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The influence of surface oxides on the distribution and release of nickel

from Nitinol wires

Svetlana A. Shabalovskaya a,c,*, He Tian b, James W. Anderegg c, Dominique U. Schryvers b,William U. Carroll d, Jan Van Humbeeck a

aMetallurgy and Material Science Department, Katholieke University Leuven, 3001 Leuven, Belgiumb EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, BelgiumcAmes Laboratory – DOE, Ames, IA 50011, USAdChemistry Department, National University of Ireland, Galway, Ireland

a r t i c l e i n f o

Article history:

Received 18 August 2008

Accepted 16 October 2008

Available online 8 November 2008

Keywords:

Nitinol

Ti oxides

Ni ion release

Corrosion

Biocompatibility

Intimal hyperplasia

a b s t r a c t

The patterns of Ni release from Nitinol vary depending on the type of material (Ni–Ti alloys with low or

no processing versus commercial wires or sheets). A thick TiO2 layer generated on the wire surface

during processing is often considered as a reliable barrier against Ni release. The present study of Nitinol

wires with surface oxides resulting from production was conducted to identify the sources of Ni release

and its distribution in the surface sublayers. The chemistry and topography of the surfaces of Nitinol

wires drawn using different techniques were studied with XPS and SEM. The distribution of Ni into

surface depth and the surface oxide thickness were evaluated using Auger spectroscopy, TEM with FIB

and ELNES. Ni release was estimated using either ICPA or AAS. Potentiodynamic potential polarization of

selected wires was performed in as-received state with no strain and in treated strained samples. Wire

samples in the as-received state showed low breakdown potentials (200 mV); the improved corrosion

resistance of these wires after treatment was not affected by strain. It is shown how processing tech-

niques affect surface topography, chemistry and also Ni release. Nitinol wires with the thickest surface

oxide TiO2 (up to 720 nM) showed the highest Ni release, attributed to the presence of particles of

essentially pure Ni whose number and size increased while approaching the interface between the

surface and the bulk. The biological implications of high and lasting Ni release are also discussed.

Ó 2008 Elsevier Ltd. All rights reserved.

1. Introduction

The release of Ni from Nitinol (a group of nearly equiatomic

Ni–Ti alloys with shape memory effect and superelasticity) is

a central issue in its biocompatibility. However, our under-

standing of the patterns of Ni release from the material evolved

slowly. Observations on Nitinol prepared under laboratory

conditions [1–3] showed that Ni release might be higher than

that from stainless steel during the first days of exposure to

biological solutions but it drops to almost undetectable levels

after 10–14 days. Studies on commercial Nitinol, however, point

at different patterns: on initial exposure Ni release increases and

does not drop even after a few months [4–7]. The amount of

released Ni also differs significantly depending on the variable Ni

surface concentrations reported for Nitinol wires (0.4–15 at.%) [8].

High-temperature annealing in air employed in wire processing

promotes the formation of a thick external TiO2 layer that inev-

itably should lead to Ni accumulation in the surface depth; this

buried Ni can be easily released through the defective surfaces

resulting from wire processing.

Another avenue leading to the accumulation of Ni in the surface

layers can be the type of surface treatment itself. Some of the tech-

niques used for deposition of coatings onto a Nitinol surface employ

low temperature (60–160 �C) pre-treatment protocols aiming at the

formation a thick TiO2 layer, similar to that grown on pure Ti. As was

demonstrated inRef. [9], the resulting surfaceshavealsoburiedNi-rich

sublayers. These result in high and lasting Ni release, exceeding the Ni

release fromnon-treatedmaterialby twotothreeordersofmagnitude.

The goal of the present study was to identify the sources of Ni

release from Nitinol wires with the original black or dark blue

lustrous oxides formed during wire processing involving numerous

drawing cycles followed by annealing in air at a temperature

slightly higher than 700 �C. As it follows from the analysis of the

literature on medical applications of Nitinol wires, these oxides are

gaining in populariry as final surfaces in medical devices such as

defect closures for heart and hernia, for instance. The surface

chemical composition of the wires and the elemental states of the

* Corresponding author. Svetlana Shabalovskaya, Ames Laboratory, Ames, IA

50011, USA.

E-mail addresses: [email protected], [email protected] (S.A. Shabalovskaya).

Contents lists available at ScienceDirect

Biomaterials

journal homepage: www.elsevier .com/locate/biomater ia ls

0142-9612/$ – see front matter Ó 2008 Elsevier Ltd. All rights reserved.

doi:10.1016/j.biomaterials.2008.10.014

Biomaterials 30 (2009) 468–477


surface constituents were evaluated with the help of XPS. Auger

spectroscopy was used for depth profiling and evaluation of

thickness of the surface oxides. Small collection spots,1 mM and less,

attainable using Auger spectrometer compared to approximately

3 mm2 collection areas in XPS allowed for working with the

required precision with thin wire samples. TEM was employed to

visualize the Ni distribution clearly in the surface depth down to

the interface with the bulk. In order to discriminate between TiO2

and TiO, Electron Energy Loss Near Edge Spectra (ELNES) were

recorded. Finally, the corrosion resistance of the original and

treated wires was evaluated.

2. Material and methods

2.1. Material and samples

Three types of Nitinol wires of 0.75 mm diameter with original oxides were

studied. According to Ref. [10] wires 1 and 2 were drawn using synthetic poly-

crystalline diamond dies, and wire 3 was drawn using single crystal natural dia-

mond. The wires were straightened at 500 �C under various levels of argon/oxygen

atmosphere. Before the surface studies the wire samples were cleaned in ultrasonic

baths of alcohol and then deionized water. Control wire samples for corrosion tests

under stress were chemically etched in a 1HFþ 4HNO3þ 5H2O solution.

2.2. DSC analysis

DSC analysis carried out using Universal V4.4A TA Instruments in the temper-

ature range fromÿ150 to 200 �C showed that at room temperature all wires were in

the high-temperature austenitic phase and the martensitic transformation on

cooling started at 10 �C.

2.3. Mechanical test

Mechanical tensile tests (Instron 4467) performed to assure superelastic wire

performance were conducted with a rate of 5 mm/min. A martensitic plateau up to

5.5% strain and a yield strength of 629 MPa were observed.

2.4. XPS analysis

XPS surface analysis was performed with a PHI 5500 spectrometer using Al Ka1radiation with a 45� electron collection angle. Survey spectra were acquired in the

range from 0 to 1100 eV at pass energy of 187 eV. The high resolution spectra of the

surface constituents were acquired from three spots for each sample with 58.7 eV

pass energy. In the XPS study, the samples were mounted over a 1�2� 0.5 cm

channel in a samplemount. The use of amonochromatic X-ray source allowed for the

excitation of electrons only from those surfaces that were detected by the analyzer.

Since thewire samples had identical diameters andwere placed in a parallelmanner,

andexamined simultaneously,wemaycompare the results forall threewires studied.

2.5. Auger analysis

An Auger microprobe (Jeol Jamp 7830F) with a base pressure of less than

1�10ÿ9 Pa was employed for depth profiling using Arþ ion bombardment at 3 kV

energy. The thickness of the surface oxides was estimated using the point at which

the oxygen peak reached half of its maximal concentration. In the absence of

standards (reference Ti oxide samples), the sputtering rate of Nitinol surface oxides

was approximated by that of SiO2 (12 nM/min) obtained under similar conditions, as

it is routinely accepted. The use of SiO2 standards allows for the comparison

between various XPS or Auger studies. All studied samples were mounted in parallel

and kept in the mount until the end of the study to avoid possible interference

associatedwith the geometry of the thinwires. The data from two collection areas of

different sizes (1 mM2 and 50 mM

2) were analyzed. Since the 1 mM2 areas could be

approximated as flat ones the use of a corresponding sputtering rate obtained on flat

SiO2 samples is justified for comparative analysis of the thickness of surface oxides.

The 50 mM2 areas were selected to have a broader few of the surface and to compare

the results of the present study with those reported earlier [7]. It is questionable

whether this larger area could be also approximated as a flat area; however, the

comparative analysis of the thickness of different wires is still valued.

2.6. Potentiodynamic cyclic polarization

The potentiodynamic (PD) curves were obtained using a PAR 273A potentiostat

after 1 h exposure of the wire samples to a 0.9 NaCl solution previously purged with

nitrogen also for 60 min. A common three electrodes configuration (working elec-

trode, counter graphite and reference Ag/AgCl electrode) was used. The PD tests

were performed at a sweeping rate of 600 mV/h beginning at ÿ0.6 V and ending on

the reverse scan at the potentials 20 mV higher than Ecorr. The details of acquiring

the PD polarization curves under stress have been described elsewhere [11].

2.7. Structure characterization

Conventional TEM studies were performed using a CM20 Philips microscope.

The HRTEM imaging and Electron Energy Loss Spectroscopy (EELS) studies were

conducted using an UltraTwin CM30 Philips FEG instrument equipped with a post-

column GIF200 detector. The TEM specimens were prepared by Focused Ion Beam

(FIB) cutting with Gaþ ions, following the so-called in-situ lift-out method by

removing material from both sides of the desired cross-section. To prevent surface

damage by the incident Gaþ ions, the area of interest is covered in-situ with

a sputtered platinum layer. Cut slices of 15� 5 mM size with 200 nM thickness were

lifted out by a needle and transported to amodified TEM grid. For regular cutting and

thinning the Gaþ ions were accelerated with 30 kV; at the final stage the voltage was

reduced to 5 kV. After continued thinning down to electron transparency, a final

cross-section thickness of 60 nM was achieved. Through adjusting the angle of the

incident ion beam, thinned flat regions of 3–5 mM were acquired.

2.8. Ni release

Ni release into 0.9 NaCl solution (45 mL) fromNitinol wires (two samples of each

wire type) with a surface area of 2.65 cm2 was evaluated using either ICP-OES

analysis or Flame Atomic Absorption measured with a Perkin Elmer A Analyst 3300

instrument. The limits for Ni detection for these two methods were 0.005 mg/L

(wires 2 and 3) and 0.02 mg/L (wire 1), respectively. Prior to immersion the cut ends

of the wires were polished to a 600 grit finish. Tubes were sealed with rubber corks

and stored at 37 �C. After periods of 0.5, 1, 2, 4, and 6 months, respectively, the

samples were removed and the solution was analyzed for Ni content.

3. Results

3.1. Ni release

All wires showed continued Ni release that increased with the

time of exposure and did not drop even after five months (Fig. 1).

The release of Ni from wire 1 exceeded that from wires 2 and 3 as

much as one hundred times.

3.2. Surface topography

Examination of the surface topography of the wires revealed

significant differences (Fig. 2). Wires 2 and 3 showed a uniform

tweed-like morphology with longitudinal strips resulting from

drawing as well as occasional scratches that could be distinguished

as crossing the wire at an angle different from the drawing direc-

tion. In contrast, wire 1 exhibited a rather irregular surface, where

the external surface layer (scale) was occasionally cracked or

missing completely. Cracking in the direction perpendicular to the

drawing direction indicates high surface stresses. The variation in

surface topography among the studied wires points at the different

processing procedures, related probably to the different dies used

for wire drawing. Although it was reported previously [10] that

wires 1 and 2 were drawn using similar dies, it is obvious from the

SEM evaluation and from the TEM study that, in fact, wires 2 and 3

were drawn in a similar manner. This conclusion is in agreement

with the behavior of Ni release that also revealed similarities

between wires 2 and 3 but not between wires 1 and 2.

3.3. Surface chemical analysis

The survey XPS spectra of all three types of studied wires

showed that their surfaces were essentially clean (Fig. 3). The total

level of contaminants like Ca, S, Cl, and Nwas not higher than 3 at.%.

However, in addition to these elements, molybdenum (Mo less than

1%) and thallium (Tl less than 0.5%) were also detected on the

surfaces of wires 2 and 3. Both elementsMo and Tl (a toxic element)

could be constituents of the lubricants used for the wire drawing.

The elemental concentrations and the Ti/Ni ratios are presented in

Table 1. Both wires 2 and 3 showed a Ti/Ni ratio between 11 and 12.

In contrast, for wire 1, where Ni was a prevailing metal on the

surface (7.8%), the Ti/Ni ratio dropped to 0.6.

S.A. Shabalovskaya et al. / Biomaterials 30 (2009) 468–477 469


The high resolution XPS spectra of the major surface constitu-

ents for wire 2 and the selected spectra for wire 1 are presented in

Fig. 4. On the surfaces of all wires, Ti was in the oxidizedþ4 state, as

the location of the 2p spectrum indicates (w458 eV, [12]). A

negligible contribution at the 456 eV indicates the presence of

a small amount of other oxides of Ti. The Ni oxidation state varied

from one wire to another, as indicated from the analysis of the Ni

2p3/2 spectra. While wire 1 had Ni only in the oxidized state

(857 eV), wire 2 revealed nickel in both the oxidized Niþ3 (857 eV)

and the elemental (852.8 eV) states. The 1s spectrum of oxygen

showed a double peak structure. The low energy part below 530 eV

represents oxygen in organic contaminants. There is a slight shift of

the high energy peak around 530 eV representing oxygen in the

metallic oxides. While for wires 2 and 3 this peak is located at

530.5 eV implying the oxygen bond to Ti, for wire 1, it is shifted to

529.5 eV pointing at a significant contribution fromNi oxide. The 1s

peaks for carbonwere registered at 285 eV (CH), 286.5 (C–O) and at

289 eV. The latter peak can be assigned to a carbon compound with

nitrogen, oxygen and hydrogen (O¼ C–N–H). The 1s spectrum of

nitrogen revealed two distinct peaks implying the presence of two

types of N on the surface, in an organic state (400.6 eV) and in

a metallic compound (407.5 eV), like Ni(NO3)$6H2O, for instance.

The Thallium 4f duplex could be easily distinguished on the

surfaces of wires 2 and 3.

3.4. Auger depth profiles

The Auger depth profiles were acquired for all three wires using

two sizes of collection spots. There were no significant differences

between the shapes of the depth profiles acquired with a 50 mM2

raster for wires 2 and 3; however, there were significant variations

in the distribution of Ni in the surface depth between the depth

profiles acquired with a 1 mM2 raster, even within the same wire

sample. These variations were more pronounced for wire 1 (Fig. 5).

In the case of wires 2 and 3 there was a gradual increase of Ni

content into surface depth from about zero percent. The situation

was different in the case of wire 1 where the Ni concentration was

Fig. 2. SEM images of the surfaces of the studied wires.

Fig. 1. Ni release into 0.9 NaCl solution from Nitinol wires 1, 2 and 3.

S.A. Shabalovskaya et al. / Biomaterials 30 (2009) 468–477470


high from the very beginning (w13%). Also, while the depth profiles

obtained from various locations for 1 mM2 spots from wires 2 and 3

could be averaged yielding the ones obtained using a 50 mM2 raster,

this was not possible in the case of wire 1. It can be seen from Fig. 5,

upper right panel, that a ‘local’ depth profile (1 mM2 spots) acquired

from a scale with non-damaged surface exhibits one Ni peak

directly on the surface (w12 at.%) and another – on the interface

with the bulk (w80 at.%), pointing at a significant accumulation of

Ni. The latter Ni depth profile resembles those typical for Nitinol

samples heat treated at T� 600 �C [9,13]. The second 1 mM2Ni depth

profile for wire 1 (lower left panel) acquired from an area with

a missing scale is drastically different. As can be seen, at this new

location the surface oxide is as much as ten times thinner. The

thickness of surface oxides for all three wires varied from site to site

over a wide range (Table 2). The thickest oxide layer was detected

on wire 1 with an intact scale and the thinnest layer was observed

on wire 3.

3.5. The TEM study

The results of the TEM study together with the elemental maps

for nickel, titanium and oxygen are shown in Figs. 6 and 7. As can be

seen for wire 2 (Fig. 6a), two areas with slightly different grayscale

can be distinguished in the external surface adjacent to the

protective Pt layer. The difference in the grayscale is due to

the different Ti oxides. This also follows from the oxygen map (d).

The total thickness of the oxide layer remaining after FIB thinning

for wire 2 varies from 50 to 100 nM as can be deduced from image

(a) and from the maps of Ti (c) and oxygen (d). A deeper layer that

looks rather irregular on the TEM image (a) but can be more clearly

recognized as the brighter contrast on the Ni map (b) and darker

contrast on the Ti map (c) at the interface with the bulk of the alloy,

was identified using energy dispersive X-ray analysis as well as

electron nanodiffraction as Ni3Ti, confirming recent results

received on Nitinol microwires [14]. As an example, Fig. 7 shows

a [212] nanodiffraction zone pattern together with the corre-

sponding simulated patterns obtained on wire 2 using the D024Ni3Ti structure as input. The faint 10–1 superlattice reflections are

a special characteristic of this structure. Also, the measured inter-

planar spacings and angles obtained from this and other zones

confirm the same conclusion. According to the Ni–Ti phase

diagram, this Ni3Ti could have been formed by decomposition in

the B2 phase during production of the wire [14]. The thickness of

this sublayer reaches 70 nM, comparable to the Ti oxide layer itself.

It is clear from the Ti map (c) that Ni3Ti can be found as close as

30 nM to the surface.

The situation is different in case of wire 1. The TEM images

presented in Fig. 8 do not reveal the full width of the oxide due to

the complete removal of the Pt protection layer and of the part of

the top surface of the oxide film by FIB thinning. Still, the total

thickness of the oxide layer could be measured as 160–190 nM from

other parts of the thinned sample, which contained the Pt protec-

tion layer but where insufficiently thinned to produce good EFTEM

images. Damaged surface areas with partially missing oxide,

possibly associated with deep scratches or cracks, can be seen. The

most important observation is related to the presence of embedded

particles beginning from a surface depth as low as 20 nM. The

elemental maps and high resolution images shown in Fig. 8e,

indicate that metallic Ni is present in these particles, which sizes

are varying from 10 to 100 nM, and larger particles are located

closer to the interface with the Ni3Ti layer.

The ELNES spectra show that the surface oxide for wire 1 is

mostly TiO2 as follows from the doubling of the Ti edge in the upper

spectrum in Fig. 9a, while for other two wires there is also

a contribution from TiO. Indeed, in the case of wire 2, the bottom

part of the oxide layer, i.e. closer to the Ni3Ti layer, reveals TiO (the

middle spectrum in Fig. 9a), while the external layer is TiO2. In wire

3, however, both oxide phases appear to bemixed, as becomes clear

from the broadening in the ELNES (lower spectrum in Fig. 9a). The

difference between both oxides is also obvious from the oxygen

EFTEM map in Fig. 6d. The brighter oxide layer, which has more

oxygen in it (TiO2) is located closer to the surface or to the sputtered

Pt layer. In Fig. 9b a typical Selected Area Electron Diffraction

(SAED) pattern from one of the oxide layers is shown for wire 1. It

reveals primarily TiO2 rings and some TiO rings in this particular

region, indicating the fine grained structure of the oxide. The

reflections that fall next to the labeled rings belong to Ni particles or

grains from Ni3Ti layer also included in the selected aperture.

3.6. Corrosion behavior

It is well known that a heat treatment can cause deterioration of

the corrosion resistance of Nitinol. Using potentiodynamic polari-

zation the corrosion behavior of one of the studied wires, namely

wire 2, which has a more regular surface oxide than wire 1. In the

original state the examined wire revealed a very low breakdown

potential (200 mV). The huge hysteresis loop (Fig. 10, curve 3) that

did not close in the limits of positive potentials indicates the

difficulties in repassivation of the areas with surface damage. The

breakdown potential in saline solution observed in the present

study is significantly lower than the one reported for Ringer’s

solution (738 mV) in Ref. [10]. This result is in agreement with an

earlier comparative study [15] that showed that saline solution is

more aggressive than Ringer’s one. Because of the very poor

measured corrosion resistance, the wires with the original oxides

were not included in the PD test performed under stress. The

treated wires, however, exhibit a passive behavior in the range of

potentials from 0.5 to 1.2 eV with a current density of 10ÿ6 A/cm2

(Fig. 10, curve 2). It seems that there was a surface breakdown at

Fig. 3. XPS survey spectra for wire 1.

Table 1

Elemental concentrations (at.%) on the surfaces of NiTi wires detected with a 45�

collection angle using XPS.

Wire C N O Ti Ni Tl Ti/Ni

1 47.04 1.04 39.29 4.83 7.81 0 0.6

2 37.28 1.88 46.46 12.82 1.06 0.51 12

3 26.57 1.4 51.7 16.66 1.53 0.50 11



1.2 V because the polarization curve did not follow the same path

on the reverse scan. However, the wire with the treated surface

repassivated successfully at a high potential of 1.1 V. The treated

wires were also tested under 3% strain in a tension mode (curve 1).

This latter test was performed immediately after the direct and

reverse cycles of the first PD polarization (unstrained state) were

completed. The strained wire sample shows almost three orders of

magnitude lower current density in the passive region but a lower

breakdown potential (0.8 V) compared to the unstrained samples.

The surface damage on the samples under strain repassivated as

easily as in the case of wire samples with no strain. Although

a positive effect of strain on the corrosion resistance of Nitinol

cannot be excluded [16], it is believed that the large decrease in the

current density observed for the strained wire in the region of

passivity (curve 1) should be assigned rather to a longer exposure

to saline solution that is known to improve Nitinol’s corrosion

resistance [17,18]. Additional studies are needed to draw a conclu-

sion regarding the possibility of improvement in corrosion resis-

tance of Nitinol wires strained in tension mode.

4. Discussion

4.1. Surface oxides and Ni distribution

First we would like to mention the differences between the

elemental concentrations on the wire surfaces observed in the

Fig. 4. The high resolution XPS spectra for the constituents of the surfaces for wires 1 and 2.



present XPS study (Table 1) and those found earlier from Auger

analysis [7]. It is obvious from a comparison of the results that the

metal concentrations on the surface are twice as low in the present

study. We believe that various factors may be involved. First, it

could be the inevitable organic contamination of the surface with

time. Second, the selection of sensitivity factors for the calculation

of elemental concentrations could also contribute to the observed

differences. Disregarding the differences in the absolute values of

the metal concentrations between the present and the earlier study

[7], it is important to note that in both studies wire 1 exhibited the

highest Ni content and wire 3 – the lowest. Additionally, the

concentrations of Ni in the external surface layers that can be

deduced directly from the Auger depth profiles shown in Fig. 5 are

in agreement with those reported in Ref. [7].

As far as the oxide thickness (Table 2) is concerned, the 1 mM2

depth profiles indicate that it can vary significantly from site to site

for the wires 2 and 3 with an intact oxide layer, and it can differ as

much as five times for wire 1 with a damaged surface. Considering

the possibility of a significant variation in the oxide thickness

within the wire samples as well as the effect of surface roughness

on the actual numbers obtained, the oxide thickness for wires 2 and

3 for a 50 mM2 raster are of the same order of magnitude in both

studies, the present and referred [7]. Taking into account a similar

surface topography and roughness [7], the higher Ni release from

wire 3 than from wire 2 can be assigned to a thinner oxide layer.

The thickness of the scale for wire 1 can be as high as 720 nM. The

lower oxide thickness determined from our TEM study may be

partially due to the sample thinning procedure.

A thick oxide layer (scale) in the case of wire 1 is cracked and Ni

particles are detected very close to the surface. It is clear that

a thicker Ti oxide with a higher oxidation degree (Tiþ4) results not

only in the formation of Ni-rich Ni3Ti intermetallics at the interface

between the surface and the bulk of the alloy, but also of pure Ni in

agreement with an earlier XRD study [13]. In the present study we

can visualize the Ni particles and can see clearly how Ni is

distributed in the surface depth, and also can correlate these

observation with the Ni release. As it follows from the present

analysis, the Ni concentration of approximately 8% detected on the

surface of wire 1 is misleading, because there is a reservoir of pure

Ni in close proximity to the surface. The dramatic increase in Ni

release fromwire 1 compared towires 2 and 3 is not only due to the

higher Ni concentration on the surface, but also due to accumulated

Ni deposits and a damaged scale. The gradual increase in Ni release

from the studied wires observed with the duration of their

immersion into the solution is due to the dissolution of external

surface layers and the disclosure of the Ni enriched surface sub-

layers. The saturation in Ni release is achieved when the external

oxide surface layers are corroded away and the interface with the

bulk is exposed.

It was demonstrated in Ref. [10] that after four months

immersion of Nitinol wires with original oxides to saline solution

their breakdown potentials increased from 200 to 800 mV. The

authors assigned this improvement in the corrosion resistance to

the reduced Ni surface content, although they questioned the

underlying mechanism. We believe that a galvanic component of

corrosion is involved. Thus it was shown that the open circuit

potentials of Nitinol can vary in a wide range from þ142 to

ÿ430 mV depending on the surface treatment [9,19], and Ni3Ti has

a significantly higher current density compared to Nitinol samples

[19]. Based on the previous studies [7,10] it was also concluded that

surface treatment of Nitinol wires is necessary to reduce the Ni

content. We would like to make a stronger statement: Ni surface

accumulation must be completely eliminated and a new surface

oxide, homogeneous in chemistry and topography, should be

formed to ensure a non-toxic material’s performance after

implantation in the human body. It is not clear, however, whether

the surface integrity suffered because of stress corrosion cracking

associated with thicker oxides or because of the type of dies

involved in the drawing process. The presence of a deadly toxic

Fig. 5. The elemental Auger depth profiles obtained from wires 1 and 2 using different sizes of collection areas (1 mM2 and 50 mM

2). The Arþ sputtering rate for the surface oxides of

Nitinol wires was approximated by that obtained for a flat SiO2 sample (see details in the text).

Table 2

Thickness of oxide layers on the surface of NiTi wires determined using various

techniques (nm).

Auger 1 mM2

raster

Auger 50 mM2

raster

TEM thinned

samples

Auger [7]

50 mM2 raster

1 80; 440 720 160–190 340 (SEM

[7])

2 120; 300 220 50–100 170

3 36; 72 84 25–50 120



element like thallium on the surfaces of wires 2 and 3 is not

acceptable.

Another avenue that can bring us to the same final results is

associatedwith seemingly low temperatures in the interval from 60

to 160 �C that are commonly employed as pre-treatment protocols

for a coating deposition, as developed originally for pure Ti. These

temperatures are not high enough to expect significant atomic

diffusion through vacancies that is needed for new phase growth.

However, Ni atoms are liberated from the Ni–Ti interatomic bonds

already at room temperature because a thin Ti-based oxide is

formed spontaneously on the Nitinol surface. These Ni atoms are

lattice defects, interstitial atoms in the structure of the Ti-surface

oxides. Besides, due to a shortage of oxygen atoms, Ti-surface

oxides are generally non-stoichiometric. The sites of the missing

oxygen atoms present structural oxygen vacancies that can be used

by Ni atoms to migrate through Ti oxides. Thus, the diffusion of Ni

atoms through surface oxides can occur effectively already at low

temperatures. Furthermore, the smaller size of Ni atoms (as

compared with Ti and oxygen atoms) also helps Ni diffusion

through an interstitial path.

Discussing the sample preparation procedure it is important to

mention that the Nitinol surface chemistry can also be easily

Fig. 6. TEM image (a) and the elemental maps for Ni (b), Ti (c) and oxygen (d) for wire 2.

Fig. 7. Experimental and simulated [212] nanodiffraction pattern of a grain in the Ni3Ti sublayer (note that the fine 10–1 superlattice reflections can be seen in both patterns).



altered by electro-dischargemachining thatmay cause damage into

surface depth up to 22 mM [20].

4.2. Biological implications

The accumulation of Ni on the surface and at the interfaces is

a cause of enhanced and lasting release that unavoidably should

affect the biological performance of implant devices. For instance,

the biocompatibility of stents can be improved not only through an

optimization of the design and the mechanical properties that

would minimize surgical trauma but also by avoiding metal ion

release, thus reducing the inflammatory response. Studies

involving coronary self-expanding stents in animals and humans

[21,22] have concluded that Nitinol stents continuously expand

after placement. In a six months follow-up period, the increase in

the Nitinol stent diameter might reach 21–24%, which is attractive

considering the possibility of restenosis (repeated narrowing or

closure of lumen). However, the self-expanding mechanism of

Nitinol that may reduce a trauma from stent deployment did not

result in a lower intimal proliferation. On the contrary, an increased

intimal hyperplasia was observed. For instance, the proliferative

response was maximal after three to six months of implantation,

while a considerable reduction occurred only after a 12 months

period [23]. Due to the interplay between a follow-up increase in

Fig. 8. TEM image (a) and the elemental maps for Ni (b), Ti (c) and oxygen (d) for wire 1 (the arrows point at metallic Ni particles). HRTEM image (e) of an inclusion revealing the Ni

face centered cubic (fcc) lattice observed in [001] orientation.



stent diameter and neointimal proliferation, the total late lumen

loss was yet smaller for Nitinol when compared to stainless steel,

even though it was not statistically significant [22]. It is not clear at

the moment whether continuing self-expansion provokes an

excessive SMC proliferation over that of balloon-expandable stents

or whether this is the result of lasting Ni release. Thus several other

studies [24,25] contrary to [21,22] showed no increase of neo-

intimal proliferation of self-expandable stents compared with rigid

ones. Among the factors that could stimulate neointimal hyper-

plasia, the hypersensitivity to stent materials was mentioned [22].

We believe that the increasing Ni release fromNitinol surfaces with

the buried Ni-rich layers which were not eliminated during surface

preparation with a peak (or saturation) at six months exposure in

vitro could induce a toxic effect in vivo, while significantly

contributing to intimal proliferation in the studies [21,22].

Importantly, it is sometimes mentioned in Nitinol literature that

‘the ceramic-like nature of the titanium-dioxide surface layer on

Nitinol must resist cracking in the pulsative oscillation and wear on

contact points with other materials’ [26,27]. It should be

emphasized that the ceramic nature of TiO2 contradicts both

Nitinol superelasticity and the cyclic pulsation of blood vessels.

Ceramic materials are held together by ionic and covalent bonds

which tend to fracture before any plastic deformation even in

traditional metallic materials occurs, i.e. at strain level significantly

lower than 1%. Superelastic Nitinol medical devices, on the other

hand, are expected to perform at least at a 2–4% strain.

5. Conclusions

The present study demonstrated that the use of Nitinol wires

with original oxides for long-term implantation is not necessarily

safe even if the external surface layers are adjusted to relatively low

Ni concentrations. Also, a long and lasting Ni release is always an

indication of the presence of Ni-rich buried surface sublayers

serving as a permanent Ni reservoir and a possible cause of reduced

corrosion resistance. The provided analysis also indicates that the

chemistry of Nitinol surfaces formed during heat treatments should

be scrutinized through evaluation of its homogeneity, and the

conclusions must be supported by numerous data from surface

analysis rather than based on a single depth profile commonly

associated with the evaluation of Nitinol surfaces for medical

devices.

Acknowledgements

The Research Fund of K.U. Leuven is acknowledged for a partial

financial support. Part of this work was also funded by the National

Science Foundation of Flanders under the project G.0465.05 ‘The

functional properties of shape memory alloys: a fundamental

approach.’ This manuscript has been also authored by Iowa State

University of Science and Technology under Contract No. DE-AC02-

07CH11358 with the U.S. Department of Energy. The authors also

thankful to G. Rondelli for the assistance with the corrosion tests.

One of the authors appreciates a productive discussion with M.

Rettenmayr.

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