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41.We acknowledge technical and scientific assistance from
S.L.Kearns,J.Levallois,and N.Mangkorntang and
collaborative support from H.H.Wen.This work was
supported by Engineering and Physical Sciences Rearch
Council(UK),the Royal Society,Laboratoire National des
Champs Magnétiques Pulsés,the French Agence
Nationale de la Recherche IceNET,and EuroMagNET.
Supporting Online Material
www.sciencemag/cgi/content/full/1165015/DC1
Materials and Methods
Figs.S1and S2
References
22August2008;accepted21November2008
Published online11December2008;
10.1126/science.1165015
Include this information when citing this paper.
Revealing the Maximum Strength
in Nanotwinned Copper
L.Lu,1*X.Chen,1X.Huang,2K.Lu1
The strength of polycrystalline materials increas with decreasing grain size.Below a critical size,smaller grains might lead to softening,as suggested by atomistic simulations.The strongest size should ari at a transition in deformation mechanism from lattice dislocation activities to grain boundary–related process.We investigated the maximum strength of nanotwinned copper samples with different twin thickness.We found that the strength increas with decreasing twin thickness,reaching a maximum at 15nanometers,followed by a softening at smaller values that is accompanied by enhanced strain hardening and tensile ductility.The strongest twin thickness originates from a transition in the yielding mechanism from the slip transfer across twin boundaries to the activity of preexisting easy dislocation sources.
T he strength of polycrystalline materials increas with decreasing grain size,as
described by the well-known Hall-Petch relation(1,2).The strengthening originates from the fact that grain boundaries block the lattice dislocation motion,thereby making plastic defor-mation more difficult at smaller grain sizes.How-ever,below a certain critical size,the dominating deformation mechanism may change from lattice dislocation activities to other mechanisms such as grain boundary–related process,and softening behavior(rather than strengthening)is expected (3,4).Such a softening phenomenon has been demonstrated by atomistic simulations,and a crit-ical grain size of maximum strength has been predicted(5–7).In pure metals,an impediment to determining the grain size that yields the highest strength is the practical difficulty of obtaining sta-ble nanostructures with extremely small structural domains(on the order of veral nanometers).The driving force for growth of nanosized grains in pure metals,originating from the high excess en-ergy of numerous grain boundaries,becomes so large that grain growth may take place easily even at ambient temperature or below.
Coherent twin boundaries(TBs),which are
defined in a face-centered cubic structure as the
(111)mirror planes at which the normal stacking
quence of(111)planes is reverd,are known to
be as effective as conventional grain boundaries
in strengthening materials.Strengthening has been
obtained in Cu when high densities of nanometer-
thick twins are introduced into submicrometer-
sized grains(8–10).In addition,coherent TBs are
much more stable against migration(a fundamen-
tal process of coarning)than conventional grain
boundaries,as the excess energy of coherent TBs is
one order of magnitude lower than that of grain
boundaries.Hence,nanotwinned structures are
energetically more stable than nanograined coun-
terparts with the same chemical composition.The
初中美术stable nanotwinned structure may provide samples
for exploring the softening behavior with very small
domain sizes.Here,we prepared nanotwinned pure
Cu(nt-Cu)samples with average twin thickness
ranging from a few nanometers to about100nm.
High-purity(99.995%)Cu foil samples com-
pod of nanoscale twin lamellae embedded in
submicrometer-sized grains were synthesized by
means of puld electrodeposition.By increasing
the deposition rate to10nm/s,we succeeded in
refining the mean twin ,the mean
spacing between adjacent TBs,hereafter referred
to as l)from a range of15to100nm down to a
range of4to10nm(e supporting online ma-
terial).The as-deposited Cu foils have an in-plane
dimension of20mm by10mm and a thickness
of30m m with a uniform microstructure.Shown
in Fig.1,A to C,are transmission electron mi-
croscopy(TEM)plane-view images of three as-
deposited samples with l values of96nm,15nm,
and4nm,respectively.The TEM images indicate
that some grains are irregular in shape,but low-
magnification scanning electronic microscopy
images,both cross ction and plane view,show
that the grains are roughly equiaxed in three di-
mensions.Grain size measurements showed a
similar distribution and a similar average diam-
eter of about400to600nm for all nt-Cu samples.
Twins were formed in all grains(e the electron
diffraction pattern in Fig.1D),and obrvations
of twins in a large number of individual grains
revealed no obvious change in the twin density
from grain to grain.Note that in all samples,the
edge-on twins that formed in different grains are
aligned randomly around the foil normal(growth)
direction(8,11),in agreement with a strong[110]
texture determined by x-ray diffraction(XRD).For
each sample,twin thickness were measured from
a large number of grains,which were detected from
numerous TEM and high-resolution TEM(HRTEM)
images,to generate a distribution.Figure1E illus-
齐白石画画
trates the492measurements for the sample with
the finest twins;the majority yielded spacings be-
tween twins smaller than10nm,with a mean of
4nm.For simplicity,each nt-Cu sample is iden-
tified by its mean twin thickness;for example,the
sample with l=4nm is referred to as nt-4.
Figure2shows the uniaxial tensile true stress–
true strain curves for nt-Cu samples of various l
values.Also included are two stress-strain curves
obtained from a coar-grained Cu(cg-Cu)and
an ultrafine-grained Cu(ufg-Cu)that has a sim-
ilar grain size to that of nt-Cu samples but is free
of twins within grains.Two distinct features are
obrved with respect to the l dependence of the
mechanical behavior of nt-Cu.The first is the oc-
currence of the l giving the highest strength.All
道路标志大全图解
stress-strain curves of nt-Cu samples in Fig.2,A洒脱的古诗词
and B,are above that of the ufg-Cu,indicating a
strengthening by introducing twins into the sub-
micrometer grains.However,such a strengthen-
ing does not show a linear relationship with l.For
l>15nm(Fig.2A),the stress-strain curves shift
upward with decreasing l,similar to the strength-
ening behavior reported previously in the nt-Cu
(9,11)and nanocrystalline Cu(nc-Cu)(12–15)
samples(Fig.3A).However,with further de-
1Shenyang National Laboratory for Materials Science,Institute of Metal Rearch,Chine Academy of Sciences,Shenyang 110016,P.R.China.2Center for Fundamental Rearch:Metal Structures in Four Dimensions,Materials Rearch Department, RisøNational Laboratory for Sustainable Energy,Technical Uni-versity of Denmark,DK-4000Roskilde,Denmark.
*To whom correspondence should be addresd.E-mail: llu@imr.ac o n J a n u a r y 3 0 , 2 0 0 9 w w w . s c i e n c e m a g . o r g D o w n l o a d e d f r o m
creas of l down to extreme dimensions (i.e.,less than 10nm),the stress-strain curves shift down-ward (Fig.2B).As plotted in Fig.3A,the mea-sured yield strength s y (at 0.2%offt)shows a maximum value of 900MPa at l ≈15nm.The cond feature is a substantial increa in tensile ductility and strain hardening when l <15nm.As en in Fig.2,the tensile elongation of the nt-Cu samples increa
s monotonically with decreasing l .When l <15nm,the uniform ten-sile elongation exceeds that of the ufg-Cu sample,reaching a maximum value of 30%at the finest twin thickness.Strain-hardening coefficient (n )values were determined for each sample by fitting the uniform plastic deformation region to s =K 1+K 2e n ,where K 1reprents the initial yield stress and K 2is the strengthening coefficient (i.e.,the strength increment due to strain hardening at strain e =1)(16,17).The n values determined for all the nt-Cu samples increa monotonically with de-creasing l (Fig.3B),similar to the trend of uniform elongation versus l .When l <15nm,n exceeds the value for cg-Cu (0.35)(16,17)and finally reaches a maximum of 0.66at l =4nm.The twin refinement –induced increa in n is opposite to the general obrvation in ultrafine-grained and nanocrystalline materials,where n continuously decreas with decreasing grain size (Fig.3B).The strength of the nt-Cu samples has been considered to be controlled predominantly by the nanoscale twins via the mechanism of slip trans-fer across the TBs (10,18),and it increas with decreasing l in a Hall-Petch –type relationship (9)similar to that of grain boundary strength-ening in nanocrystalline metals (12).Our re-sults show that such a relationship breaks down when l <15nm,although other structural pa-rameters such as grain size and texture are un-changed.The grain sizes of the nt-Cu samples are in the submicrometer regime,which is too large for grain boundary sliding to occur at room temperature,as expected for nanocrystalline ma-terials with grain sizes below 20nm (3).There-fore,the obrved softening cannot
be explained by the initiation of grain boundary –mediated mechanisms such as grain boundary sliding and grain rotation,as propod by molecular dynam-ics (MD)simulations for nanocrystalline mate-rials (3).
To explore the origin of the twin thickness giv-ing the highest strength,we carried out detailed structural characterization of the as-deposited sam-ples.HRTEM obrvations showed that in each sample TBs are coherent S 3interfaces associated with the prence of Shockley partial dislocations (as steps),as indicated in Fig.1D.The partial dislocations have their Burgers vector parallel to the twin plane and are an intrinsic structural fea-ture of twin growth during electrodeposition.The distribution of the preexisting partial dislocations is inhomogeneous,but their density per unit area of TBs is found to be rather constant among sam-ples with different twin densities.This suggests that the deposition parameters and the twin re-finement have a negligible effect on the nature of
TBs.Therefore,as a conquence of decreasing l ,the density of such TB-associated partial dislocations per unit volume increas.
We also noticed that grain boundaries in the nt-Cu samples with l ≤15nm are characterized by straight gments (facets)that are
often
True strain (%)T r u e s t r e s s (M P a )
True strain (%)
Fig.2.Uniaxial tensile true stress –true strain curves for nt-Cu samples tested at a strain rate of 6×10−3s −1.(A )Curves for samples with mean twin thickness varying from 15to 96nm;(B )curves for samples with mean twin thickness varying from 4to 15nm.For comparison,curves for a twin-free ufg-Cu with a mean grain size of 500nm and for a cg-Cu with a mean grain size of 10m m are
included.
C
200 nm
A B
b
b
D
Fig.1.TEM images of as-deposited Cu samples with 15nm.(C )l =4nm.(D )The same sample as (C)but at higher resolution,with a corresponding electron diffraction pattern (upper right int)and a HRTEM image of the outlined area showing the prence of Shockley partials at the TB (lower right int).(E )Distribution of the lamellar twin thickness determined from TEM and HRTEM images for l =4nm.
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associated with dislocation arrays (19),whereas in samples with coarr twins,grain boundaries are smoothly curved,similar to conventional grain boundaries.The microstrain measured by XRD was a n
egligible 0.01%for samples with l ≥15nm,but incread gradually from 0.038%to 0.057%when l decread from 10to 4nm,which also indicates a gradual increa in the de-fect density.
Recent experimental studies and MD simula-tions (3,20,21)have shown that an increa in the density of preexisting dislocations in nano-scale materials will cau softening.In the nt-Cu samples studied,both the dislocation arrays asso-ciated with the grain boundaries and the steps associated with the preexisting partial dislocations along TBs could be potential dislocation sources,which are expected to affect the initiation of plas-tic deformation (22)and to provide the disloca-tions required for the dislocation-TB interactions that cau work hardening.The preexisting par-tial dislocations can act as readily mobile dislo-cations,and their motion may contribute to the plastic yielding when an external stress is applied to the sample.The plastic strains induced by the motion of preexisting partial dislocations can be estimated as e =r 0b s d/M (where r 0is the initial dislocation density,b s is the Burgers vector of Shockley partial dislocation,d is the grain size,and M is the Taylor factor).Calculations showed that for the samples with l >15nm,the preexist-ing dislocations induce a negligibly small plastic strain (<0.05%).However,for the nt-4specimen,a remarkable amount of plastic strain,as high as 0.1to 0.2%,can be induced just by the motions of high-density preexisting dislocations at TBs (roughly 1014m −2),which could control the mac-roscopic yielding of the sample.
The above anal-ysis suggests that for extremely small values of l ,a transition in the yielding mechanism can result in an unusual softening phenomenon in which the preexisting easy dislocation sources at TBs and
grain boundaries dominate the plastic deforma-tion instead of the slip transfer across TBs.
Shockley partial dislocations are always in-volved in growth of twins during crystal growth,thermal annealing,or plastic deformation.Shock-ley partials might be left at TBs when the twin growth is interrupted.Therefore,the prence of Shockley partials at some TBs is a natural phe-nomenon.Although the preexisting dislocations may have a small effect on the mechanical be-havior of the samples with thick twins,the effect will be much more pronounced in the samples with nanoscale twins and/or with high preexist-ing TB dislocation densities such as tho en in deformation twins (23).
To understand the extraordinary strain hard-ening,we analyzed the deformation structures of the tensile-deformed samples.In samples with coar twins,tangles and networks of perfect dis-locations were obrved within the lattice between the TBs (Fig.4A),and the dislocation density was estimated to be on the order of 1014to 1015m −2.In contrast,high densities of stacking faults and Shockley part
ials associated with the TBs were found to characterize the deformed structure of the nt-4sample (Fig.4,B and C),indicating the interactions between dislocations and TBs.Recent MD simulations (18,24,25)showed that when an extended dislocation (two Shockley partials connected by a stacking fault ribbon)is forced by an external stress into a coherent TB,it recom-bines or constricts into a perfect dislocation con-figuration at the coherent TB and then slips through the boundary by splitting into three Shockley par-tials.Two of them glide in the slip plane of the adjacent twin lamella,constituting a new extended dislocation,whereas the third one,a twinning par-tial,glides along the TB and forms a step.It is expected that with increasing strain,such an in-teraction process will generate a high density of partial dislocations (steps)along TBs and stack-ing faults that align with the slip planes in the twin lamellae,which may (or may not)connect to the TBs.Such a configuration of defects was obrved,as shown in Fig.4C.The density of partial dislocations in the deformed nt-4sample
Fig.3.Variation of (A )
yield strength and (B )
strain hardening coef-ficient n as a function of mean twin thickness for the nt-Cu samples.For comparison,the yield strength and n values for
nc-Cu [▲(12),◀(13),
▶(14),and ◆(15)],ufg-Cu[▾(9)],andcg-Cu samples reported in the literature are included.
A maximum in the yield
stress is en for the
nt-Cu with l =15nm,
but this has not been
obrved for the nc-Cu,even when the grain size is as small as 10
nm.
0.0
0.2
0.40.6
0.8
n
or d (nm)
0200400600800
σy (M P a )
λ or d (nm)020406080
100120110100100010000
λ2 nm
T
T
B发财树怕冷吗
200 nm
A C
Fig.4.(A )A typical bright TEM image of the deformed nt-96sample showing the tangling of lattice disl
ocations.(B )An HRTEM image of the nt-4sample tensile-deformed to a plastic strain of 30%,showing a high density of stacking faults (SF)at the TB.(C )The arrangement of Shockley partials and stacking faults at TBs within the lamellae in the nt-4sample.Triangles,Shockley partial dislocations associated with stacking faults;⊥,partials with their Burgers vector parallel to the TB plane.REPORTS
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was estimated to be 5×1016m −2on the basis of the spacing between the neighboring partials and l .This is two orders of magnitude higher than that of the preexisting dislocations and the lattice dislocations stored in the coar twins.Such a finding suggests that decreasing the twin thick-ness facilitates the dislocation-TB interactions and affords more room for storage of dislocations,which sustain more pronounced strain hardening in the nt-Cu (26,27).
The obrvations suggest that the strain-hardening behavior of nt-Cu samples is governed by two competing process:dislocation-dislocation interaction hardening in coar twins,and dislocation-TB interaction hardening in fine twins.With a refining of l ,the contribution from the latter mech-anism
increas and eventually dominates the strain hardening,as revealed by the continuous increa of n values (Fig.3B).However,the former hard-ening mechanism usually leads to an inver trend,diminishing with size refinement (17).
Twins are not uncommon in nature,and they appear in various metals and alloys with different crystallographic structures.Extremely thin twin lamellae structures can possibly be achieved under proper conditions during crystal growth,plastic deformation,pha transformations,or thermal annealing of deformed structures.Our finding of the twin thickness giving maximum strength il-lustrates that the scale-dependent nature of plastic deformation of nanometer-scale materials is not necessarily related to grain boundary –mediated process.This finding also provides insight into the development of advanced nanostructured materials.
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把握今天
Rearch:Metal Structures in Four Dimensions (X.H.).We thank N.Hann,Z.Jin,W.Pantleon,and B.Ralph for stimulating discussions,X.Si and H.Ma for sample preparation,S.Zheng for TEM obrvations,and Y.Shen for conducting some of the tensile tests.
Supporting Online Material
www.sciencemag/cgi/content/full/323/5914/607/DC1Materials and Methods Table S1References
24October 2008;accepted 30December 200810.1126/science.1167641
Control of Graphene ’s Properties by Reversible Hydrogenation:Evidence for Graphane
D.C.Elias,1*R.R.Nair,1*T.M.G.Mohiuddin,1S.V.Morozov,2P.Blake,3M.P.Halsall,1A.C.Ferrari,4D.W.Boukhvalov,5M.I.Katsnelson,5A.K.Geim,1,3K.S.Novolov 1†Although graphite is known as one of the most chemically inert materials,we have found that graphene,a single atomic plane of graphite,can react with atomic hydrogen,which transforms this highly conductive zero-overlap mimetal into an insulator.Transmission electron microscopy reveals that the obtained graphene derivative (graphane)is crystalline and retains the hexagonal lattice,but its period becomes markedly shorter than that of graphene.The reaction with hydrogen is reversible,so that the original metallic state,the lattice spacing,and even the quantum Hall effect can
be restored by annealing.Our work illustrates the concept of graphene as a robust atomic-scale scaffold on the basis of which new two-dimensional crystals with designed electronic and other properties can be created by attaching other atoms and molecules.
G
raphene,a flat monolayer of carbon atoms tightly packed into a honeycomb lattice,continues to attract immen interest,most-ly becau of its unusual electronic properties and effects that ari from its truly atomic thick-ness (1).Chemical modification of graphene has been less explored,even though rearch on car-bon nanotubes suggests that graphene can be al-tered chemically without breaking its resilient C-C bonds.For example,graphene oxide is graphene denly covered with hydroxyl and other groups (2–6).Unfortunately,graphene oxide is strongly disordered,poorly conductive,and difficult to reduce to the original state (6).However,one can imagine atoms or molecules being attached to the atomic scaffold in a strictly periodic manner,which should result in a different electronic struc-ture and,esntially,a different crystalline mate-rial.Particularly elegant is the idea of attaching atomic hydrogen to each site of the graphene lattice to create graphane (7),which changes the hybridization of carbon atoms from sp 2into sp 3,thus removing the conducting p -bands and open-ing an energy gap (7,8).
Previously,absorption of hydrogen on gra-phitic surfaces was investigated mostly in con-junction with hydrogen storage,with the rearch focud on physisorbed molecular hydrogen (9–11).More recently,atomic hydrogen chem-isorbed on carbon nanotubes has been studied theoretically (12)as well as by a variety of exper-imental techniques including infrared (13),ultra-violet (14,15),and x-ray (16)spectroscopy and scanning tunneling microscopy (17).We report the reversible hydrogenation of single-layer graphene and obrved dramatic changes in its transport properties and in its electronic and atomic struc-ture,as evidenced by Raman spectroscopy and transmission electron microscopy (TEM).
Graphene crystals were prepared by u of micromechanical cleavage (18)of graphite on top of an oxidized Si substrate (300nm SiO 2)and then identified by their optical contrast (1,18)and distinctive Raman signatures (19).Three types of samples were ud:large (>20m m)crystals for Raman studies,the standard Hall bar de-vices 1m m in width (18),and free-standing mem-branes (20,21)for TEM.For details of sample fabrication,we refer to earlier work (18,20,21).
1
School of Physics and Astronomy,University of Manchester,M139PL,Manchester,UK.2Institute for Mi
croelectronics Tech-nology,142432Chernogolovka,Russia.3Manchester Centre for Mesoscience and Nanotechnology,University of Manches-ter,M139PL,Manchester,UK.4Department of Engineering,Cambridge University,9JJ Thomson Avenue,Cambridge CB3OFA,UK.5Institute for Molecules and Materials,Radboud University Nijmegen,6525ED Nijmegen,Netherlands.*The authors contributed equally to this work.
†To whom correspondence should be addresd.E-mail:Kostya@manchester.ac.uk
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