Binary Zn-Ti alloys for orthopedic applications: Corrosion and degradation behaviors, friction and wear performance, and cytotoxicity

Abstract

Zinc (Zn) and its biocompatible and biodegradable alloys have substantial potential for use in ortho-pedic implants. Nevertheless, pure Zn with a hexagonal close-packed crystal structure has only two independent slip systems, therefore exhibiting extremely low elongation and yield strength in its as-cast condition, which restricts its clinical applications. In this study, as-cast Zn-xTi (titanium) (x = 0.05, 0.10, 0.20, and 0.30 wt.%) binary alloys were hot-rolled and their microstructures, mechanical properties , wear resistance, and cytocompatibility were comprehensively investigated for orthopedic implant applications. The microstructures of both as-cast and hot-rolled Zn-xTi alloys consisted of an ␣-Zn matrix phase and a TiZn 16 phase, while Zn-0.2Ti and Zn-0.3Ti exhibited a finer ␣-Zn phase due to the grain-refining effect of Ti. The hot-rolled Zn-0.2Ti alloy exhibited the highest yield strength (144.5 MPa), ultimate strength (218.7 MPa), and elongation (54.2%) among all the Zn-xTi alloys. The corrosion resistance of Zn-xTi alloys in Hanks' solution decreased with increasing addition of Ti, and the hot-rolled Zn-0.3Ti alloy exhibited the highest corrosion rates of 432 m/y as measured by electrochemical testing and 57.9 m/y as measured by immersion testing. The as-cast Zn-xTi alloys showed lower wear losses than their hot-rolled counterparts. The extracts of hot-rolled Zn-xTi alloys at concentrations of ≤ 25% showed no cytotoxicity to MG-63 osteosarcoma cells and the extracts of Zn-xTi alloys exhibited enhanced cytocompatibility with increasing Ti content.

Full text

Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Contents
lists
available
at
ScienceDirect
Journal
of
Materials
Science
&
Technology
jo
ur
na
l
homepage:
www.jmst.org
Research
Article
Binary
Zn–Ti
alloys
for
orthopedic
applications:
Corrosion
and
degradation
behaviors,
friction
and
wear
performance,
and
cytotoxicity
Kun
Wanga,b,1,
Xian
Tongc,a,∗,1,
Jixing
Lina,∗,
Aiping
Weia,
Yuncang
Lid,
Matthew
Dargusche,
Cuie
Wend,∗∗
aDepartment
of
Material
Engineering,
Zhejiang
Industry
&
Trade
Vocational
College,
Wenzhou
325003,
China
bKey
Laboratory
of
Materials
Physics,
Institute
of
Solid
State
Physics,
Chinese
Academy
of
Sciences
Hefei,
CN
230031,
China
cSchool
of
Materials
Science
and
Engineering,
Xiangtan
University,
Xiangtan
411105,
China
dSchool
of
Engineering,
RMIT
University,
Melbourne,
Victoria
3001,
Australia
eCentre
for
Advanced
Materials
Processing
and
Manufacturing
(AMPAM),
The
University
of
Queensland,
Brisbane,
Queensland
4072,
Australia
a
r
t
i
c
l
e
i
n
f
o
Article
history:
Received
30
June
2020
Received
in
revised
form
28
August
2020
Accepted
12
September
2020
Available
online
20
October
2020
Keywords:
Cytotoxicity
Degradation
behavior
Mechanical
property
Wear
resistance
Zn–Ti
alloy
a
b
s
t
r
a
c
t
Zinc
(Zn)
and
its
biocompatible
and
biodegradable
alloys
have
substantial
potential
for
use
in
ortho-
pedic
implants.
Nevertheless,
pure
Zn
with
a
hexagonal
close-packed
crystal
structure
has
only
two
independent
slip
systems,
therefore
exhibiting
extremely
low
elongation
and
yield
strength
in
its
as-
cast
condition,
which
restricts
its
clinical
applications.
In
this
study,
as-cast
Zn–xTi
(titanium)
(x
=
0.05,
0.10,
0.20,
and
0.30
wt.%)
binary
alloys
were
hot-rolled
and
their
microstructures,
mechanical
proper-
ties,
wear
resistance,
and
cytocompatibility
were
comprehensively
investigated
for
orthopedic
implant
applications.
The
microstructures
of
both
as-cast
and
hot-rolled
Zn–xTi
alloys
consisted
of
an
␣-Zn
matrix
phase
and
a
TiZn16 phase,
while
Zn–0.2Ti
and
Zn–0.3Ti
exhibited
a
finer
␣-Zn
phase
due
to
the
grain-
refining
effect
of
Ti.
The
hot-rolled
Zn–0.2Ti
alloy
exhibited
the
highest
yield
strength
(144.5
MPa),
ultimate
strength
(218.7
MPa),
and
elongation
(54.2%)
among
all
the
Zn–xTi
alloys.
The
corrosion
resis-
tance
of
Zn–xTi
alloys
in
Hanks’
solution
decreased
with
increasing
addition
of
Ti,
and
the
hot-rolled
Zn–0.3Ti
alloy
exhibited
the
highest
corrosion
rates
of
432
␮m/y
as
measured
by
electrochemical
testing
and
57.9
␮m/y
as
measured
by
immersion
testing.
The
as-cast
Zn–xTi
alloys
showed
lower
wear
losses
than
their
hot-rolled
counterparts.
The
extracts
of
hot-rolled
Zn–xTi
alloys
at
concentrations
of
≤
25%
showed
no
cytotoxicity
to
MG-63
osteosarcoma
cells
and
the
extracts
of
Zn–xTi
alloys
exhibited
enhanced
cytocompatibility
with
increasing
Ti
content.
©
2020
Published
by
Elsevier
Ltd
on
behalf
of
The
editorial
office
of
Journal
of
Materials
Science
&
Technology.
1.
Introduction
Biodegradable
metal
implant
materials
with
satisfactory
bio-
compatibility
can
degrade
gradually
in
vivo,
elicit
suitable
host
responses,
and
dissolve
completely
after
tissue
healing
[1,2].
The
main
biodegradable
metals
are
magnesium
(Mg)-,
iron
(Fe)-,
and
zinc
(Zn)-based
alloys
and
composites.
Currently,
Mg-
and
Fe-based
∗Corresponding
authors
at:
Department
of
Material
Engineering,
Zhejiang
Indus-
try
&
Trade
Vocational
College,
Wenzhou
325003,
China.
∗∗ Corresponding
author.
E-mail
addresses:
tx847595271@163.com
(X.
Tong),
linjixing@163.com
(J.
Lin),
cuie.wen@rmit.edu.au
(C.
Wen).
1Joint
first
author.
metal
implants
have
entered
the
stage
of
clinical
medical
applica-
tions;
in
particular,
Mg-based
biomaterials
have
found
applications
in
bone-fixation
and
stent
devices
[3–7].
Ezechieli
et
al.
[5]
reported
that
MgYREZr
(MAG-NEZIX®and
Magmaris
by
Biotronik)
has
been
successfully
commercialized
in
medical
screws.
Haude
et
al.
[4]
reported
that
the
mean
in-segment
late
lumen
loss
was
reduced
at
the
end
of
6
months
after
the
implantation
of
a
second-generation
drug-eluting
absorbable
metal
scaffold
(DREAMS
2
G)
(Mg-based)
into
the
coronary
target
lesions,
and
no
definite
or
probable
scaf-
fold
thrombosis
was
observed.
For
Fe-based
implants,
Wu
et
al.
[8]
reported
that
pure
Fe
stents
showed
no
thrombosis,
inflam-
mation,
or
necrosis
after
being
implanted
into
healthy
mini-swine
for
4
weeks.
Lin
et
al.
[9]
reported
that
an
Fe-based
drug-eluting
coronary
scaffold
(IBS
scaffold)
coated
with
a
Zn
barrier
layer
and
a
poly(D,
l-lactide)
(PDLLA)
coating
showed
a
significantly
shortened
https://doi.org/10.1016/j.jmst.2020.10.031
1005-0302/©
2020
Published
by
Elsevier
Ltd
on
behalf
of
The
editorial
office
of
Journal
of
Materials
Science
&
Technology.

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
corrosion
period
and
smaller
amounts
of
corrosion
products
with-
out
causing
any
biological
problems
after
implantation
in
a
rabbit
abdominal
aorta
for
up
to
13
months.
Although
both
Mg
and
Fe
play
important
roles
in
the
metabolic
processes
of
the
human
body
as
essential
elements,
and
Mg-
and
Fe-based
biomaterials
have
good
biocompatibility
and
high
mechanical
strength
[10–12],
these
materials
have
various
disadvantages.
Mg-based
alloys
degrade
too
fast,
which
leads
to
the
generation
of
excessive
hydrogen
(H2)
gas
bubbles
and
inadequate
mechanical
integrity
before
sufficient
heal-
ing
[13,14].
On
the
other
hand,
Fe-based
alloys
degrade
too
slow,
which
can
severely
affect
the
natural
bone
growth
and
healing
[15].
Zn
and
its
alloys
have
attracted
great
interest
as
potential
biodegradable
medical
materials
due
to
their
moderate
degra-
dation
rates
and
acceptable
cytocompatibility
[16,17].
Zn
plays
an
essential
role
in
physiological
activities
and
in
all
six
enzyme
classes
as
one
of
the
essential
trace
elements,
such
as
participa-
tion
in
the
metabolism
of
RNA
and
DNA,
signal
transduction,
gene
expression,
and
regulation
in
apoptosis
[18–20].
In
addition,
Zn
also
has
an
osteoinductive
effect
on
bone
tissue
[21].
However,
pure
Zn
has
low
strength
and
poor
ductility,
which
hinders
its
application
in
medical
implant
applications
[22,23].
Furthermore,
pure
Zn
has
a
low
recrystallization
temperature,
which
may
cause
microstructural
and
mechanical
instability,
affecting
its
reliabil-
ity
in
medical
applications.
Micro-alloying
is
one
of
the
strategies
for
improving
the
mechanical
properties
of
Zn-based
alloys
such
as
the
addition
of
low
concentrations
of
Mg
[24],
strontium
(Sr)
[24],
calcium
(Ca)
[24],
copper
(Cu)
[25],
manganese
(Mn)
[26],
sil-
ver
(Ag)
[27],
Fe
[28],
lithium
(Li)
[29],
and
germanium
(Ge)
[30],
among
others.
Thermomechanical
processing
is
another
strategy
for
enhancing
mechanical
properties,
including
rolling
[24,29,30],
extrusion
[24,26],
drawing
[27],
and
equal
channel
angular
pressing
(ECAP)
[31],
among
others.
Titanium
(Ti)
and
some
of
its
alloys
are
considered
ideal
biomedical
implant
materials
due
to
their
high
corrosion
resis-
tance,
mechanical
properties,
and
biocompatibility,
and
are
widely
used
in
orthopedic
internal
fixation
surgery
[32].
In
addition,
Ti
is
commonly
used
as
an
alloying
element
to
refine
the
grain
size
of
Mg
alloys
[33].
Yin
[34]
comprehensively
investigated
the
microstructure,
mechanical
properties,
and
fractography
of
as-cast
and
hot-extruded
Zn–xTi
(x
=
0.01,
0.1,
and
0.3
wt.%)
alloys
for
biodegradable
stent
applications,
and
demonstrated
that
the
hot-
extruded
Zn–0.1Ti
alloy
showed
the
best
combination
of
tensile
properties,
with
an
ultimate
tensile
strength
of
207
MPa,
yield
strength
of
163
MPa,
and
elongation
of
44%.
However,
the
corro-
sion
and
degradation
behaviors,
friction
and
wear
performance,
and
biocompatibility
of
these
Zn–Ti
alloys
were
not
tested.
Li
et
al.
[35]
reported
that
a
high
concentration
of
Ti
ions
was
toxic
to
osteoblast-
like
SaOS2cells
and
the
threshold
of
Ti
ion
concentration
was
15.5
␮g/l.
Therefore,
Ti-containing
biodegradable
metals
must
be
tested
for
their
cytotoxicity
and
biocompatibility
for
biomedical
applica-
tions.
Hence,
the
cell
viability
of
Ti-containing
Zn
alloys
must
be
assessed
in
the
development
of
mechanically
compatible,
biocom-
patible,
and
biofunctional
implant
materials.
In
this
study,
Zn–xTi
(x
=
0.05,
0.1,
0.2,
and
0.3
wt.%)
binary
alloys
were
prepared
by
casting
and
hot-rolling,
and
their
corrosion
and
degradation
properties,
friction
and
wear
performance,
and
in
vitro
cytotoxicity
were
comprehensively
assessed
to
determine
their
suitability
for
biodegradable
implant
material
applications.
2.
Experimental
procedure
2.1.
Material
preparation
Zn–xTi
(x
=
0.05,
0.1,
0.2,
and
0.3
wt.%,
hereafter
denoted
Zn–0.05Ti,
Zn–0.1Ti,
Zn–0.2Ti,
and
Zn–0.3Ti,
respectively)
alloy
ingots
were
prepared
via
gravity
casting
by
melting
pure
Zn
(purity
99.99%,
Huludao
Zinc
Industry
Co.,
China)
and
a
Zn–3Ti
master
alloy
(Hunan
High
Broad
New
Material
Co.,
China)
at
560 ◦C
under
an
argon
(Ar)
atmosphere.
The
melted
metals
were
poured
into
a
steel
mold
preheated
to
200 ◦C
for
solidification.
The
as-cast
(AC)
ingots
were
homogenized
at
340 ◦C
for
10
h
before
hot-rolling,
followed
by
quenching
in
air.
Plates
with
dimensions
of
60
mm
×
20
mm
×
10
mm
were
cut
via
electrical
discharge
machining
(EDM)
for
hot-rolling.
The
plates
were
rolled
to
a
final
thickness
of
2
mm
at
a
rolling
reduction
of
0.5
mm
per
pass
after
each
preheating
at
200
◦C
for
5
min.
The
chemical
compositions
of
the
Zn–xTi
metal
ingots
were
determined
via
X-ray
fluorescence
(XRF;
S4
Pioneer,
Bruker,
Germany).
The
Ti
content
was
0.054
±
0.011
wt.%
for
Zn–0.05Ti,
0.114
±
0.020
wt.%
for
Zn–0.1Ti,
0.205
±
0.038
wt.%
for
Zn–0.2Ti,
and
0.321
±
0.067
wt.%
for
Zn–0.3Ti.
For
electrochemical,
immersion,
and
in
vitro
cytotoxicity
tests,
disc
samples
of
the
Zn–xTi
alloys
with
a
diameter
of
8
mm
and
a
thickness
of
2
mm
were
cut
from
the
AC
ingots
and
hot-rolled
(HR)
plates.
All
the
disc
samples
were
polished
up
to
2000
grit
using
sil-
icon
carbide
(SiC)
abrasive
grinding
papers,
followed
by
ultrasonic
cleaning
in
ethanol
for
20
min
and
drying
in
air
prior
to
testing.
2.2.
Microstructure
characterization
The
microstructures
of
the
Zn–xTi
alloys
were
observed
via
optical
microscopy
(OM;
DM2500C,
Leica,
Germany)
and
scan-
ning
electron
microscopy
(SEM;
Pro
X
FEI,
Phenom,
Netherlands)
after
polishing
to
1
␮m
with
diamond
paste
and
etching
with
0.5%
HNO3/alcohol
solution,
and
the
AC
Zn–xTi
alloy
samples
were
cut
from
the
same
area
(i.e.,
near
the
bottom)
of
the
ingots
by
elec-
trical
discharge
machining
(EDM).
The
chemical
compositions
of
various
phases
were
determined
via
energy-dispersive
spectrom-
etry
(EDS;
X-Max,
Oxford,
England)
at
15
keV.
The
microstructure
of
AC
Zn–0.3Ti
alloy
was
also
observed
via
transmission
electron
microscopy
(TEM)
(Talos
F200X,
FEI,
USA)
at
200
keV.
TEM
thin
film
samples
with
a
diameter
of
3
mm
were
ground
and
argon-
ion
milled
using
a
precision
ion
polishing
system
(PIPS,
Gatan
691,
USA).
The
average
grain
size
and
second-phase
content
were
mea-
sured
using
Image
Pro-Plus
software
(Image
Pro-Plus
6.0;
Media
Cybernetics,
Silver
Spring,
MD,
USA).
The
grain
size
distributions
of
the
HR
samples
were
examined
by
a
field
emission
scanning
electron
microscope
(FE-SEM,
Scios
DualBeam,
FEI,
USA)
equipped
with
an
electron
backscatter
diffraction
device
(EBSD)
(HKL
Chan-
nel
5
system,
Oxford,
England)
at
20
keV.
The
EBSD
samples
were
ground
and
electrolytically
polished
with
a
solution
of
7%
perchlo-
ric
acid
and
93%
alcohol
at
a
current
density
of
0.45
A/cm2and
a
voltage
of
20
V
for
1
min.
The
constituent
phases
were
identified
using
X-ray
diffractometry
(XRD;
D8
ADVANCE,
Bruker,
Germany)
with
Cu–Ka
radiation
at
a
scan
speed
of
2◦/min
and
a
scan
step
of
0.02◦over
a
2␪
range
from
10
to
90◦.
2.3.
Mechanical
testing
Tensile
tests
of
the
Zn–xTi
alloys
were
performed
on
a
univer-
sal
material
testing
machine
(Instron
3369,
USA)
with
a
crosshead
speed
of
1
mm/min
at
room
temperature.
The
tensile
samples
with
a
gage
length
of
25
mm
and
a
width
of
6
mm
were
prepared
via
EDM
according
to
ASTM
E8/E
8M-16
[36]
and
the
HR
samples
were
cut
parallel
to
the
rolling
direction
from
the
plate.
The
fracture
sur-
faces
of
the
tensile
samples
after
testing
were
observed
by
SEM
(Pro
X
FEI,
Phenom,
Netherlands).
Vickers
hardness
of
the
Zn–xTi
alloys
was
determined
using
a
micro-hardness
tester
(MicroMet
6000,
Buehler,
America)
with
an
applied
load
of
100
g
and
a
loading
time
of
10
s.
217

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
2.4.
Electrochemical
testing
Electrochemical
tests
of
the
Zn–xTi
alloys
were
conducted
using
an
electrochemical
workstation
(ParStat
2273,
Princeton,
USA)
in
Hanks’
solution
with
a
pH
value
of
7.40
at
37 ◦C.
A
typical
three-
electrode
cell
system
was
used
for
electrochemical
measurements,
including
a
disc
sample
with
an
exposed
area
of
0.5
cm2as
the
working
electrode,
a
saturated
calomel
electrode
(SCE)
as
the
ref-
erence
electrode,
and
a
platinum
electrode
as
a
counter
electrode
(CE).
Potentiodynamic
polarization
tests
were
performed
over
a
potential
range
from
−0.3
V
to
0.8
V
vs.
open
circuit
potential
(OCP)
at
a
scanning
rate
of
1
mV/s
after
reaching
a
stable
OCP
value.
Linear
polarization
resistance
(LPR)
of
the
samples
was
mea-
sured
at
a
potential
scan
range
of
±
20
mV
vs.
OCP
with
a
potential
rate
of
10
mV/min
according
to
the
method
described
in
a
previous
study
[37].
For
the
samples
without
Tafel
region
in
the
anodic
branches,
the
corrosion
current
density
(Icorr)
was
determined
using
the
equation
given
by:
Icorr=
bc/(2.303×Rp)
(1)
where
Rpis
the
polarization
resistance
obtained
via
the
LPR
method
and
bcis
the
cathodic
Tafel
slope.
The
corrosion
rate
(Vcorr)
of
the
samples
was
calculated
accord-
ing
to
the
method
described
by
Liu
et
al.
[37],
which
can
be
expressed
by
the
following
equation:
Vcorr=
(3.15×105)×(a×Icorr)/(N×F×D)
(2)
where
a
is
the
atomic
weight
(g/mol)
of
the
samples,
N
is
the
num-
ber
of
equivalent
exchange,
F
is
the
Faraday’s
constant,
and
D
is
the
density
(g/cm3)
of
the
samples.
2.5.
Immersion
testing
Immersion
tests
were
performed
by
immersing
the
Zn–xTi
alloy
samples
in
Hanks’
solution
at
37 ◦C
for
30
d
with
a
ratio
of
surface
area
to
solution
volume
of
1
cm2/20
mL
according
to
the
Standard
ASTM
G31–72
[38].
The
morphology
and
composition
of
the
cor-
rosion
products
on
the
sample
surfaces
after
immersion
tests
were
characterized
by
SEM
and
EDS.
The
constituent
phases
of
the
cor-
rosion
products
after
immersion
tests
were
identified
by
a
XRD
(D/max
2500,
Rigaku,
Japan)
with
Cu–K␣
radiation
and
the
scan-
ning
spectra
were
recorded
over
a
2␪
range
of
10–90◦at
a
scan
speed
of
2◦min−1.
A
chromic
acid
solution
composed
of
200
g/L
chromium
trioxide
(CrO3)
and
10
g/L
silver
nitrate
(AgNO3)
was
used
to
remove
the
corrosion
products
on
the
corroded
surfaces
of
the
samples,
and
the
samples
were
subsequently
cleaned
with
alcohol
and
dried
in
air.
The
weights
of
the
samples
before
and
after
the
immersion
corrosion
tests
were
measured
using
an
electronic
analytical
balance
with
an
accuracy
of
0.1
mg.
The
degradation
rate
was
calculated
via
weight
loss
measurements
according
to
ASTM
G31−72
[38].
2.6.
Friction
and
wear
testing
Friction
and
wear
tests
of
the
Zn–xTi
alloys
were
conducted
in
a
high-speed
reciprocating
friction
tester
(HSR-2
M;
Lanzhou
Zhongke
Kaihua
Technology
Development
Co.,
China).
Disc
samples
with
15
mm
diameter
and
10
mm
thickness
were
mounted
with
epoxy
resin,
polished
to
0.5
␮m
using
a
diamond
paste,
and
cleaned
with
ethanol
before
wear
testing.
Wear
tests
were
performed
at
a
load
of
200
gf,
a
friction
stroke
length
of
5
mm,
a
friction
pair
of
zir-
conia
(ZrO2)
with
3
mm
diameter,
and
a
friction
frequency
of
1
Hz
for
10
min.
Corrosion
wear
tests
were
carried
out
by
immersing
the
samples
in
a
sink
containing
50
mL
of
Hanks’
solution
at
a
constant
temperature
of
37 ◦C.
The
detailed
procedure
for
the
method
was
Fig.
1.
XRD
patterns
of
AC
and
HR
Zn–xTi
alloys.
described
in
our
previous
study
[39].
The
wear
loss
was
calculated
from
the
weights
of
the
samples
before
and
after
the
friction
and
wear
tests.
The
morphology
and
compositions
of
the
sample
sur-
faces
after
friction
and
wear
tests
were
characterized
by
SEM
(Pro
X
FEI,
Phenom,
Netherlands)
and
EDS.
2.7.
Cytotoxicity
evaluation
assay
Cytotoxicity
assay
of
the
HR
Zn–xTi
alloys
was
conducted
using
the
indirect
method
according
to
ISO
10993−5:
2009
[40]
and
the
detailed
procedure
was
described
in
our
previous
study
[39].
Briefly,
human
osteoblast-like
MG-63
cell
lines
(GDC074,
Shanghai
Cell
Bank
of
Chinese
Academy
of
Sciences,
China)
were
used
for
the
cytotoxicity
assay.
The
cells
were
cultured
in
0.1
mM
alpha-
minimum
essential
medium
(␣-MEM;
Invitrogen,
Carlsbad,
USA)
supplemented
with
10%
fetal
bovine
serum
(FBS)
(Cat
Number:
S711−001S.
Lonsera,
Lonsa
Science
SRL,
Uruguay)
and
antibiotics
(100
units/mL
penicillin
and
100
␮g/mL
streptomycin,
Cat
Num-
ber:
PH1513,
Aladdin,
Shanghai,
China)
at
37 ◦C
in
a
humidified
atmosphere
with
5%
CO2.
The
extracts
of
the
HR
Zn–xTi
alloys
were
prepared
by
immersing
the
samples
in
the
culture
medium
with
an
extraction
ratio
of
1.25
cm2/mL
and
cultured
for
3
d.
MG-63
cells
with
a
density
of
5
×
103cells
per
well
were
seeded
into
a
96-well
culture
plate
and
incubated
for
1
d
to
allow
for
attachment.
Then
the
culture
medium
was
replaced
by
alloy
extracts
with
concentrations
of
100%,
25%,
and
12.5%,
and
incubated
for
1
day.
A
cell-counting
solution
kit
(CCK–8,
10
␮L,
Dojindo,
Japan)
was
added
to
each
well
and
the
plate
was
incubated
continually
for
2
h
under
the
same
conditions
after
replacing
the
original
culture
medium
with
fresh
culture
medium.
Finally,
the
absorbance
of
each
well
was
measured
using
a
plate
reader
(Gen5,
Biotek,
USA)
at
450
nm.
2.8.
Statistical
analysis
In
this
study,
the
experimental
data
is
expressed
as
the
mean
±
standard
deviation.
One-way
analysis
of
variance
(ANOVA,
V17.0,
SPSS
Inc.,
USA)
was
used
to
determine
the
significance
of
the
differences
observed
between
groups.
P
<
0.05
was
accepted
as
statistically
significant.
218

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Fig.
2.
Optical
micrographs
of
AC
Zn–xTi
alloys:
(a)
Zn–0.05Ti;
(b)
Zn–0.1Ti;
(c)
Zn–0.2Ti;
and
(d)
Zn–0.3Ti.
Fig.
3.
Optical
micrographs
of
HR
Zn–xTi
alloys:
(a)
Zn–0.05Ti;
(b)
Zn–0.1Ti;
(c)
Zn–0.2Ti;
and
(d)
Zn–0.3Ti.
3.
Results
3.1.
Microstructure
of
Zn–xTi
Fig.
1
shows
XRD
patterns
of
the
AC
and
HR
Zn–xTi
alloys.
All
of
the
Zn–xTi
alloy
samples
were
mainly
composed
of
an
␣-Zn
matrix
phase
and
an
intermetallic
TiZn16 phase.
The
intensity
of
the
diffraction
peaks
of
the
TiZn16 phase
increased
with
increas-
ing
Ti
addition;
hence,
the
amount
of
the
TiZn16 phase
increased
gradually
with
the
increasing
Ti
addition.
Fig.
2
shows
optical
micrographs
of
the
AC
Zn–xTi
alloys.
These
alloys
were
mainly
composed
of
a
white
␣-Zn
matrix
phase
and
a
grey
TiZn16 phase
that
was
distributed
along
the
grain
boundaries
of
the
␣-Zn
phase.
The
average
grain
size
was
86.4
±
10.9
␮m
for
219

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Fig.
4.
EBSD
images
and
corresponding
grain
size
distributions
of
HR
Zn–xTi
alloys:
(a)
and
(a’)
Zn–0.05Ti;
(b)
and
(b’)
Zn–0.1Ti;
(c)
and
(c’)
Zn–0.2Ti;
(d)
and
(d’)
Zn–0.3Ti.
Zn–0.05Ti,
71.2
±
12.3
␮m
for
Zn–0.1Ti,
19.4
±
1.9
␮m
for
Zn–0.2Ti,
and
46.9
±
8.2
␮m
for
Zn–0.3Ti.
It
can
be
seen
that
with
increasing
addition
of
Ti,
the
average
grain
size
of
the
␣-Zn
phase
in
the
Zn–xTi
alloys
decreased
initially
and
subsequently
increased,
exhibiting
the
finest
grain
size
of
∼19.4
␮m
with
0.2%
addition
of
Ti.
Therefore,
a
suitable
level
of
Ti
addition
can
significantly
refine
the
␣-Zn
phase
in
Zn–xTi
alloys,
while
excessive
addition
of
Ti
weakens
the
grain-
refining
effect
on
the
␣-Zn
phase.
In
addition,
with
increasing
Ti
addition
in
the
AC
Zn–xTi
alloys,
the
TiZn16 phase
became
longer
and
its
volume
fraction
increased.
The
size
and
volume
fraction
of
the
TiZn16 phase
were
2.3
±
0.4
␮m
and
11.8
±
0.9%
for
Zn–0.05Ti,
3.8
±
0.7
␮m
and
18.3
±
1.4%
for
Zn–0.1Ti,
4.0
±
0.5
␮m
and
23.4
±
1.2%
for
Zn–0.2Ti,
and
9.6
±
1.8
␮m
and
29.6
±
1.8%
for
Zn–0.3Ti,
respectively.
Fig.
3
shows
optical
micrographs
of
the
HR
Zn–xTi
alloys.
Fig.
4
shows
the
representative
EBSD
images
of
the
HR
Zn–xTi
alloys.
After
hot-rolling,
both
the
␣-Zn
phase
and
intermetallic
compound
TiZn16 phase
were
elongated
and
distributed
along
the
rolling
direction.
As
the
Ti
addition
increased,
the
grain
size
of
the
␣-Zn
phase
in
the
HR
Zn–xTi
alloys
decreased
initially
and
subsequently
increased,
and
the
TiZn16 strips
became
longer.
The
HR
Zn–0.2Ti
alloy
showed
the
smallest
grain
size
among
the
four
HR
alloys.
In
addition,
the
HR
Zn–0.3Ti
alloy
exhibited
coarse
TiZn16 lumps
in
addition
to
the
elongated
TiZn16 strips,
and
an
increase
in
its
content
than
the
HR
Zn–0.2Ti
alloy
(Figs.
4c
and
4d).
Fig.
5
shows
SEM
images
and
EDS
analysis
results
of
the
AC
and
HR
Zn–xTi
alloys.
The
corresponding
EDS
spot
analysis
results
for
the
different
phases
(spots
marked
in
Figs.
5d
and
5
h)
of
the
AC
and
HR
Zn–0.3Ti
alloys
are
shown
in
Fig.
5(i).
A
substantial
amount
of
the
long
strip-like
phase
(marked
spot
2)
distributed
in
the
matrix
(marked
spot
1)
is
observed
in
the
AC
Zn–0.3Ti
alloy
(Fig.
5d).
A
coarse
lumpy
phase
(marked
spot
5)
and
a
long
strip-
like
phase
(marked
spot
4)
distributed
in
the
matrix
(marked
spot
3)
are
observed
in
the
HR
Zn–0.3Ti
alloy
(Fig.
5h),
and
the
amount
of
the
long
strip-like
phase
is
less
than
that
in
the
AC
Zn–0.3Ti
alloy.
In
addition,
many
cracks
are
clearly
observable
in
the
coarse
lumpy
phase.
Fig.
5(i)
shows
EDS
analysis
results
for
the
various
spots
marked
in
Figs.
5(d)
and
5(h).
Spot
1
contains
Zn
and
a
small
amount
of
Ti,
and
the
relative
atomic
content
of
Ti
is
0.1
±
0.0%.
Spot
2
is
the
long
strip-like
second
phase,
which
contains
Ti
and
Zn,
and
the
relative
atomic
content
of
Ti
is
2.9
±
0.1%.
Spot
3
is
the
matrix
phase
in
Fig.
5(d),
which
contains
a
small
amount
of
Ti
with
a
relative
atomic
content
of
0.2
±
0.0%.
Spot
4
is
the
strip-like
second
phase,
which
contains
Ti
and
Zn,
and
the
relative
atomic
content
of
Ti
is
2.8
±
0.2%.
Spot
5
is
the
coarse
lumpy
phase,
which
contains
Ti
and
Zn,
and
the
relative
atomic
content
of
Ti
is
5.5
±
0.1%.
Accord-
ing
to
the
Zn–Ti
phase
diagram
[41]
and
the
XRD
patterns
in
Fig.
1,
the
matrix
phase
in
spot
1
and
spot
3
is
the
␣-Zn
phase.
Spot
2
and
spot
4
contain
small
amounts
of
Ti,
less
than
the
Ti
content
in
the
TiZn16 phase.
The
EDS
spot
scans
might
include
the
compositions
of
both
the
matrix
phase
and
the
strip
phase
due
to
their
small
size,
thereby
resulting
in
the
atomic
content
of
the
Ti
in
the
strip
phase
being
lower
than
that
in
the
TiZn16 phase,
so
the
strip
phase
can
be
inferred
as
the
TiZn16 phase.
Spot
5
is
a
Ti-rich
second-phase
particle
and
the
atomic
ratio
of
Ti
and
Zn
is
approximately
1:16,
so
it
can
be
inferred
that
this
phase
is
a
TiZn16 phase.
Fig.
5(j)
shows
the
EDS
mapping
result
of
the
HR
Zn–0.3Ti
alloy.
The
EDS
mapping
result
of
the
HR
Zn–0.3Ti
alloy
mainly
contained
Zn
and
Ti
elements
(Fig.
5(h)).
The
majority
of
Ti
element
is
distributed
in
the
coarse
lumpy
phase
and
the
others
are
distributed
in
the
strip-like
second
phase
and
matrix.
Fig.
6(a)
shows
the
TEM
bright-field
image
and
selected
area
electron
diffraction
(SAED)
pattern
of
the
structural
components
of
TiZn16 phase.
Fig.
6(b)
shows
the
high-resolution
TEM
image
of
AC
Zn–0.3Ti
alloy.
In
the
high-resolution
TEM
image
of
the
TiZn16
phase,
the
interplanar
spacing
(d)
of
(222)
plane
with
a
d
value
of
0.283
nm
is
consistent
with
the
interplanar
spacing
calculated
by
the
lattice
constant
of
TiZn16 (a
=
0.7720
nm,
b
=
1.1449
nm,
c
=
220

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Fig.
5.
SEM
images
and
EDS
analysis
results
of
AC
and
HR
Zn–xTi
alloys:
(a)-(d)
SEM
images
of
AC
alloys;
(e)-(h)
SEM
images
of
HR
alloys;
(i)
EDS
analysis
results
of
various
phases;
and
(j)
EDS
mapping
results
for
Zn
and
Ti
elements
in
(h).
Fig.
6.
(a)
TEM
bright-field
image
and
SAED
pattern
of
the
structural
components
of
TiZn16 phase;
(b)
high-resolution
TEM
image
of
AC
Zn–0.3Ti
alloy;
and
(c)
EDS
profile
in
(a).
1.1755
nm)
[42],
confirming
the
formation
of
ZnTi16 as
the
strips.
Fig.
6(c)
shows
the
EDS
profile
marked
in
Fig.
6(a).
It
is
a
Ti-rich
phase
that
contains
a
relative
atomic
content
of
Ti
is
6.0
±
0.4%,
which
can
be
inferred
as
the
TiZn16 phase.
3.2.
Mechanical
properties
of
Zn–xTi
Fig.
7
presents
the
tensile
properties
and
hardness
of
the
AC
and
HR
Zn–xTi
alloys.
The
tensile
stress–strain
curves
of
the
AC
and
HR
Zn
alloys
are
plotted
in
Fig.
7(a)
and
the
yield
strength
(␴ys),
ulti-
mate
tensile
strength
(␴uts),
and
elongation
(␧)
of
the
AC
and
HR
Zn
alloys
are
shown
in
Fig.
7(b)
and
Table
1.
The
HR
Zn–xTi
alloys
showed
significantly
higher
␴ys,
␴uts,
and
␧
than
the
AC
Zn–xTi
alloys.
In
addition,
the
␴ys,
␴uts,
and
␧
values
of
the
AC
and
HR
Zn–xTi
alloys
initially
increased
and
subsequently
decreased
with
the
increase
in
Ti
additions,
and
the
AC
and
HR
Zn–0.2Ti
alloys
exhibited
the
highest
strength
and
elongation
among
all
alloys.
Fig.
7(c)
and
Table
1
show
the
hardness
values
of
the
AC
and
HR
Zn–xTi
alloys.
The
hardness
of
the
HR
samples
was
slightly
lower
than
that
of
their
AC
counterparts.
It
is
worth
noting
that
the
AC
and
HR
Zn–xTi
alloys
exhibited
an
increase
in
hardness
with
the
increase
in
Ti
addition.
221

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Fig.
7.
Tensile
stress–strain
curves,
tensile
properties,
and
hardness
of
AC
and
HR
Zn–xTi
alloys:
(a)
tensile
stress–strain
curves;
(b)
yield
strength
(␴ys),
ultimate
tensile
strength
(␴uts),
and
elongation
(␧);
and
(c)
hardness.
Table
1
Mechanical
properties
and
hardness
of
AC
and
HR
Zn–xTi
alloys.
Samples
Yield
strength,
␴ys (MPa)
Ultimate
tensile
strength,
␴uts (MPa)
Elongation,
␧
(%)
Hardness
(HV)
AC
Zn–0.05Ti 62.4
±
1.4
81.8
±
1.8
0.8
±
0.2
42.4
±
1.9
AC
Zn–0.1Ti
84.1
±
1.6
93.6
±
2.7
0.9
±
0.2
47.9
±
1.3
AC
Zn–0.2Ti
105.9
±
1.7
126.9
±
2.4
1.3
±
0.3
54.8
±
2.7
AC
Zn–0.3Ti
90.4
±
2.3
113.6
±
2.6
1.2
±
0.2
59.3
±
1.7
HR
Zn–0.05Ti
129.9
±
1.9
180.1
±
3.2
39.5
±
1.9
41.6
±
2.5
HR
Zn–0.1Ti
132.4
±
3.5
200.7
±
2.7
48.5
±
3.1
47.2
±
1.2
HR
Zn–0.2Ti
144.5
±
3.1
218.7
±
3.5
54.2
±
3.0
52.1
±
1.6
HR
Zn–0.3Ti
134.6
±
3.7
217.3
±
4.6
32.9
±
3.8
58.7
±
1.4
Fig.
8
shows
the
SEM
images
of
the
tensile
fracture
surfaces
of
the
AC
and
HR
Zn–xTi
alloys.
It
can
be
seen
that
many
cleavage
planes,
tear
ridges,
and
regular
grains
appeared
on
the
fracture
surfaces
of
AC
Zn–xTi
alloys
with
no
obvious
plastic
deformation,
indicating
a
transgranular
cleavage
or
intergranular
mode
of
fracture.
Among
the
four
AC
alloys,
the
fracture
surface
of
Zn–0.2Ti
alloy
shows
clearer
plastic
deformation,
which
is
in
consistent
with
the
elon-
gation
values
of
the
AC
alloys.
In
the
HR
Zn–xTi
alloys,
the
fracture
surfaces
of
all
the
alloys
show
many
dimples
and
a
few
cleavage
facets,
showing
a
ductile
or
mixed
ductile-cleavage
mode
of
frac-
ture.
The
HR
Zn–0.2Ti
alloy
shows
significantly
denser,
larger,
and
deeper
dimples
with
almost
no
cleavage
plane
on
most
of
the
frac-
ture
surface
(Fig.
8g),
which
is
consistent
with
the
highest
tensile
elongation
among
the
four
HR
alloys,
indicating
a
ductile
mode
of
fracture.
3.3.
Corrosion
behavior
of
Zn–xTi
Fig.
9(a)
shows
potentiodynamic
polarization
curves
of
the
AC
and
HR
Zn–xTi
alloys
in
Hanks’
solution.
The
corresponding
elec-
trochemical
performance
parameters
are
listed
in
Table
2.
The
corrosion
potential
(Ecorr),
corrosion
current
density
(Icorr),
and
corrosion
rate
(Vcorr)
for
AC
Zn–xTi
alloys
were
in
the
range
of
–1.049∼–1.036
V,
19.7∼27.4
␮A/cm2,
and
293∼408
␮m/y,
respec-
tively.
After
hot-rolling,
the
HR
Zn–xTi
alloys
showed
more
negative
values
of
Ecorr and
larger
values
of
Icorr and
Vcorr than
their
AC
coun-
terparts,
thereby
exhibiting
poorer
corrosion
resistance.
The
Ecorr ,
Icorr ,
and
Vcorr values
for
the
HR
Zn–xTi
alloys
were
in
the
range
of
–1.082∼–1.054
V,
21.1∼29.0
␮A/cm2,
and
314∼432
␮m/y,
respec-
tively.
For
both
AC
and
HR
conditions
of
the
Zn–xTi
alloys,
Ecorr
decreased
and
Icorr and
Vcorr increased
with
the
increase
in
Ti
con-
tent;
hence,
the
corrosion
resistance
of
the
Zn–xTi
alloys
decreased
with
increasing
Ti
content.
Fig.
9(b)
and
Table
2
show
the
degradation
rates
of
the
AC
and
HR
Zn–xTi
alloys
after
immersion
in
Hanks’
solution
for
30
d.
The
degradation
rates
of
the
Zn–xTi
alloys
show
the
same
trend
as
the
results
obtained
from
the
electrochemical
polarization
tests,
namely,
the
degradation
rates
of
the
Zn–xTi
alloys
increased
with
increasing
addition
of
Ti,
and
the
HR
Zn–xTi
alloys
showed
a
higher
degradation
rate
than
their
AC
counterparts.
222

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Fig.
8.
SEM
images
of
fracture
surfaces
of
AC
and
HR
Zn–xTi
alloys:
(a)-(d)
SEM
images
of
AC
alloys;
(e)-(h)
SEM
images
of
HR
alloys.
Fig.
9.
Corrosion
and
degradation
behaviors
of
AC
and
HR
Zn–xTi
alloys
in
Hanks’
solution:
(a)
potentiodynamic
polarization
curves;
(b)
degradation
rates
after
30
d
immersion.
Table
2
Electrochemical
performance
parameters
of
AC
and
HR
Zn–xTi
alloys
in
Hanks’
solution.
Samples
Corrosion
potential,
Ecorr (V
vs.
SCE)
Corrosion
current
density,
Icorr
(␮A/cm2)
Polarization
resistance,
Rp
(k·cm2)
bcCorrosion
rate,
Vcorr (␮m/y)
Degradation
rate
(␮m/y)
AC
Zn–0.05Ti
–1.036
±
0.078
19.7
±
0.4
1.345
±
0.024
61.0
±
0.3
293
±
6
39.6
±
3.4
AC
Zn–0.1Ti
–1.038
±
0.094
21.6
±
0.6
1.262
±
0.021
62.8
±
0.4
322
±
9
41.5
±
2.7
AC
Zn–0.2Ti
–1.045
±
0.072
22.8
±
0.7
1.198
±
0.035
62.9
±
0.7
340
±
10
47.9
±
2.1
AC
Zn–0.3Ti
–1.049
±
0.082
27.4
±
0.3
0.983
±
0.032
62.0
±
0.3
408
±
4
51.2
±
2.9
HR
Zn–0.05Ti
–1.054
±
0.109
21.1
±
0.9
1.420
±
0.029
69.0
±
0.9
314
±
13
47.5
±
1.8
HR
Zn–0.1Ti
–1.061
±
0.094
21.8
±
0.5
1.406
±
0.031
70.6
±
0.5
325
±
7
50.6
±
2.4
HR
Zn–0.2Ti
–1.072
±
0.103
24.8
±
0.6
1.175
±
0.026
67.1
±
0.5
369
±
9
52.4
±
3.0
HR
Zn–0.3Ti
–1.082
±
0.115
29.0
±
0.5
1.015
±
0.027
67.8
±
0.4
432
±
7
57.9
±
2.4
Fig.
10
shows
the
SEM
images,
EDS
spectrum,
and
XRD
pattern
of
the
corrosion
products
on
the
AC
and
HR
Zn–xTi
alloys
surface
after
immersion
in
Hanks’
solution
for
30
d.
The
corrosion
surface
of
the
AC
and
HR
samples
were
uniformly
covered
by
corrosion
products
without
severe
localized
corrosion.
AC
and
HR
Zn–xTi
alloys
showed
an
increase
in
the
number
of
corrosion
products
on
the
sample
surface
with
increasing
addition
of
Ti,
and
the
HR
Zn–xTi
alloys
showed
more
corrosion
products
than
their
AC
counterparts,
which
is
in
consistent
with
the
degradation
rates
of
the
alloy
samples.
The
EDS
analysis
result
for
the
corrosion
products
on
HR
Zn–0.2Ti
marked
in
Fig.
10g
indicates
the
presence
of
Zn,
O,
P,
Ca,
Mg,
and
Cl,
and
the
relative
atomic
contents
of
P
and
Ca
are
9.0
±
0.3%
and
4.5
±
0.2%,
respectively
(Fig.
10i).
Fig.
10j
shows
the
XRD
pattern
of
the
corrosion
products
on
the
HR
Zn–0.2Ti
alloy
surface,
which
includes
␣-Zn,
Zn(OH)2,
ZnO,
Zn3(PO4)2,
Zn5(OH)8Cl2·H2O
and
Zn5(OH)6(CO3)2phases.
223

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Fig.
10.
SEM
images,
EDS
spectrum
and
XRD
pattern
of
corrosion
products
on
the
surface
of
AC
and
HR
Zn–xTi
alloys
after
30
d
immersion:
(a)-(d)
SEM
images
of
AC
alloys;
(e)-(h)
SEM
images
of
HR
alloys;
(i)
EDS
analysis
result
of
corrosion
products
in
(g);
and
(j)
XRD
pattern
of
corrosion
products
on
HR
Zn–0.2Ti
alloy
surface.
Fig.
11.
Friction
behaviors
of
AC
and
HR
Zn–xTi
alloys
during
wear
testing
in
Hanks’
solution:
(a)
friction
coefficient
curves;
and
(b)
friction
coefficients
and
wear
loss.
3.4.
Friction
and
wear
behavior
of
Zn–xTi
Fig.
11
shows
the
friction
behaviors
of
the
AC
and
HR
Zn–xTi
alloys
during
wear
testing
in
Hanks’
solution.
The
friction
coef-
ficient
curves
of
all
the
alloys
showed
a
steady
trend
with
prolongation
of
testing
time
and
almost
all
samples
became
rel-
atively
stable
after
500
s
of
testing
(Fig.
11a).
The
corresponding
friction
coefficients
and
wear
losses
of
the
alloys
after
wear
testing
are
shown
in
Fig.
11(b)
and
Table
3.
The
friction
coefficient
initially
decreased
and
subsequently
increased
with
increasing
addition
of
Ti,
while
the
wear
loss
decreased
with
increasing
addition
of
Ti.
It
can
be
seen
that
the
HR
Zn–xTi
alloys
showed
a
decrease
in
the
value
224

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Table
3
Friction
coefficient
and
weight
loss
of
AC
and
HR
Zn–xTi
alloys.
Samples
Friction
coefficient
Weight
loss
(mg)
AC
Zn–0.05Ti
1.027
±
0.171
2.3
±
0.4
AC
Zn–0.1Ti 0.940
±
0.196 2.2
±
0.3
AC
Zn–0.2Ti
0.861
±
0.141
2.0
±
0.2
AC
Zn–0.3Ti
0.974
±
0.168
1.9
±
0.3
HR
Zn–0.05Ti
0.943
±
0.159
2.4
±
0.2
HR
Zn–0.1Ti
0.802
±
0.162
2.3
±
0.2
HR
Zn–0.2Ti
0.776
±
0.133
2.2
±
0.3
HR
Zn–0.3Ti
0.892
±
0.164
2.0
±
0.5
of
the
friction
coefficients
and
increased
wear
losses
compared
to
their
AC
counterparts;
hence
the
AC
Zn–xTi
alloys
exhibited
higher
wear
resistance
than
the
HR
Zn–xTi
alloys.
Fig.
12
shows
the
SEM
images
of
the
AC
and
HR
Zn–xTi
sam-
ple
surface
and
a
representative
EDS
spectrum
of
the
Zn–0.05Ti
surface
after
wear
testing
in
Hanks’
solution.
It
can
be
seen
that
the
surfaces
of
the
AC
and
HR
samples
were
mainly
composed
of
clear
furrows
and
severe
plastic
deformation
in
most
areas
of
the
surfaces
formed
by
abrasive
wear,
and
small
pits
and
flakes
or
band-shaped
wear
debris
formed
by
adhesive
wear,
indicat-
ing
a
mixed
wear
mechanism
of
typical
abrasive
wear,
adhesive
wear
and
peeling
wear
mechanism
[43].
In
addition,
there
were
black
granular
or
irregular
wear
debris
at
the
edge
of
the
wear
grooves.
HR
samples
showed
wider
grooves
than
those
of
the
AC
samples,
indicating
high
wear
loss
among
the
AC
and
HR
Zn–xTi
alloys.
The
EDS
spectrum
of
the
wear
product
on
the
wear
surface
of
HR
Zn–0.05Ti
alloy
marked
in
Fig.
12(e)
shows
the
pres-
ence
of
Zn,
O,
Cl,
P,
and
Ca.
The
wear
products
may
contain
the
compounds
of
ZnO
or
Zn(OH)2and
ZnCl2or
Zn5(OH)8Cl2·H2O
formed
in
the
physiological
environment
of
Hanks’
solution
during
wear
test.
3.5.
Cytotoxicity
of
Zn–xTi
Fig.
13
shows
the
cell
viability
of
MG-63
cells
after
culturing
with
100%,
25%,
and
12.5%
extracts
of
HR
Zn–xTi
alloys
for
1
d.
The
cell
viability
of
MG-63
cells
was
clearly
affected
by
the
concentration
of
the
extract
and
increased
with
the
dilution
of
the
extract.
The
cell
viability
of
MG-63
cells
was
43.13
±
1.44%
for
Zn–0.05Ti,
38.05
±
0.27%
for
Zn–0.1Ti,
36.96
±
1.38%
for
Zn–0.2Ti,
and
33.29
±
0.50%
for
Zn–0.3Ti
with
the
100%
extracts,
indicating
grade
3
cytotoxic-
ity.
It
is
worth
noting
that
the
100%
extracts
of
HR
Zn–xTi
alloys
show
decreased
cell
viability
with
increasing
Ti
addition.
Accord-
ing
to
ISO
10993−5
[40],
cell
viability
of
80%–100%
is
considered
slightly
cytotoxic
(grade
1
cytotoxicity),
between
50%
and
80%
is
mild
(grade
2
cytotoxicity),
30%–50%
is
moderate
(grade
3
cytotox-
icity),
while
lower
than
30
%
is
severe
cytotoxicity
(grade
4).
When
the
concentration
of
the
extracts
was
diluted
to
25%
and
12.5%,
the
cytotoxicity
of
the
HR
Zn–xTi
alloys
significantly
decreased.
The
cell
viability
of
the
25%
extracts
was
78.61
±
1.61%
for
Zn–0.05Ti,
80.22
±
1.41%
for
Zn–0.1Ti,
81.32
±
1.28%
for
Zn–0.2Ti,
and
84.77
±
0.43%
for
Zn–0.3Ti,
all
of
which
approximate
or
exceed
80%,
indicating
grade
1
cytotoxicity,
except
the
Zn–0.05Ti
alloy,
which
showed
a
cell
viability
approximating
80%.
The
cell
viability
of
the
12.5%
extracts
was
93.29
±
1.86%
for
Zn–0.05Ti,
92.90
±
1.84%
for
Zn–0.1Ti,
95.77
±
0.99%
for
Zn–0.2Ti,
and
97.57
±
1.46%
for
Zn–0.3Ti,
indicating
satisfactory
cytocompatibility.
It
is
worth
not-
ing
that
the
25%
and
12.5%
extracts
of
the
HR
Zn–xTi
alloys
showed
enhanced
cell
viability
with
the
addition
of
Ti
in
concentrations
of
25%
and
12.5%,
except
for
the
Zn–0.1Ti
alloy
at
the
concentration
of
12.5%.
Fig.
12.
SEM
images
and
EDS
spectrum
of
AC
and
HR
Zn–xTi
alloys
after
wear
testing
in
Hanks’
solution:
(a)-(d)
SEM
images
of
AC
alloys;
(e)-(h)
SEM
images
of
HR
alloys;
and
(i)
representative
EDS
spectrum
of
wear
surface
in
(e).
225

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Fig.
13.
Cell
viability
of
MG-63
cells
after
culturing
with
100%,
25%
and
12.5%
extracts
of
HR
Zn–xTi
alloys
for
1
d.
4.
Discussion
4.1.
Microstructures
and
mechanical
properties
of
Zn–xTi
According
to
the
Zn–Ti
phase
diagram
[41],
the
solid
solubility
of
Ti
in
the
Zn
matrix
is
approximately
0.007∼0.015
wt.%
at
300
◦C.
The
melt
undergoes
a
eutectic
reaction
of
L
→
␣-Zn
+
TiZn16
when
the
melting
temperature
reaches
418.6 ◦C,
and
excessive
Ti
will
form
fine
particles
of
TiZn16 with
the
Zn
and
this
phase
will
be
dispersed
on
the
grain
boundaries
of
the
␣-Zn
phase
[34].
During
the
crystallization
process,
the
Ti
mainly
forms
crystal
nuclei
of
the
TiZn16 phase
and
serves
as
a
crystalline
core
to
promote
non-
uniform
nucleation
in
melts,
thereby
refining
the
grain
size
of
Zn
alloys
[44].
This
is
confirmed
by
the
experimental
results
that
the
average
grain
size
of
the
␣-Zn
phase
was
19.4
±
1.9
␮m
for
AC
Zn–0.2Ti,
71.2
±
12.3
␮m
for
AC
Zn–0.1Ti,
and
86.4
±
10.9
␮m
for
AC
Zn–0.05Ti.
However,
a
further
increase
in
addition
of
Ti
to
AC
Zn–0.3Ti
caused
a
coarsening
of
the
␣-Zn
phase
(46.9
±
8.2
␮m).
This
was
mainly
caused
by
the
reaction
between
the
Zn
and
the
excessive
Ti,
which
generated
TiZn16 compounds,
accompanied
by
the
release
of
a
large
amount
of
latent
heat
of
crystallization
during
solidification
of
the
Zn–Ti
alloy
[45].
This
heat
reduced
the
degree
of
supercooling
of
the
Zn
melt
and
this
reduced
the
nucleation
rate,
leading
to
a
coarsening
of
the
␣-Zn
phase.
Excessive
addition
of
Ti
also
led
to
a
rapid
growth
in
the
TiZn16 phase,
as
observed
in
Fig.
2(d).
Spittle
et
al.
[46]
reported
that
Ti
and
Zn
react
and
form
an
irregular
TiZn16 phase
with
high
Ti
content
and
a
slow
cooling
rate.
The
ultimate
tensile
strength
(␴uts)
and
yield
strength
(␴ys)
were
33.6
MPa
and
29.3
MPa
for
AC
pure
Zn,
and
153.1
MPa
and
84.2
MPa
for
HR
pure
Zn
[30].
With
the
addition
of
Ti,
the
AC
Zn–xTi
alloys
showed
higher
strength
than
the
AC
pure
Zn
(Table
2).
This
was
mainly
due
to
the
strengthening
effects
of
grain
refinement
and
second-phase
precipitation
strengthening.
The
addition
of
Ti
contributed
to
the
refinement
of
the
␣-Zn
phase
and
the
precipita-
tion
of
the
second
intermetallic
phase
of
TiZn16.
According
to
the
Hall–Petch
relationship,
the
␴ys of
an
alloy
is
inversely
proportional
to
the
square
root
of
its
grain
size
[47,48],
while
the
intermetallic
TiZn16 phase
contributes
to
precipitation
strengthening
[49].
Furthermore,
with
increasing
additions
of
Ti,
the
strength
and
elongation
of
the
Zn–xTi
alloys
initially
increased
and
subsequently
decreased,
and
the
HR
Zn–0.2Ti
alloy
showed
the
highest
␴uts
(218.7
±
3.5
MPa),
␴ys (144.5
±
3.1
MPa),
and
␧
(54.2
±
3.0%)
among
all
the
Zn–xTi
alloys
(Table
2).
This
was
mainly
because
exces-
sive
addition
of
Ti
resulted
in
coarsening
of
the
␣-Zn
phase
and
growth
of
the
second
phase
of
TiZn16 into
strips
distributed
along
the
grain
boundaries.
Cheng
et
al.
[50]
reported
that
fine
and
uni-
form
second-phase
particles
distributed
along
the
grain
boundaries
can
effectively
pin
the
grain
boundaries
and
thus
refine
the
grain
size.
However,
such
a
pinning
effect
will
decrease
with
the
growth
of
the
second-phase
particles.
Coarsening
of
the
␣-Zn
phase
and
growth
of
the
second-phase
particles
will
cause
a
decrease
in
both
the
strength
and
elongation
of
the
Zn
alloy.
In
addition,
due
to
unco-
ordinated
deformation
between
the
brittle
intermetallic
phase
and
the
matrix
in
the
process
of
plastic
deformation,
stress
concentra-
tions
and
voids
may
form
more
easily
at
the
interface
between
the
second
phase
and
the
matrix,
resulting
in
premature
fracture
of
the
alloy
and
thus
reduced
plastic
deformation
ability
[51,52].
There-
fore,
the
Zn–0.3Ti
alloy
showed
lower
␴ys, ␴uts, and
␧
as
compared
to
the
Zn–0.2Ti
alloy.
It
is
worth
noting
that
the
mechanical
properties
of
the
HR
Zn–xTi
alloys
were
significantly
better
than
those
of
their
AC
coun-
terparts.
This
was
mainly
because
the
second-phase
particles
were
more
homogenously
distributed
in
the
matrix
after
hot-rolling,
which
enhanced
the
effect
of
the
second
phase
on
preventing
grain-
boundary
sliding,
leading
to
an
increase
in
deformation
resistance,
which
improved
the
mechanical
properties
[49].
In
the
Zn–xTi
alloys,
the
solid
solubility
of
Ti
in
Zn
is
very
low,
so
the
amount
of
Ti
dissolved
in
the
Zn
matrix
is
minimal
[53].
Most
of
the
Ti
is
dispersed
along
the
grain
boundary
in
the
form
of
the
TiZn16 sec-
ond
phase,
which
is
refined
after
hot-rolling
and
plays
a
role
in
precipitation
strengthening.
The
HR
Zn–xTi
alloys
showed
a
slight
decrease
in
hardness
compared
to
their
AC
counterparts,
which
resulted
from
the
combined
effect
of
preheating
and
the
recov-
ery
and
recrystallization
caused
by
the
heat
of
deformation
during
hot-rolling,
leading
to
softening
of
the
HR
alloys
[22,54].
Gu
et
al.
[55]
and
Yanson
et
al.
[56]
reported
that
the
femur
cor-
tical
bone
shows
a
tensile
strength
of
35∼383
MPa,
yield
strength
of
104.9∼114.3
MPa,
hardness
of
16∼50.2
HV,
and
elongation
of
5∼23%,
while
cancellous
bone
has
a
tensile
strength
of
1.5∼38
MPa
and
elongation
of
0.01∼1.57%.
Compared
with
the
mechan-
ical
properties
of
natural
bone,
the
tensile
strength
and
hardness
of
the
HR
Zn–xTi
alloys
are
higher
than
those
of
cortical
bone.
There-
fore,
the
mechanical
properties
of
HR
Zn–xTi
alloys
can
satisfy
the
requirements
of
the
mechanical
properties
for
orthopedic
implants.
4.2.
Corrosion
behavior
of
Zn–xTi
The
standard
electrode
potentials
of
the
alloying
elements
are
0.763
V
for
Zn
and
−1.630
V
for
Ti
[57].
Therefore,
the
addition
of
Ti
into
our
Zn–xTi
alloys
played
the
leading
roles
during
corrosion
that
the
undissolved
Ti
reacted
with
Zn
and
formed
the
second
phase
of
TiZn16 at
the
grain
boundary.
Fig.
14
shows
a
schematic
illustration
of
the
corrosion
mechanism
of
Zn–xTi
alloys
immersed
in
Hanks’
solution.
In
the
physiological
environment
of
Hanks’
solution,
the
Zn–xTi
alloys
underwent
oxygen
reduction
corrosion
during
the
degradation
process.
There
is
an
interface
between
the
second
phase
and
␣-Zn
matrix,
and
micro-battery
corrosion
occurred
in
Hanks’
solution
due
to
the
potential
difference
between
the
two
phases,
thereby
accelerating
the
corrosion
process
of
the
Zn
alloy.
The
second
phase,
with
a
lower
self-corrosion
potential,
usually
serves
as
an
anode
and
is
usually
preferentially
corroded,
and
the
␣-Zn
matrix
serves
as
the
cathode.
With
increasing
Ti
addition,
the
size
and
content
of
the
second-phase
particles
increased
sig-
nificantly,
resulting
in
an
increase
in
the
active
anode
area,
which
caused
more
severe
micro-battery
corrosion
and
a
decrease
in
the
corrosion
resistance
of
the
Zn–xTi
alloys.
Zhang
et
al.
[58]
reported
that
the
corrosion
rates
of
Zn–2Cu–xTi
(x
=
0,
0.05,
and
0.1
wt.%)
alloys
substantially
increased
with
an
increase
in
Ti
content.
After
the
hot-rolling
treatment,
the
HR
Zn–xTi
alloys
showed
a
decrease
in
corrosion
resistance
compared
to
their
AC
counterparts
226

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
Fig.
14.
Schematic
illustration
of
corrosion
mechanism
of
Zn–xTi
alloys
immersed
in
Hanks’
solution.
due
to
the
increased
likelihood
of
crack
generation
during
hot-
rolling.
The
hard
and
brittle
intermetallic
TiZn16 phase
was
prone
to
micro-cracking
during
the
rolling
process
(Fig.
3d)
and
the
cracks
in
the
structure
caused
an
increase
in
the
contact
area
between
the
Hanks’
solution
and
the
substrate
along
the
cracks,
thereby
accel-
erating
the
corrosion
[59].
In
addition,
the
rolling
process
may
have
weakened
the
corrosion
barrier
effect
of
the
second-phase
network
that
could
inhibit
the
localized
corrosion
of
the
Zn
alloy,
thereby
forming
shear
bands
that
introduced
structural
defects
and
residual
stress
into
the
materials,
thereby
reducing
the
corrosion
resistance
[30].
4.3.
Friction
and
wear
behavior
of
Zn–xTi
There
is
wear
between
the
implant
material
and
the
surrounding
tissue
after
plastic
surgery
or
bone
implantation.
Wear
can
cause
debris
at
the
interface
between
the
bone
tissue
and
the
implant,
and
this
can
trigger
a
series
of
adverse
reactions
such
as
inflamma-
tion
and
bone
resorption
[60].
This
is
mainly
because
macrophages
are
activated
after
phagocytosis
of
debris
and
produce
a
series
of
conditioning
media
that
lead
to
the
generation
of
mature
osteo-
clasts
[61,62].
Inflammation
and
bone
resorption
around
implant
materials
are
the
main
causes
of
failure
of
implant
surgery
[63].
Therefore,
the
implant
material
must
have
suitable
wear
resis-
tance
to
reduce
the
formation
of
wear
debris,
thereby
preventing
inflammation
and
bone
resorption.
According
to
the
Holm–Archard
equation
[64],
the
wear
resistance
of
a
material
is
dependent
on
its
hardness,
and
a
material
with
higher
hardness
shows
lower
wear
loss.
In
this
study,
the
wear
loss
of
both
the
AC
and
HR
Zn–xTi
alloys
was
inversely
proportional
to
their
hardness
and
was
lower
than
that
of
pure
Zn,
indicating
good
wear
performance
[39].
In
addition,
the
coarser
␣-Zn
matrix
phase
and
TiZn16 second
phase
induced
a
stronger
obstruction
effect
on
the
friction
pair,
thereby
resulting
in
a
greater
friction
coefficient
which
was
also
less
stable.
As
shown
in
Fig.
2
and
Fig.
3,
both
the
AC
and
HR
Zn–0.3Ti
alloys
exhibited
the
coarsest
second
phases,
thereby
showing
the
highest
coefficients
of
friction
in
Hanks’
solution.
4.4.
Cytocompatibility
of
Zn–xTi
Biodegradable
orthopedic
implant
materials
should
have
the
following
material
properties
[65,66]:
(1)
satisfactory
biocompati-
bility;
(2)
no
carcinogenicity,
teratogenicity,
or
toxic
side
effects;
(3)
sufficient
mechanical
properties;
(4)
gradual
degradation
and
even-
tually
absorption
by
the
tissue
after
the
completion
of
the
function
in
the
body;
and
(5)
easy
processing
and
sterilization.
Therefore,
Zn-
based
biodegradable
metals
must
not
only
satisfy
the
mechanical
and
corrosion
resistance
requirements,
but
also
have
satisfactory
biocompatibility.
Ti
has
good
biocompatibility
as
a
beneficial
ele-
ment
in
the
human
body
[67,68].
In
addition,
Ti
can
stimulate
phagocytic
cells
and
significantly
enhance
immune
responses
[69].
The
recommended
intake
and
content
levels
in
healthy
adults
are
approximately
12∼16
mg/d
and
2∼4
g
for
Zn
ions,
and
0.3∼2
mg/d
and
15
mg
for
Ti
ions,
respectively;
but
excessive
intake
of
these
metal
ions
will
cause
systemic
toxicity
[35,70–72].
Ti
and
some
of
its
alloys
are
widely
used
in
bone
fixation
and
replacement
due
to
their
stable
chemical
properties,
satisfactory
corrosion
resistance,
and
good
mechanical
properties
[73].
However,
Termine
et
al.
[74]
reported
that
the
Ti
ions
that
are
released
by
Ti
implants
inhibit
the
expression
of
the
osteoblast
phenotype
and
the
deposition
of
mineralized
matrix,
and
significantly
inhibit
the
synthesis
of
osteo-
calcin,
which
may
lead
to
bone
resorption
around
the
implant.
Blumenthal
et
al.
[75]
reported
on
the
effects
of
Ti4+ and
V5+ ions
on
hydroxyapatite
(HA)
formation
and
found
that
the
release
of
Ti
ions
into
the
implant–tissue
interface
could
interfere
with
the
normal
osteoid
mineralization
and
remodeling
processes.
Liao
et
al.
[76]
reported
that
a
concentration
of
Ti
ions
greater
than
or
equal
to
10
ppm
inhibited
the
cell
proliferation
of
rat
calvaria
cells,
while
a
concentration
of
Ti
ions
less
than
or
equal
to
5
ppm
had
no
effect
or
stimulated
proliferation.
Therefore,
it
is
inferred
that
low
Ti
ion
concentrations
have
minimal
effect
on
mature
osteoblasts,
but
high
Ti
ion
concentrations
will
accelerate
the
bone
resorption
and
weaken
the
bone-formation
ability
of
osteoblasts.
Most
excess
Ti
ions
are
excreted
through
feces
and
urine
[77].
Woodman
et
al.
[78]
reported
that
the
level
of
Ti
ions
in
the
urine
of
baboons
released
by
Ti-based
segmental
bone
replacements
was
sixfold
higher
than
that
of
controls
without
implants
at
2
weeks
after
surgery,
while
this
level
subsequently
decreased
and
approached
the
preopera-
tive
level;
hence,
most
Ti
ions
can
be
excreted
through
the
urine.
In
this
study,
based
on
the
degradation
rates
(Fig.
9b),
the
highest
concentration
of
Ti
ions
was
0.68
␮g/d
released
by
an
HR
Zn–0.3Ti
screw
with
dimensions
of
2.5
mm
in
diameter
and
18
mm
in
length
and
a
surface
area
of
approximately
2.0
cm2[79],
far
below
the
recommended
daily
intake.
In
the
undiluted
extracts
of
the
HR
Zn–xTi
alloys,
MG-63
cells
showed
grade
3
cytotoxicity;
this
was
because
the
concentration
of
metal
ions
in
the
undiluted
extracts
exceeded
the
ion
concentration
tolerance
limit
of
MG-63
cells,
thereby
inhibiting
cell
prolifer-
ation
and
growth,
indicating
moderate
cytotoxicity.
In
addition,
the
cytotoxicity
of
the
HR
Zn–xTi
alloys
gradually
increased
with
increasing
Ti
addition.
This
may
have
been
due
to
the
increased
metal
ion
concentrations
of
Zn
and
Ti
in
the
extracts,
which
may
have
resulted
from
two
effects:
first,
the
higher
Ti-content
Zn
alloy
released
more
Ti
ions
during
degradation;
and,
second,
the
degra-
dation
rate
increased
with
increasing
Ti
addition
(Fig.
9),
resulting
in
an
increase
in
both
Ti
and
Zn
ion
concentrations.
Jin
et
al.
[80]
reported
that
an
appropriate
amount
of
released
Zn2+ ions
from
titanium
with
dual
Zn/Ag
ion
implantation
significantly
improved
the
osteogenic
activity
of
rat
bone
marrow
mesenchymal
stem
cells
(rBMSCs)
via
the
long-range
interactions.
As
the
concentration
of
the
extracts
was
diluted
to
25%,
the
cell
viability
of
the
HR
Zn–xTi
alloys
were
approximately
or
higher
than
80%,
and
these
alloys
can
be
considered
non-toxic.
Tang
et
al.
[23]
also
reported
that
cell
viability
gradually
increased
with
dilution
227

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
of
the
extracts.
When
the
concentration
of
the
alloy
extract
was
diluted
to
12.5%,
the
MG-63
cells
showed
further
enhanced
cell
viability,
all
over
90%.
5.
Conclusion
In
this
study,
we
comprehensively
investigated
the
microstruc-
ture,
mechanical
properties,
corrosion
resistances
and
degradation
rates,
friction
and
wear
performance,
and
cytotoxicity
of
binary
Zn–xTi
(x
=
0.05,
0.1,
0.2,
and
0.3%)
alloys
under
the
conditions
of
as-casting
and
hot-rolling
for
orthopedic
implant
and
stent
appli-
cations.
The
main
conclusions
are
as
follows:
1
The
microstructures
of
the
AC
and
HR
Zn–xTi
alloys
were
mainly
composed
of
an
␣-Zn
matrix
phase
and
an
intermetallic
TiZn16
phase.
The
grain
size
of
the
␣-Zn
matrix
decreased
with
increas-
ing
Ti
addition
up
to
0.2%,
then
increased
at
0.3%
Ti
addition.
Both
the
size
and
content
of
the
TiZn16 phase
gradually
increased
with
increasing
Ti
addition.
After
hot-rolling,
both
the
␣-Zn
phase
and
TiZn16 phase
were
elongated
and
distributed
along
the
rolling
direction.
2
The
HR
Zn–0.2Ti
alloy
showed
the
highest
strength
and
elonga-
tion
among
all
the
Zn–xTi
alloys,
with
␴ys =
144.5
MPa,
␴uts =
218.7
MPa,
and
␧
=
54.2%.
The
hardness
of
the
AC
and
HR
Zn–xTi
alloys
increased
with
an
increase
in
Ti
addition.
3
The
AC
and
HR
Zn–xTi
alloys
exhibited
decreased
corrosion
potential,
and
increased
corrosion
current
density
and
corrosion
rate,
with
increasing
Ti
addition.
The
HR
Zn–xTi
alloys
showed
higher
corrosion
current
densities
and
corrosion
rates
than
their
AC
counterparts.
The
HR
Zn–xTi
alloys
showed
a
corrosion
cur-
rent
density
in
the
range
of
21.1∼29.0
␮A/cm2and
a
corrosion
rate
in
the
range
of
314∼432
␮m/y.
4
The
AC
and
HR
Zn–xTi
alloys
exhibited
increased
degradation
rates
with
increasing
Ti
addition,
and
the
HR
Zn–xTi
alloys
showed
higher
degradation
rates
than
their
AC
counterparts.
The
HR
Zn–xTi
alloy
showed
a
degradation
rate
ranging
from
47.5
to
57.9
␮m/y.
5
Wear
tests
in
Hanks’
solution
indicated
that
the
friction
coeffi-
cients
of
the
AC
and
HR
Zn–xTi
alloys
decreased
with
increasing
Ti
addition
up
to
0.2%Ti,
then
increased
at
0.3%Ti.
The
wear
loss
decreased
with
increasing
addition
of
Ti
and
the
AC
Zn–xTi
alloys
showed
lower
wear
loss
than
their
HR
counterparts.
6
The
extracts
of
both
AC
and
HR
Zn–xTi
alloys
at
concentrations
of
≤
25%
showed
no
cytotoxicity
to
MG-63
cells.
Acknowledgments
This
work
was
supported
by
the
Wenzhou
Science
and
Technol-
ogy
Bureau
through
the
project
ZG2019022
and
2018ZG008.
CW
and
YL
also
acknowledge
the
financial
support
for
this
research
by
the
Australian
Research
Council
(ARC)
through
the
Discovery
Project
DP170102557
and
Future
FellowshipFT160100252.
MD
and
CW
also
acknowledge
the
support
of
the
ARC
Research
Hub
for
Advanced
Manufacturing
of
Medical
Devices
(IH150100024).
References
[1]
Y.F.
Zheng,
X.N.
Gu,
F.
Witte,
Mater.
Sci.
Eng.,
R
77
(2014)
1–34.
[2]
L.
Tan,
X.
Yu,
P.
Wan,
K.
Yang,
J.
Mater.
Sci.
Technol.
29
(6)
(2013)
503–513.
[3]
M.
Heiden,
E.
Walker,
L.
Stanciu,J.
Biotechnol.
Biomater.
5
(2)
(2015),
1000178.
[4]
M.
Haude,
H.
Ince,
A.
Abizaid,
R.
Toelg,
P.A.
Lemos,
C.
von
Birgelen,
E.
Eeckhout,
Lancet
387
(10013)
(2016)
31–39.
[5]
M.
Ezechieli,
M.
Ettinger,
C.
König,
A.
Weizbauer,
P.
Helmecke,
R.
Schavan,
C.
Becher,
Knee
Surg,
Sport
Tr.
A
24
(12)
(2016)
3976–3981.
[6]
B.
Liu,
Y.F.
Zheng,
Acta
Biomater.
7
(3)
(2011)
1407–1420.
[7]
M.
Schinhammer,
A.C.
Hänzi,
J.F.
Löffler,
P.J.
Uggowitzer,
Acta
Biomater.
6
(5)
(2010)
1705–1713.
[8]
C.
Wu,
X.
Hu,
H.
Qiu,
Y.
Ruan,
Y.
Tang,
A.
Wu,
Q.
Wang,
J.
Am.
Coll.
Cardiol.
60
(17)
(2012)
B166.
[9]
W.J.
Lin,
D.Y.
Zhang,
G.
Zhang,
H.T.
Sun,
H.P.
Qi,
L.P.
Chen,
W.
Zheng,
Mater.
Des.
91
(2016)
72–79.
[10]
R.R.
Watson,
V.R.
Preedy,
S.
Zibadi,
Magnesium
in
Human
Health
and
Disease,
Springer,
2012.
[11]
A.
Sigel,
H.
Sigel,
R.K.O.
Sigel,
Interrelations
Between
Essential
Metal
Ions
and
Human
Diseases,
Springer,
Netherlands,
2013.
[12]
R.
Crichton,
R.R.
Crichton,
J.R.
Boelaert,
Inorganic
Biochemistry
of
Iron
Metabolism:
From
Molecular
Mechanisms
to
Clinical
Consequences,
John
Wiley
&
Sons,
2001.
[13]
X.N.
Gu,
W.R.
Zhou,
Y.F.
Zheng,
Y.
Cheng,
S.C.
Wei,
S.P.
Zhong,
L.J.
Chen,
Acta
Biomater.
6
(12)
(2010)
4605–4613.
[14]
J.
Kuhlmann,
I.
Bartsch,
E.
Willbold,
S.
Schuchardt,
O.
Holz,
N.
Hort,
D.
Höche,
W.R.
Heineman,
F.
Witte,
Acta
Biomater.
9
(10)
(2013)
8714–8721.
[15]
M.
Schinhammer,
A.C.
Hänzi,
J.F.
Löffler,
P.J.
Uggowitzer,
Acta
Biomater.
6
(5)
(2010)
1705–1713.
[16]
P.K.
Bowen,
J.
Drelich,
J.
Goldman,
Adv.
Mater.
25
(18)
(2013)
2577–2582.
[17]
D.
Vojtech,
J.
Kubasek,
J.
Serak,
P.
Novak,
Acta
Biomater.
7
(2011)
3515–3522.
[18]
E.
Mocchegiani,
M.
Muzzioli,
R.
Giacconi,
Trends
Pharmacol.
Sci.
21
(2000)
205–208.
[19]
H.
Haase,
L.
Rink,
Biofactors
40
(1)
(2014)
27–40.
[20]
H.
Tapiero,
K.D.
Tew,
Biomed.
Pharmacother.
57
(9)
(2003)
399–411.
[21]
J.Z.
Ilich,
J.E.
Kerstetter,
J.
Am.
Coll.
Nutr.
19
(2000)
715–737.
[22]
J.C.T.
Farge,
W.M.
Williams,
Can.
Metall.
Q.
5
(4)
(1966)
265–272.
[23]
Z.
Tang,
H.
Huang,
J.
Niu,
L.
Zhang,
H.
Zhang,
J.
Pei,
G.
Yuan,
Mater.
Des.
117
(2017)
84–94.
[24]
H.F.
Li,
X.H.
Xie,
Y.F.
Zheng,
Y.
Cong,
F.Y.
Zhou,
K.J.
Qiu,
X.
Wang,
S.H.
Chen,
L.
Huang,
L.
Tian,
L.
Qin,
Sci.
Rep.
5
(2015)
10719.
[25]
Z.B.
Tang,
J.L.
Niu,
H.
Huang,
H.
Zhang,
J.
Pei,
J.M.
Ou,
G.Y.
Yuan,
J.
Mech.
Behav.
Biomed.
Mater.
72
(2017)
182–191.
[26]
S.
Sun,
Y.
Ren,
L.
Wang,
B.
Yang,
H.
Li,
G.
Qin,
Mater.
Sci.
Eng.,
A
701
(2017)
129–133.
[27]
P.
Li,
C.
Schille,
E.
Schweizer,
F.
Rupp,
A.
Heiss,
C.
Legner,
L.
Scheideler,
Int.
J.
Mol.
Sci.
19
(3)
(2018)
755.
[28]
A.
Kafri,
S.
Ovadia,
G.
Yosafovich-Doitch,
E.
Aghion,
J.
Mater.
Sci.
Mater.
Med.
29
(7)
(2018)
94.
[29]
S.
Zhao,
J.M.
Seitz,
R.
Eifler,
H.J.
Maier,
R.J.
Guillory
II,
E.J.
Earley,
J.W.
Drelich,
Mater.
Sci.
Eng.,
C
76
(2017)
301–312.
[30]
X.
Tong,
D.C.
Zhang,
X.T.
Zhang,
Y.C.
Su,
Z.M.
Shi,
K.
Wang,
J.G.
Lin,
Y.C.
Li,
J.X.
Lin,
C.E.
Wen,
Acta
Biomater.
82
(2018)
197–204.
[31]
M.
Dambatta,
S.
Izman,
D.
Kurniawan,
H.
Hermawan,
J.
King
Saud
Univ.,
Sci.
29
(4)
(2017)
455–461.
[32]
N.G.
Durmus,
T.J.
Webster,
Nanomedicine
7
(6)
(2012)
791–793.
[33]
X.
Ai,
G.
Quan,
Open
Mater.
Sci.
J.
6
(1)
(2012)
6–13.
[34]
Z.Y.
Yin,
Microstructural
Evolution
and
Mechanical
Properties
of
Zn–Ti
Alloys
for
Biodegradable
Stent
Applications,
M.S.
Thesis,
Michigan
Technological
University,
2017.
[35]
Y.
Li,
C.
Wong,
J.
Xiong,
P.
Hodgson,
C.
Wen,
J.
Dent.
Res.
89
(5)
(2010)
493–497.
[36]
ASTM
E8/E
8M–16,
Standard
test
methods
for
tension
testing
of
metallic
materials,
ASTM
International,
West
Conshohocken,
PA
(2011).
[37]
L.J.
Liu,
Y.
Meng,
A.A.
Volinsky,
H.J.
Zhang,
L.N.
Wang,
Corros.
Sci.
153
(2019)
341–356.
[38]
ASTM
G31–72,
Standard
practice
for
laboratory
immersion
corrosion
testing
of
metals,
ASTM
International,
West
Conshohocken,
PA
(2004).
[39]
J.X.
Lin,
X.
Tong,
Z.M.
Shi,
D.C.
Zhang,
L.S.
Zhang,
K.
Wang,
A.P.
Wei,
L.F.
Jin,
J.G.
Lin,
Y.C.
Li,
C.E.
Wen,
Acta
Biomater.
106
(2020)
410–427.
[40]
ISO
10993–5,
Biological
evaluation
of
medical
devices.
Part
5:
tests
for
in
vitro
cytotoxicity,
International
Organisation
for
Standardization,
Geneva,
Switzerland
(2009)
2009.
[41]
G.
Boczkal,
Mod.
Aspects
Bulk
Cryst.
Thin
Film
Prep.
(2012)
141–162.
[42]
M.
Saillard,
G.
Develey,
C.
Becle,
J.M.
Moreau,
D.
Paccard,
Acta
Crystallogr.,
Sect.
B:
Struct.
Crystallogr.
Cryst.
Chem.
37
(1)
(1981)
224–226.
[43]
P.
BK,
S.
Das,
Mater.
Trans.,
JIM
36
(8)
(1995)
1048–1057.
[44]
G.
Boczkal,
Arch.
Metall.
Mater.
58
(4)
(2013)
1019–1022.
[45]
A.
Safari,
R.
Saidur,
F.A.
Sulaiman,
Y.
Xu,
J.
Dong,
Renew.
Sust.
Energ.
Rev.
70
(2017)
905–919.
[46]
J.A.
Spittle,
Metallography
5
(5)
(1972)
423–447.
[47]
E.O.
Hall,
Proc.
Phys.
Soc.
B
64
(1951)
747–753.
[48]
N.J.
Petch,
J.
Iron
Steel
Inst.
174
(1953)
25–28.
[49]
W.D.
Callister,
D.G.
Rethwisch,
Materials
Science
and
Engineering:
an
Introduction,
eight
ed.,
John
Wiley
and
Sons,
Inc,
Hoboken,
NJ,
2014.
[50]
W.L.
Cheng,
S.S.
Park,
B.S.
You,
B.H.
Koo,
Mater.
Sci.
Eng.,
A
527
(18–19)
(2010)
4650–4653.
[51]
W.
Zhang,
M.
Wang,
W.
Chen,
Y.
Feng,
Y.
Yu,
J.
Alloys.
Compd.
669
(2016)
79–90.
[52]
W.
Liu,
J.
Zhang,
L.
Wei,
C.
Xu,
X.
Zong,
J.
Hao,
Mater.
Sci.
Eng.,
A
681
(2017)
97–102.
[53]
J.
Venezuela,
M.S.
Dargusch,
Acta
Biomater.
87
(2019)
1–40.
[54]
W.
Wang,
A.L.
Helbert,
T.
Baudin,
F.
Brisset,
R.
Penelle,
Mater.
Charact.
64
(2012)
1–7.
[55]
X.N.
Gu,
Y.F.
Zheng,
Front.
Mater.
Sci.
China
4
(2)
(2010)
111–115.
[56]
K.A.
Yanson,
G.R.
Bite,
I.V.
Knets,
Y.Z.
Saulgozis,
Polym.
Mech.
9
(6)
(1973)
966–971.
[57]
P.
Van ´
ysek,
Electrochemical
Series,
CRC
Press
LLC.,
2008.
[58]
L.
Zhang,
X.Y.
Liu,
H.
Huang,
W.
Zhan,
Mater.
Lett.
244
(2019)
119–122.
228

K.
Wang
et
al.
Journal
of
Materials
Science
&
Technology
74
(2021)
216–229
[59]
S.E.
Harandi,
P.C.
Banerjee,
C.D.
Easton,
R.S.
Raman,
Mater.
Sci.
Eng.,
C
80
(2017)
335–345.
[60]
T.W.
Bauer,
Clin.
Orthop.
Relat.
Res.
405
(2002)
138–143.
[61]
M.R.
Parsek,
E.P.
Greenberg,
Trends
Microbiol.
13
(1)
(2005)
27–33.
[62]
W.D.
Hamilton,
Nature
228
(5277)
(1970)
1218–1220.
[63]
W.H.
Harris,
Acta
Orthop.
Scand.
65
(1)
(1994)
113–123.
[64]
S.
Kumar,
M.
Chakraborty,
V.
Subramanya
Sarma,
B.S.
Murty,
Wear
265
(1–2)
(2008)
134–142.
[65]
S.C.
Lee,
M.
Shea,
M.A.
Battle,
K.
Kozitza,
E.
Ron,
T.
Turek,
W.C.
Hayes,
J.
Biomed,
Mater.
Res.
28
(10)
(1994)
1149–1156.
[66]
R.
Kenley,
L.
Marden,
T.
Turek,
L.
Jin,
E.
Ron,
J.O.
Hollinger,
J.
Biomed.
Mater.
Res.
28
(10)
(1994)
1139–1147.
[67]
T.
Albrektsson,
P.I.
Branemark,
H.A.
Hansson,
J.
Lindström,
Acta
Orthop.
Scand.
52
(2)
(1981)
155–170.
[68]
S.
Asadpour,
M.
Chamsaz,
M.H.
Entezari,
M.J.
Haron,
N.
Ghows,
Arabian
J.
Chem.
9
(2016)
S1833–S1839.
[69]
Y.
Wang,
C.
Yao,
L.
Ding,
C.
Li,
J.
Wang,
M.
Wu,
Y.
Lei,
J.
Biomed.
Nanotechnol.
13
(4)
(2017)
367–380.
[70]
D.
Bian,
W.R.
Zhou,
J.X.
Deng,
Y.
Liu,
W.T.
Li,
X.
Chu,
P.
Xiu,
H.
Cai,
Y.H.
Kou,
B.G.
Jiang,
Y.F.
Zheng,
Acta
Biomater.
64
(2017)
421–436.
[71]
P.
Trumbo,
A.A.
Yates,
S.
Schlicker,
M.
Poos,
J.
Am.
Diet.
Assoc.
101
(3)
(2001)
294–301.
[72]
H.
Tapiero,
K.D.
Tew,
Biomed.
Pharmacother.
57
(2003)
399–411.
[73]
J.
Lin,
S.
Ozan,
Y.
Li,
D.
Ping,
X.
Tong,
G.
Li,
C.
Wen,
Sci.
Rep.
6
(2016)
37901.
[74]
J.D.
Termine,
A.B.
Belcourt,
K.M.
Conn,
H.K.
Kleinman,
J.
Biol.
Chem.
256
(20)
(1981)
10403–10408.
[75]
N.C.
Blumenthal,
V.
Cosma,
J.
Biomed.
Mater.
Res.
23
(S13)
(1989)
13–22.
[76]
H.
Liao,
T.
Wurtz,
J.
Li,
J.
Biomed.
Mater.
Res.
47
(2)
(1999)
220–227.
[77]
A.
MacNicoll,
M.
Kelly,
H.
Aksoy,
E.
Kramer,
H.
Bouwmeester,
Q.
Chaudhry,
J.
Nanopart.
Res.
17
(2)
(2015)
66.
[78]
J.L.
Woodman,
J.J.
Jacobs,
J.O.
Galante,
R.M.
Urban,
J.
Orthop.
Res.
1
(4)
(1983)
421–430.
[79]
Y.
Oba,
A.
Yasue,
K.
Kaneko,
R.
Uchida,
A.
Shioyasono,
K.
Moriyama,
Orthod.
Waves
67
(1)
(2008)
1–8.
[80]
G.
Jin,
H.
Qin,
H.
Cao,
S.
Qian,
Y.
Zhao,
X.
Peng,
P.K.
Chu,
Biomaterials
35
(27)
(2014)
7699–7713.
229