Supplemental Materials
The alloy studied was Ti-5.at% Al, loaded in double slip orientation. For
micron/submicron samples, the pillars had square cross sections with the edge length
(L) or diameter (d) ranging from 5.0 μm to 150 nm, and an aspect ratio of ~2. In situ
Transmission Electron Microscopy (TEM) tensile samples had width around 200 nm.
All samples were cut out from a bulk single crystal with well defined orientation
determined using the method described in previous work (10-11). The samples were
loaded in approximately the “double slip” orientation, as two equivalent prismatic slip
systems, (01-10)[2-1-10] and (-1100)[11-20] are expected to have the highest
propensity to be activated (1,2). Larger samples were tested in compression in a MTS
Nanoindenter XP device outfitted with a flat punch diamond tip, at a nominal axial
strain rate ~ 110-4s-1 at room temperature. Cross-sectional TEM samples of these
deformed pillars were produced by Focus Ion Beam (FIB) using the “lift out” method
and analyzed under a JEOL 3010 TEM. In the same TEM, samples with sizes below
200 nm were compressed or pulled by using a Hysitron Picoindenter in the method
that has been described before (3,4). All mechanical tests were under displacementcontrolled mode. Further thinning was performed in a Triple-beam Focus Ion Beam
with final milling by using 2KV Ar+ source to reduce the FIB damage, for atomic
resolution imaging of the microstructure of the deformed samples in a Philips CM 300
high-resolution TEM (HRTEM).
We first note that slip systems 1/3(01-10)[2-1-10] and 1/3(-1100)[11-20] on prismatic
planes would be the first ones to be activated. Dislocations from these two slip
systems will intersect with each other and form junctions, and in bulk samples the
large number of junctions formed will produce a rather robust dislocation network
that leads to stain hardening and stable flow (2). This phenomenon can be also
observed in our micron-sized samples. Figure S1a shows the interlaced strain field
produced by the dislocation network in one micron-sized pillar (the two dislocation
types are marked). However, with continued deformation and rising stresses the
pinning at the limited number of junctions eventually loses its grip, such that the
dislocations eventually pulls away in an avalanche, producing uncontrollable strain
bursts (Fig. 1a in main text) in lieu of sustainable/stable plastic strain. Cross-sectional
TEM of the micron-sized deformed samples revealed a large number of “U” and “H”
shaped dislocations. Examples are shown in Figure S1b and Fig. 3a in the main text.
Different from that in bulk samples, it appears that the interaction and tangling
between dislocations is weaker in our small samples, due to the limited chances for
the dislocations to intersect and form junctions in the small volume. A dislocation
would encounter and be pinned by only a couple of dislocations from the other slip
system, leaving two long single arms on both ends. Once the applied stress is enough
to overcome the barrier from the single pinning point at the dislocations intersection,
the arms could escape and continue to slip until leaving the sample from the surface.
In other words, the cross-sectional TEM observations in Figure S1b indicate no
multiple-pinned segments left behind, consistent with very limited number of
junctions. As such, the dislocation groups resulting from deformation can easily
disintegrate, resulting in collective dislocation avalanche, or a pronounced strain burst,
in a way similar to those in face centered cubic (FCC) and body centered cubic (BCC)
pillars. With decreasing sample size, the dislocations became even less organized and
retained.
Figure S1 (a) Typical HRTEM image that shows the interlaced strain field
produced by the dislocation network in one micron-sized pillar (the two dislocation
types are marked). (b) The TEM image of the cross section of one micro-pillar after
deformation. “U” and “H” shaped dislocations are marked. g=[4-2-20].
In addition, according to the E. A. Metzbower’s work (5), the stacking fault
probability a is the function of the aluminum content CAl as a = 0.003exp(0.133CAl),
here CAl is the weight percent (wt pct). It is also reported that the stacking faults
energy  is assumed to have an inverse relationship (6) with stacking fault probability
a as = K/a. K can be determined to be 0.9 mJ/m2 by the pure Ti with = 300 mJ/m2,
a = 0.003. Thus, with our present sample Ti-5.0at% of CAl = 2.9 wt pct, the calculated
stacking faults energy  is assuming to be decreased to 204 mJ/m2. Although the
stacking faults of Ti-5.0at% alloy is decreased compared to the pure Ti, it is still quite
high in comparison with some other metals such as Ag (22 mJ/m2) (7), Cu (55 mJ/m2)
(7) and Al (166 mJ/m2) (8).
Reference:
1 X. Tan, H. Gu, C. Laird, N.D.H. Munroe, Metallurgical and Materials Transactions A
29A, 507(1998).
2 L. Xiao, Y. Umakoshiz, Philosophical Magazine, 82, 2379(2002).
3 Q. Yu, Z.W. Shan, J. Li, X. X. Huang, L. Xiao, J. Sun, E. Ma, Nature, 351, 335(2010).
4 D. Kiener, A.M. Minor, Nanoletters, 11, 3816(2011).
5 E. A. Metzbower, Metall. Trans., 2, 3099 (1971).
6 R.P. Adler, H.M. Otte, C.N.J. Wagner, Metall. Trans., 1, 2375 (1970).
7 P. C. J. Gallagher, Metall. Trans., 1, 2429 (1970).
8 L. E. Murr, Acta Metall., 21, 791 (1973).