means that Si film annealed at 973 K aftereing deposited has a substantial effect on εR, which is ascribed tohe tightly bonded Si film that restrict transformation elongationf the Ti–50.3Ni substrate as will be mentioned later.

After thermal cycling tests under the applied stress of 40 MPa,he surface of the samples were examined by FE-SEM observationsnd then results obtained are shown in Fig. 4. Fig. 4(a) is FE-SEMhotograph obtained from the as-deposited sample. Many largeracks are observed as shown by single headed arrows. It is alsoound that some parts of the Si film are detached from the Ti–Niubstrate as shown by double headed arrows, which means thathe bonding between the Si film and the substrate is poor. In con-rast to Fig. 4(a), only small amount of cracks are found in thennealed sample as shown in Fig. 4(b). It is to be noted here thathe Si film is not detached from the substrate after 2.2% tensileeformation.

The Ti–Ni substrate is elongated by the B2–B19′ martensiticransformation on cooling under 40 MPa. The Si film deposited onhe surface of the Ti–Ni substrate would be cracked during the cool-ng procedure because the Si film is so brittle that it is not elongatedy 2.2%. On heating, the elongation of the Ti–Ni substrate is recov-red by the B19′–B2 reverse transformation of the Ti–Ni substrate.uring the heating procedure, behavior of the Si film cracked onooling is considered to depend on the bond strength between thei film and Ti–Ni substrate. When the bond strength is high, the Silm would keep strong contact with the Ti–Ni substrate, thus theracks would be closed with the shape recovery of the Ti–Ni sub-trate. When the bond strength is low, however, the Si film wouldake loose contact with the Ti–Ni substrate, thus the cracks are not

losed completely.

From Fig. 4, therefore, it is concluded that the Si thin film

nnealed at 973 K for 7.2 ks after being deposited on the Ti–Ni sub-trate is mechanically stable even after 2.2% tensile deformation.he difference between Fig. 4(a) and (b) is ascribed to the fact thatdiffusion bonding occurs with the formation of Ti5Si4 and Ni2Si in

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ompounds 497 (2010) L13–L16

the annealed sample, which would increase bond strength betweenthe Si film and the Ti–Ni substrate.

Electrochemical cycle life is actually compared with cells com-posed of as-deposited Si film electrode and annealed electrode.Fig. 5 shows charge–discharge cycle performances of two cells. Theas-deposited electrode exhibits rapid capacity fading at the initialcycling (∼13 cycles) and its capacity retention is 30% after 50 cycles,whereas the annealed Si electrode shows remarkably improvedcycle performance and its capacity retention is 77%. The improvedelectrochemical cycle life of annealed electrode is ascribed to theenhanced adhesion between the Si film electrode and the TiNi sub-strate (current collector).

4. Conclusions

In summary, only a few small cracks were observed on theSi thin film annealed at 973 K for 7.2 ks after being deposited onthe Ti–50.3Ni(at%) substrate after 2.2% tensile deformation, whilemany large cracks were observed on the as-deposited Si thin film.Some of the Si film was detached from the Ti–Ni substrate in the as-deposited sample, while the Si film is not detached in the annealedsample. The enhanced adhesion between the Si film and the Ti–Nisubstrate, which were ascribed to an inter-diffusion bonding withthe formation of Ti5Si4 and Ni2Si, lead the longer cycle life of Sifilm electrode than the as-deposited Si electrode. Ms of the Ti–Nisubstrate was raised by annealing the Si film deposited on the sub-strate, which was ascribed to a tensile stress developed by thedifference of thermal expansion coefficient between the Si film andthe Ti–Ni substrate.

Acknowledgements

This research was supported by WCU (World Class Univer-sity) program through the National Research Foundation of Koreafunded by the Ministry of Education, Science and Technology (Grantnumber: R32-2008-000-20093-0) and also by Pioneer ResearchCenter for Nano-morphic Biological Energy Conversion and Storage.

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