Ultrafine-grained titanium of high interstitial contents with a good combination of strength and ductility

Article abstract of DOI:10.1063/1.2832657

A dehydrided Ti powder of very high oxygen content was successfully consolidated using back pressure equal channel angular processing into a fully dense bulk ultrafine-grained Ti showing apparent compressive ductility as well as high true yield and ultimate strengths of 1350 and 1780 MPa, respectively. Interstitial solid solution strengthening contributed to the majority of the increase in strength with additional contribution from ultrafine grains. Significantly, the material also exhibited much improved ductility for such a high interstitial content, thanks probably to the nonequilibrium grain boundaries and bimodal grain structure introduced during severe plastic deformation.

Full text of DOI:10.1063/1.2832657

Ultrafine-grained titanium of high interstitial contents with a good combination of
strength and ductility
W. Xu, X. Wu, D. Sadedin, G. Wellwood, and K. Xia
Citation: Applied Physics Letters 92, 011924 (2008); doi: 10.1063/1.2832657
View online: http://dx.doi.org/10.1063/1.2832657
View Table of Contents: http://scitation.aip.org/content/aip/journal/apl/92/1?ver=pdfcov
Published by the AIP Publishing
Articles you may be interested in
Sn and Nb modified ultrafine Ti-based bulk alloys with high-strength and enhanced ductility
Appl. Phys. Lett. 102, 061908 (2013); 10.1063/1.4792592
High strength and utilizable ductility of bulk ultrafine-grained Cu–Al alloys
Appl. Phys. Lett. 92, 201915 (2008); 10.1063/1.2936306
Tougher ultrafine grain Cu via high-angle grain boundaries and low dislocation density
Appl. Phys. Lett. 92, 081903 (2008); 10.1063/1.2870014
Ductile ultrafine-grained Ti-based alloys with high yield strength
Appl. Phys. Lett. 91, 051906 (2007); 10.1063/1.2766861
Compressive strength measurements in aluminum for shock compression over the stress range of 4 – 22 GPa
J. Appl. Phys. 98, 033524 (2005); 10.1063/1.2001729
This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to IP:
142.150.190.39 On: Sat, 20 Dec 2014 09:16:24

APPLIED PHYSICS LETTERS 92, 011924 ͑2008͒
Ultrafine-grained titanium of high interstitial contents with a good
combination of strength and ductility
W. Xu,1,a͒ X. Wu,1 D. Sadedin,2 G. Wellwood,3 and K. Xia1,b͒
1Department of Mechanical and Manufacturing Engineering and ARC Centre of Excellence for Design
in Light Metals, University of Melbourne, Victoria 3010, Australia
2Department of Chemical and Biomolecular Engineering, University of Melbourne, Victoria 3010, Australia
3Division of Minerals, Commonwealth Scientific and Industrial Research Organization (CSIRO),
Clayton South, Victoria 3169, Australia
͑Received 21 November 2007; accepted 17 December 2007; published online 10 January 2008͒
A dehydrided Ti powder of very high oxygen content was successfully consolidated using back
pressure equal channel angular processing into a fully dense bulk ultrafine-grained Ti showing
apparent compressive ductility as well as high true yield and ultimate strengths of 1350 and
1780 MPa, respectively. Interstitial solid solution strengthening contributed to the majority of the
increase in strength with additional contribution from ultrafine grains. Significantly, the material
also exhibited much improved ductility for such a high interstitial content, thanks probably to the
nonequilibrium grain boundaries and bimodal grain structure introduced during severe plastic
deformation. © 2008 American Institute of Physics. ͓DOI: 10.1063/1.2832657͔
Ultrafine-grained/nanostructured ͑UFG/NS͒ metallic ma-
terials have been the subject of extensive research thanks to
their superior mechanical properties compared to their
coarse-grained counterparts.1–3 Due to the high strength-to-
density ratio and excellent biocompatibility for structural and
implant applications, high performance titanium and its al-
loys have been synthesized by controlling at least one micro-
structural length scale down to the submicron/nanometer
regime.4–8 Ti-based UFG/NS composites, such as Ti–Cu–Ni–
Sn–Nb ͑Ref. 5͒ and Ti–Fe–Sn ͑Ref. 8͒ alloys, exhibited very
promising combined mechanical properties with ultimate
compressive strength over 2200 MPa and plastic strain
higher than 8%. However, most Ti-based UFG/NS compos-
ites are multicomponent alloys containing high content of
cytotoxic alloying elements such as Ni, Cu, or Fe, which are
not biocompatible and, therefore, not desirable for implant
applications.9 Recently, increasing interests in substantial
grain refinement using equal channel angular pressing have
opened a promising route to producing bulk UFG/NS mate-
rials with a good combination of mechanical properties.3,4
Especially, some of the intrinsically brittle intermetallics ex-
hibited certain ductility after severe plastic deformation
͑SPD͒.1 In this letter, back pressure ͑BP͒ equal channel an-
gular ͑ECA͒ processing was employed to consolidate a de-
hydrided ͑DH͒ Ti powder of very high oxygen content into a
fully dense bulk UFG-Ti exhibiting enhanced strength as
well as apparent compressive ductility.
tents ͑0.71 wt. % O and 0.04 wt. % N͒ was consolidated
for one pass at 600 °C with a back pressure of 50 MPa.
The contents of oxygen and nitrogen were analyzed using a
carrier gas hot extraction ͑O/N analyzer/Leco͒ for consoli-
dated DH-Ti ͑1.34 wt. % O and 0.3 wt. % N͒ and CP-Ti
͑0.73 wt. % O and 0.07 wt. % N͒. For DH-Ti, the increases
in oxygen and nitrogen contents were believed to originate
from the entrapped air during BP-ECA processing of ul-
trafine particles.10 Density of the consolidated samples was
measured using the Archimedes’ method. Room-temperature
mechanical properties were characterized under quasistatic
uniaxial compression loading using a MTS testing machine
at a strain rate of 3.0ϫ10−3 s−1. The Vickers microhardness
͑HV͒ was measured at room temperature under a static load
of 100 g for 25 s. Microstructural and fractographic analyses
were conducted using x-ray diffractometry ͑XRD͒ ͑Cu K␣͒,
optical microscopy, scanning electron microscopy ͑FEI
Quanta, 30 kV͒ and transmission electron microscopy
͑TEM͒ ͑FEI Tecnai F20, 200 kV͒.
Densities of both consolidated CP-Ti ͑4.48 g/cm3͒ and
DH-Ti ͑4.53 g/cm3͒ were very close to the value
͑4.51–4.54 g/cm3͒ of commercially pure Ti,11 indicating a
high efficiency of producing dense materials by BP-ECA
processing. The Rietveld refinement12 of XRD patterns
yielded an average crystallite size of 80 nm for both consoli-
dated CP-Ti and DH-Ti. CP-Ti exhibited a microstructure
consisting of elongated particles in both the longitudinal
͓Fig. 1͑a͔͒ and cross ͓Fig. 1͑b͔͒ sections with a substantial
amount of mechanical twins. Figures 1͑c͒ and 1͑d͒ show
TEM images and corresponding selected-area diffraction
͑SAD͒ patterns in the cross section. Finer elongated grains
͓Fig. 1͑c͔͒ were revealed in localized regions, possibly de-
rived from deformation twinning.7 The SAD pattern dis-
played elongated and clustered diffraction spots, implying
the presence of low-angle grain boundaries and lattice distor-
tions. Most of the elongated grains in Fig. 1͑c͒ were, there-
fore, viewed as subgrains with low-angle misorientations.
Figure 1͑d͒ shows a region comprising domains subdivided
by high-density dislocations. The SAD pattern revealed that
A dehydrided Ti powder ͑1.15 wt. % O and 0.09 wt. %
N͒ was produced through dehydrogenation of commercial
TiH2 powder ͑purity of 98%, Sigma-Aldrich͒ under vacuum
at 650 °C for 10 h followed by ring milling for 2 min at
room temperature. The average crystallite size of the DH-Ti
powder was determined to be 130 nm using x-ray diffraction.
Consolidation of the DH-Ti powder was carried out using
BP-ECA processing in air for one pass at 630 °C with a back
pressure of 200 MPa. For comparison, a micrometer-sized
commercially pure ͑CP͒ Ti powder of lower interstitial con-
a͒
Electronic mail: weixu@unimelb.edu.au.
Electronic mail: k.xia@unimelb.edu.au.
b͒
This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to IP:
0003-6951/2008/92͑1͒/011924/3/$23.00 92, 011924-1 © 2008 American Institute of Physics
142.150.190.39 On: Sat, 20 Dec 2014 09:16:24

011924-2
Xu et al.
Appl. Phys. Lett. 92, 011924 ͑2008͒
FIG. 1. ͓͑a͒ and ͑b͔͒ Optical micrographs in the longitudinal and cross
sections, respectively. ͓͑c͒ and ͑d͔͒ TEM bright-field images in the cross
section and corresponding SAD patterns ͑the insets͒ of the BP-ECA consoli-
dated CP-Ti.
the less sharp domain boundaries had characteristics of low-
angle grain boundaries.
In contrast, although elongated particles were also ob-
served in the longitudinal section of DH-Ti ͓Fig. 2͑a͔͒, these
particles had a finer interparticle spacing and a shorter par-
ticle length owing to smaller powder size. Most of the par-
ticles in the cross section ͓Fig. 2͑b͔͒ were irregular but near
equiaxed and particularly very fine particles filled in the in-
terparticle regions of coarse ones, forming a homogeneous
bimodal structure. Unlike CP-Ti, deformation twins were sel-
dom observed in DH-Ti except in several big particles. The
TEM examination ͓Fig. 2͑c͔͒ revealed the presence of equi-
axed grains of several micrometers in size. The SAD dotted
pattern implied most grains had high-angle boundaries. It is
noteworthy that ultrafine grains were found to uniformly dis-
tribute mostly at the triple grain junctions ͓Figs. 2͑c͒–2͑e͔͒,
reflecting a distinct difference in grain size of the DH-Ti
powder. The spreading of boundary thickness extinction con-
tours of some ultrafine grains ͓Fig. 2͑d͔͒ was characteristic of
nonequilibrium grain boundaries and was suggested to cor-
relate with a high level of elastic strain and lattice distortion.1
The volume fraction of the ultrafine grains was estimated to
be about 30%. Figure 2͑f͒ shows the HRTEM image of an
ultrafine grain and corresponding fast Fourier transformation
͑FFT͒ pattern from which the ultrafine grain was identified as
␣-Ti with the ͓121͔ zone axis.
FIG. 2. ͓͑a͒ and ͑b͔͒ Optical micrographs in the longitudinal and cross
sections, respectively. ͓͑c͒ and ͑d͔͒ TEM bright-field and ͑e͒ dark-field im-
ages in the cross section and corresponding SAD pattern ͑inset͒. ͑f͒ HRTEM
image of one ultrafine grain and corresponding FFT pattern ͑inset͒ of the
BP-ECA consolidated DH-Ti.
flutes connecting cleavage planes ͓Fig. 3͑b͔͒ were prominent
throughout the fracture surface together with frequently ob-
served shallow dimples ͓Fig. 3͑c͔͒. The fluted fracture, con-
sidered as a result of low-energy ductile rupture, was re-
ported to occur under mechanical overload in ␣-Ti of high
oxygen content that hindered cross slip and promoted planar
slip.13 The presence of flutes and shallow dimples implied
that localized plastic deformation happened in DH-Ti, lead-
ing to some limited ductility. In contrast, the rupture surface
of CP-Ti showed both cleavages and elongated dimples ͑not
shown͒ without visible fluted areas. HV values were mea-
sured to be 3207 137 MPa for CP-Ti and 5129 216 MPa
for DH-Ti, in agreement with yield strengths.
Incorporating interstitials in the lattice of Ti may result
in a substantial increase in strength but a drastic reduction in
ductility.14 Interstitially dissolved oxygen of 0.8 wt. % or ni-
trogen of 0.4 wt. % caused severe embrittlement of Ti with-
out any sign of plastic deformation.14 However, the ECA
consolidated DH-Ti exhibited enhanced strength and hard-
ness and at the same time retained a certain degree of duc-
tility in the presence of very high interstitial content. Based
on the correlation between Vickers hardness and equivalent
oxygen content,14 the hardness of DH-Ti of 1.34 wt. % O
and 0.3 wt. % N was estimated to be 4872 MPa, slightly
lower than the measured one. It suggests that strength en-
Figure 3͑a͒ exhibits the room-temperature compressive
true stress-strain curves for DH-Ti and CP-Ti. CP-Ti showed
yield strength ͑␴
offset͒ of 630 MPa and ultimate
0.2%
strength of 952 MPa. In contrast, remarkably enhanced
strength was achieved in DH-Ti with yield strength of
1350 MPa and ultimate strength reaching 1780 MPa. Both
DH-Ti and CP-Ti showed apparent compressive ductility.
Fractography indicated a mixed mode of brittle and ductile
fracture for DH-Ti. Although cleavage and intergranular
This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to IP:
fracture features ͑not shown͒ were predominant, very fine
hancement is largely attributable to interstitial solid solution
142.150.190.39 On: Sat, 20 Dec 2014 09:16:24

011924-3
Xu et al.
Appl. Phys. Lett. 92, 011924 ͑2008͒
In summary, a dehydrided Ti powder of very high oxy-
gen content ͑1.15 wt. %͒ was consolidated using BP-ECA
processing into a fully dense bulk UFG-Ti with a bimodal
grain structure, showing significantly enhanced strength to-
gether with apparent ductility in the presence of high oxygen
͑1.34 wt. %͒ and nitrogen ͑0.3 wt. %͒. The strength en-
hancement was mainly derived from interstitial solid solution
strengthening although ultrafine grains also contributed. The
ductility attained was likely to originate from the non-
equilibrium grain boundary configuration and the bimodal
grain structure. The good combination of high strength and
improved ductility makes BP-ECA processing very promis-
ing in producing low-cost Ti and its alloys of high interstitial
content.
The authors greatly appreciate the financial support from
the Australian Research Council with the ARC Centre of
Excellence for Design in Light Metals. Sincere thanks are
also due to Dr. C. Wen and Dr. Y. Li at Deakin University for
compressive testing.
FIG. 3. ͑a͒ Room-temperature compressive true stress-strain curves for the
BP-ECA consolidated DH-Ti and CP-Ti. ͓͑b͒ and ͑c͔͒ fracture surfaces
showing flutes and shallow dimples, respectively, in the BP-ECA consoli-
dated DH-Ti.
1R. Z. Valiev, R. K. Islamgaliev, and I. V. Alexandrov, Prog. Mater. Sci.
45, 103 ͑2000͒.
2Y. M. Wang, M. W. Chen, F. H. Zhou, and E. Ma, Nature ͑London͒ 419,
912 ͑2002͒.
3R. Z. Valiev and T. G. Langdon, Prog. Mater. Sci. 51, 881 ͑2006͒.
4D. Jia, Y. M. Wang, K. T. Ramesh, E. Ma, Y. T. Zhu, and R. Z. Valiev,
Appl. Phys. Lett. 79, 611 ͑2001͒.
strengthening although the ultrafine grains also contribute.
The attained ductility, on the other hand, was probably due to
the nonequilibrium grain boundaries introduced by SPD as
well as the bimodal grain structure. It was suggested that
these grain boundaries were in a high-energy nonequilibrium
configuration comprising an excess of dislocations and other
defects.1,3 Grain boundary sliding was proposed to be the
possible deformation mode in SPD processed UFG/NS
materials,15,16 resulting in a certain level of ductility even
in some intrinsically brittle intermetallics.1 In the present
DH-Ti, grain boundary sliding might be influenced by two
factors. Grain boundaries may act as potential sinks for in-
terstitials and a large amount of interstitials accumulating at
grain boundaries were found to retard grain boundary
sliding.17 On the other hand, high-density dislocations near
those nonequilibrium grain boundaries were suggested to
significantly facilitate grain boundary sliding15 to produce
certain ductility. With respect to the attained ductility, it
seems that the second factor predominates. However, the ac-
tual deformation mechanisms are likely to be complicated
and experimental verifications of grain boundary sliding are
imperative.
5G. He, J. Eckert, W. Löser, and L. Schultz, Nat. Mater. 2, 33 ͑2003͒.
6Y. T. Zhu, Y. R. Kolobov, G. P. Grabovetskaya, V. V. Stolyarov, N. V.
Girsova, and R. Z. Valiev, J. Mater. Res. 18, 1011 ͑2003͒.
7D. H. Shin, I. Kim, J. Kim, Y. S. Kim, and S. L. Semiatin, Acta Mater. 51,
983 ͑2003͒.
8J. Das, K. B. Kim, F. Baier, W. Löser, and J. Eckert, Appl. Phys. Lett. 87,
161907 ͑2005͒.
9D. Kuroda, M. Niinomi, M. Morinaga, Y. Kato, and T. Yashiro, Mater. Sci.
Eng., A 243, 244 ͑1998͒.
10W. Xu, T. Honma, X. Wu, S. P. Ringer, and K. Xia, Appl. Phys. Lett. 91,
031901 ͑2007͒.
11Metals Reference Book, 3rd ed. ͑ASM International, Materials Park,
1993͒, p. 153; W. F. Smith and J. Hashemi, Foundations of Materials
Science and Engineering, 4th ed. ͑McGraw-Hill, New York, 2006͒, p. 438.
12L. B. McCusker, R. B. Von Dreele, D. E. Cox, D. Louër, and P. Scardi, J.
Appl. Crystallogr. 32, 36 ͑1999͒.
13D. A. Meyn and E. J. Brooks, Fractography and Materials Science, edited
by L. N. Gilbertson and R. D. Zipp ͑ASTM STP 733, Philadelphia, 1981͒,
pp. 5–31.
14H. Conrad, J. Mater. Res. 26, 123 ͑1981͒.
15R. Z. Valiev, I. V. Alexandrov, Y. T. Zhu, and T. C. Lowe, J. Mater. Res.
17, 5 ͑2002͒.
16N. Q. Chinh, P. Szommer, Z. Horita, and T. G. Langdon, Adv. Mater.
͑Weinheim, Ger.͒ 18, 34 ͑2006͒.
17J. Monk, B. Hyde, and D. Farkas, J. Mater. Sci. 41, 7741 ͑2006͒.
This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to IP:
142.150.190.39 On: Sat, 20 Dec 2014 09:16:24