In situ observations of crack propagation
and role of grain boundary microstructure
in nickel embrittled by sulfur
Shigeaki Kobayashi, Tatsuya Maruyama,
Sakae Saito, Sadahiro Tsurekawa &Tadao Watanabe
Journal of Materials Science
Full Set - Includes `Journal of Materials
Science Letters'
ISSN 0022-2461
Volume 49
Number 11
J Mater Sci (2014) 49:4007-4017
DOI 10.1007/s10853-014-8056-z

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J Mater Sci (2014) 49:4007–4017
DOI 10.1007/s10853-014-8056-z

INTERGRANULAR AND INTERPHASE BOUNDARIESIn situ observations of crack propagation and role of grainboundary microstructure in nickel embrittled by sulfur
Shigeaki Kobayashi • Tatsuya Maruyama •Sakae Saito • Sadahiro Tsurekawa •Tadao Watanabe

Received: 24 September 2013 / Accepted: 15 January 2014 / Published online: 28 January 2014
Ó Springer Science+Business Media New York 2014

Abstract In situ observations of crack propagation in
sulfur-doped coarse-grained nickel were performed for the
specimens with grain boundary microstructure pre-determined by SEM/EBSD analysis. The role of grain boundary
microstructure was studied in the crack propagation innickel embrittled by grain boundary segregation of sulfur.
It was found that the main crack tends to predominantly
propagate along random boundaries, and the crack propagation rate can be locally accelerated at the grain boundary
network with a high connectivity of random boundaries.
On the other hand, the cracks can propagated along fracture-resistant low-R coincidence site lattice (CSL) boundary only when the trace of the grain boundary is arranged
being almost parallel to slip bands in the adjacent grains.
The local crack propagation rate was found to become
lower when a crack propagated along low-R CSL boundaries. Moreover, when the crack propagation is inhibited by
low-R CSL boundaries, the branching of propagating crack
occurs at partially cracked triple junctions. The crack
S. Kobayashi (&)  T. Maruyama  S. Saito
Division of Mechanical and Electrical Engineering, Department
of Innovative Engineering, Faculty of Engineering, Ashikaga
Institute of Technology, 268-1 Omae, Ashikaga,
Tochigi 326-8558, Japane-mail: skoba@ashitech.ac.jp
S. Tsurekawa
Department of Materials Science and Engineering, Graduate
School of Science and Technology, Kumamoto University,
Kumamoto 860-8555, JapanT. Watanabe
Key Laboratory for Anisotropy and Texture of Materials,
Northeastern University, Shenyang 110004, China
T. Watanabe
Tohoku University, Sendai, Japanpropagation can locally slow down due to the occurrence of
crack branching. The optimum grain boundary microstructure for the control of sulfur segregation-induced
brittle fracture is discussed on the basis of new findings
obtained from the in situ observations on crack propagationand fracture processes in polycrystalline nickel.

Introduction
Grain boundary fragility due to grain boundary segregationof solute or impurity atoms in engineering materials has
been extensively studied and discussed since early time ofthe history of grain boundary research [1]. It is well known
that grain boundary segregation of sulfur atoms causes
severe embrittlement and degradation of performancereliability in polycrystalline nickel-based alloys which areused as high temperature structural materials [2–4].
Quantitative experimental works of grain boundary segregation in metallic materials have been performed previously, particularly using Auger electron spectroscopy(AES) analysis since 1980s [5–7]. It has been revealed thatthe degree of grain boundary segregation of detrimental
elements at specific grain boundaries strongly depends onthe grain boundary character and structure in some alloy
systems such as Fe–Sn alloy [5, 6], Fe–Si alloy [7] andother metallic materials, as fully discussed in the mostrecent book on grain boundary segregation by Lejček [8].
This has been confirmed by experiments for grain boundary
segregation of sulfur in nickel [9, 10]. It has been reported
that the degree of segregation of sulfur at special (lowenergy) boundaries tends to be lower than that at general(high-energy) boundaries for the nickel-sulfur system [10].
Moreover, the resistance to segregation-induced intergranular fracture at low-R coincidence site lattice (CSL)

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boundaries with lower energy was found to be higher thanthose at random high-energy boundaries [11–15]. Therefore, the grain boundary segregation-induced embrittlement in polycrystalline materials can be controlled by grainboundary engineering (GBE) by manipulating at least‘‘grain boundary character distribution (GBCD)’’ and‘‘grain boundary connectivity’’, as documented in recent
reviews on the GBE [16, 17].
The usefulness of GBE for control of grain boundary
degradation phenomena such as intergranular corrosion has
been proved by many researchers [18–22], and discussed inrecent Watanabe’s review on GBE [17]. However, unfortunately, such a promising application of the GBE forcontrol of segregation-induced embrittlement has been notmuch attempted so far [23–26]. Quite recently, the present
authors have revealed that the fracture toughness of sulfurdoped ultrafine-grained nickel specimens produced byelectrodeposition and subsequent annealing is strongly
influenced by the fraction of low-R CSL boundaries [25].
Furthermore, we have investigated the effect of grainboundary microstructure on the fracture resistance of sulfur-doped ordinary polycrystalline nickel, in order toestablish the utility of GBE for the control of segregationinduced embrittlement in polycrystalline materials [26]. In
the course of the previous studies, we have introduced a
new grain boundary microstructural parameter that canquantitatively describe the connectivity of fracture susceptive random boundaries on the basis of fractal analysis.
We have found that the random boundary network withmaximum connectivity (maximum random boundary connectivity, MRBC) is of fractal nature reflecting the distributions of a large variety of grain size and grain shape. It
has been revealed that the fractal dimension of MRBC canbe effectively used to predict and control the fracture
resistance and intergranular brittleness in polycrystalline
nickel embrittled by grain boundary segregation of sulfur[26]. However, these findings for the fracture processes in
sulfur-doped nickel must be confirmed further by experiment, if possible, through in situ direct observations of thecrack propagation, which is the percolation process in grainboundary network, probably affected by the GBCD andgrain boundary connectivity associated with the spatial
distribution of high-energy random boundaries or lowenergy specific grain boundaries.
In the present work, in situ observations of crack
propagation were performed for the sulfur-doped polycrystalline nickel in which individual grain boundaries
were characterized using the field emission gun-scanningelectron microscopy (FEG-SEM)/electron backscattereddiffraction (EBSD)/orientation imaging microscopy (OIM),
in order to reveal the roles of grain boundary microstructure in crack propagation processes in nickel embrittled bysulfur segregation. The roles of different types of grain

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boundaries and triple junctions in crack deflection andbranching were investigated in focusing our attention to thevariety of nature of the grain boundary network in polycrystals. Of particular interest was the effect of the grainboundary connectivity on the local propagation rate ofcracks in fracture processes of grain boundary network.
The importance of grain boundary microstructure for thecontrol of segregation-induced intergranular fracture inpolycrystalline materials is discussed on the basis of new
findings obtained from in situ observations in nickel.

Experimental procedureSpecimen preparation
A fully annealed polycrystalline nickel sheet with thepurity of 99.75 % was subjected to cold rolling to 75 % in
reduction ratio and subsequent annealing at 1273 K (0.74
Tm, Tm is a melting point of nickel) for 10.8 ks. The
specimen and sulfur powders (0.04 g) were enclosed
together in a Pyrex glass tube with 8 mm internal diameterand 70 mm length in a vacuum of 3 9 10-3 Pa. Sulfur
powders were put in an aluminum foil boat, to prevent thenickel specimen from direct contact with sulfur powders.
They were heated at 873 K (0.51 Tm) for 72 h. This heat
treatment could introduce the sulfur content of 290 ppm
into the nickel specimen. The equilibrium sulfur content at
grain boundaries was roughly estimated to be 95 at% using
the Langmuir-McLean equation [8]. However, it is very
likely that the condition of sulfur-doping heat treatment
carried out in the present work might not be sufficient toreach the equilibrium state. Using Auger electron spectroscopy (AES), Loier and Boos [4] have determined andreported that the nickel specimens with the bulk sulfur
concentration of 70 and 100 ppm possessed the intergranular sulfur concentration of 10.0 and 14.5 at%, respectively, even after heating at 1473 K for 20 days. Therefore,
the level of the grin boundary sulfur concentration should
be at least more than 10.0 at% order, when polycrystalline
nickel specimens with sulfur exhibit sulfur-induced intergranular fracture. Moreover, as simply predicted by the
Seah-Hondros diagram for the Ni–S system, it is very
likely that the presence of sulfide phase (such as Ni3S2 withlow-melting temperature) at grain boundaries can enhanceintergranular fracture [27].
The single-edge-notched specimens whose dimensions
were 40.0 mm length, 4.0 mm wide and 3.0 mm thick were
cut out by a spark machine from the heat treated sheet forcrack propagation test. A notch with 0.3 mm wide and
0.5 mm depth was introduced at the center of the longitudinal direction of the specimen. In order to characterize thegrain boundary microstructure by SEM/EBSD analyses, theAuthor's personal copy
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surface of the specimens was mechanically polished using
emery papers of 320–1000 grade and 0.06 lm alumina
powder slurry. Then specimen surfaces were electrolytically polished in an electrolytic solution of 79 vol% acetic
acid, 15 vol% methanol, and 6 vol% perchloric acid at acurrent density of 4.0 mA/mm2 at 280 K for 20 s. Finally,
they were weakly etched using a solution of 50 vol% acetic
acid and 50 vol% nitric acid at room temperature for 3 s, toreveal the grain boundary configuration.
Characterization of grain boundary microstructure
The grain boundary microstructure in the studied sulfurdoped nickel specimens was quantitatively characterizedusing FE-SEM/EBSD/OIM system. Based on the relativeorientation relationship and the misorientation angle
between adjoining grains, individual grain boundaries areconventionally defined as low-angle (h \ 15o, R1), highangle special/low-R (3 B R B 29) CSL boundaries, andhigh-angle general/random boundaries (h [ 15o, without
any CSL relationship), based on coincidence lattice site
relationship [28], under the Brandon’s criterion, Dh = 15/
R1/2 [29]. We have already obtained much experimental
evidence for the finding that low-R CSL boundaries(R \ 29) show special properties which are never observed
for random boundaries, although, strictly speaking, thegrain boundary character is not fully defined by the relativeorientation relationship and R-value. We need some
information on the grain boundary plane. However, theeffectiveness of GBCD evaluated by the conventional
characterization method of grain boundary microstructure
has been well proved in recent attempts of GBE for fracture
control and high performance [30–33].
In situ observations of fatigue crack propagationin sulfur-doped nickel specimens
Crack propagation tests under cyclic loading were carriedout using a servo-hydraulic machine (Shimadzu, Servopulser). Four-point bending load was cyclically applied tothe sulfur-doped nickel specimens. The gauge distances ofthe four-point bending jig were 30 mm between the outer
supporting pins and 10 mm between the inner loading pins.
Sinusoidal load was applied in air at room temperature, at aconstant load amplitude (31.5 N), load ratio (0.1), andfrequency (5 Hz).
In situ observations of crack propagation processes(intergranular fracture and/or transgranular fracture) weremade using the 2D-motion capture system (Detect, MotionPro) consisting of a CCD camera mounting a high magnification/long focal distance lens and a computer with an
A/D converter and a video capture card. A schematic
illustration of this observation system is shown in Fig. 1.

Fig. 1 Schematic illustration of the experimental set-up used forin situ observation of crack propagation during cyclic bending test

The observed images of cracks during cyclic loading were
captured in the computer at a regular cycle interval via
CCD camera and video capture card, only when the sinusoidal load reached to a preset value. Every step of crack
propagation was captured and evaluated in connection withthe characteristic features of grain boundary microstructures determined for each specimen prior to fracture test
under cyclic loading.

Results and discussion
Grain boundary microstructures of sulfur-doped nickel
specimen
Figure 2a–c shows the grain boundary micrograph, thegrain size distribution, the inverse pole figure with GBCD
data for the sulfur-doped nickel specimen, respectively. In
the grain boundary micrograph, Fig. 2a, the types of grainboundaries are shown by distinct colors corresponding tothose indicated in the color bars on the right-hand side of
the micrograph. For example, low-angle (R1), R3 CSL
boundaries, and random boundaries are shown by red,green and black lines, respectively. It was found that thesulfur-doped nickel specimen had coarse-grained structure
and contained a large number of annealing twins borderedby coherent R3 boundaries, typically for nickel with middle
level of stacking fault energy (128 9 10-3 J/mm2 [34]).
As shown in Fig. 2b, the grain size distribution of thesulfur-doped nickel specimen ranged from 20 to 320 lm.
An average grain size of the specimen was 84.4 lm so thatsome 35 grains might be connected in the specimen
thickness.
The surface orientation of tested nickel specimens
appeared to orient toward {001} and {101} (Fig. 2c). As
indicated by the attached table of GBCD, the total fractionof low-R CSL boundaries in this specimen was of 46 %,123

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Fig. 2 a Grain boundary microstructure (GBM) map, b grain size
distribution, c inverse pole figure with grain boundary character
distribution (GBCD) data of pre-deformed sulfur-doped nickel
specimenand R3 boundaries occupied a high fraction of 31 % of thetotal fraction. The occurrence of a high fraction of R3
boundaries may be ascribed to the dominant evolution of
annealing twins. In the present investigation, these {111}coherent annealing twin boundaries are hereafter classified
as R3 boundaries.
In situ observations of crack propagation processes
in sulfur-doped polycrystalline nickel during cyclicloading
Figure 3a shows one of in situ observations on crack
propagation processes in the sulfur-doped nickel specimen(Specimen-1) during cyclic loading. In this figure, theloading axis is horizontal direction. The character of individual grain boundaries surrounding the coarse grain in thecenter (with the grain diameter of almost 500 lm) is indicated by different symbols in the micrographs. Low-R CSL

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and random boundaries are denoted by R plus numeral andR, respectively. Low-angle boundaries are indicated as R1
CSL boundaries. The pink lines show slip lines induced bycyclic loading. Figure 3b shows the observed change inlocal crack propagation rate for the Specimen-1 with thegrain boundary microstructure already shown in Fig. 3a.
Here, let us examine the observed individual crack paths
indicated by different kinds of lines, such as solid line anddashed line, for the crack propagation processes in thisspecimen.
The pre-crack stayed at the triple junction T1 where thethree random boundaries interconnected (Fig. 3a/(i)). The
crack advanced along two random boundaries from thetriple junction T1 toward the triple junction T2 on the righthand side during the numbers of cyclic loads N reach to2000 cycles (Fig. 3a/(ii)). The average crack propagation
rate in this period was 6.2 9 10-3 lm/cycle, as shown in
Fig. 3b. The crack did not propagate toward the triple
junction T3 composed of R1, R3, and random boundaries
on the left-hand side of the triple junction T1. From ourprevious study on the triple junction hardening in intrinsically brittle molybdenum [35, 36], it was found that the
triple junctions composed of more than two low-R CSL
boundaries show higher fracture resistance. The stressconcentration associated with dislocation pile-ups is suggested to be higher around the triple junctions with higher
connectivity of low-R CSL boundaries, because of themore difficult accommodation of lattice dislocations atthese low-energy CSL boundaries [37, 38]. Nevertheless,the crack tends to propagate selectively along the path witha higher connectivity of ‘‘intrinsically weak high-energy’’
random boundaries [32] toward the triple junction T2,
because the triple junction T3 inhibited the crack propagation. The exactly same finding was already obtained
from early in situ observations on crack propagation in Bidoped coarse-grained copper [39].
As shown in Fig. 3a/(iii), the crack kept remaining at the
triple junction T2 where R1 and random boundaries lying in
the crack propagation direction, for further more than 1000
cycles up to N = 3000. Then slip lines were newly generated in the grains locating in front of the advancing cracktip, when the number of cyclic loading reached 5000 cycles(Fig. 3a/(iv)). Further extension of these slip lines wasinhibited by the random boundary, R1, which was almostaligned in the horizontal direction. The numbers of the slip
lines tend to increase with increasing the number of cyclicloading (Fig. 3a/(iv–vi)). The crack advanced slowly along
the random boundary R1, resulting from the stress concentration caused by the interaction between lattice dislocations within slip lines and the random boundary R1. The
crack propagation was arrested by the triple junction T4
where R1 boundary laid in the crack propagation direction(Fig. 3a/(vii)). As shown in Fig. 3b, the crack propagation

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Fig. 3 a A bird’s-eye view ofintergranular crack propagationin the grain boundary network
associated with heterogeneous
grain structure, and b change inthe local crack propagation rate
in sulfur-doped nickel specimen
A during cyclic bending test

was accelerated along the single random boundary (R1) on
the right-hand side of triple junction T2 by development ofslip lines, and then the propagation rate of crack gradually
lowered, as it approached the triple junction T4.

On the other hand, immediately after the stoppage ofcrack propagation occurred at the triple junction T4, anothercrack propagation path occurred along the R1 boundaries on
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Fig. 4 a A bird’s-eye view ofcrack propagation in the grainboundary network associatedwith homogeneous grain
structure, and b change in thelocal crack propagation rate insulfur-doped nickel specimen A
during cyclic bending test. Ft isthe propagation rate oftransgranular crack

the branching of crack occurred at the triple junction T2 after
N = 9000, as seen from Fig. 3a/(viii). The slip bands almost
parallel to the path of cracked R1 boundary were observed in
the adjacent grains. The more difficult cracking of the R1

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boundary than random boundary probably resulted from theheavy stress concentration produced by extensive interaction
between the R1 boundary and localized crystal slip during
cyclic deformation (Fig. 3a/(viii)). Nevertheless, the crack

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propagation was completely stopped later on the R1boundary.
As for further stage of crack propagation, another crack
path appeared around the triple junction T5, on the lefthand area, after the stoppage of crack propagation at the R1boundary (Fig. 3a/(ix)). It is notable that the nucleation andpropagation of crack were never observed on the specimen
surface with high connectivity of the low-R boundaries (inthe area (I)). On the other hand, the main crack immediately and continuingly advanced along random boundaries
in the boundary network with high connectivity of randomboundaries (Fig. 3a/(viii–x)). It is confirmed that intergranular crack predominantly and selectively tends topropagate along random boundary network, and neverpropagates along low-R CSL boundaries connecting withthis active crack propagation path, resulting in the fatigue
crack propagation rate of more than 1.9 lm/cycles which ismuch higher than that obtained from the area mentioned
before, as seen from Fig. 3b.
Let us examine the case of another specimen, similarly
tested under cyclic loading. Figure 4a demonstrates thein situ observations on the crack propagation processes
obtained from another sulfur-doped nickel specimen(Specimen-2) subjected to cyclic loading. Figure 4b shows
the observed change in the local crack propagation rate atdifferent steps of crack propagation in relation to the local
nature of grain boundary microstructure for the Specimen2, shown in Fig. 4a.
The pre-crack (precisely termed ‘‘premature-crack’’)stayed at the triple junction T1 where R3 and randomboundaries (R1) interconnected in front of the crack tip(Fig. 4a/(ii)). The crack advanced along the R3 boundary
locating along the extended direction of the propagation ofpre-crack (Fig. 4a/(ii)). The crack propagation rate along
this R3 boundary was 1.80 9 10-3 lm/cycle as shown in
Fig. 4b. As soon as the crack propagation was stopped on
the R3 boundary, another crack propagation path immediately becomes active along the random boundary (R1)connecting the triple junction T1 at the left-hand side. As aresult, the branching of crack occurred at the triple junction
T1 (Fig. 4a/(iii)). The crack propagated along the randomboundary (R1) then stopped at the triple junction T2.
As observed in the case of the Specimen 1, those slip
lines which appeared parallel to the surface trace of the R1boundary formed in the grain associated with the triple
junction T3. This R1 boundary cracked accompanying withthe formation of a large number of slip lines within nearby
grain (Fig. 4a/(iii–iv)). Moreover, the crack propagation
along R1 boundary was immediately directed into the graininterior of one of adjoining grains, resulting in a change infracture mode, i.e., from intergranular to transgranularfracture, and then finally stopped within the grain interior(Fig. 4a/(v)).

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The local propagation rate was Fi = 5.1 9 10-2 lm/
cycle for the intergranular crack along the R1 boundary, andFt = 6.9 9 10-3 lm/cycle for the transgranular crack in
the grain interior, as shown in Fig. 4b. The first and second
crack propagation paths joined together by further cyclicloading (Fig. 4a/(vii)). The third crack propagation pathalong random boundary (R2) began to operate, after thesecond crack propagation path became inactive in the graininterior (Fig. 4a/(vi)). The crack immediately advanced
along random boundary network in which ten randomboundaries were continuously connected to each other (R2–
R11) with no interruption. The intergranular crack propagation rate was accelerated from 5.6 9 10-2 lm/cycle to
0.7 lm/cycle. These results have evidenced that thebranching of random boundary connectivity induced by theinteraction with low-R CSL boundaries is useful and powerful for controlling the crack propagation in sulfur-doped
polycrystalline nickel, as previously achieved by the present
authors, on the similar principle in order to improve thesuperplasticity of intrinsically brittle Al–Li alloy throughenhanced interaction with newly introduced low-angle (R1)boundaries by strain rate change during high temperature
deformation [40].
Effect of grain boundary microstructure on fracture
resistance in sulfur-doped nickel during static bending
tests
The present investigation has revealed that the crack
propagation process of intergranular brittle fracture in sulfur-doped nickel was predominantly affected by grainboundary microstructure. In particular, the lower connectivity of random boundaries induced the more frequent
crack branching and lower crack propagation rate. Thissection is concerned with SEM observations of fracturesurfaces of sulfur-doped nickel specimens with different
grain boundary microstructures, in order to confirm theeffect of grain boundary connectivity on the fracture mode
operating in polycrystalline nickel embrittled by sulfur. The
utility of GBE for controlling sulfur segregation-induced
embrittlement in polycrystalline nickel is discussed.
(1) Figure 5a and b shows the stress–strain curves
obtained by static four-point bending test, taken from therelated work recently reported by the present authors [26].
The possible crack propagation processes are schematicallyshown in Fig. 5c and d with SEM micrographs (Fig. 5e, f)taken from typical fracture surfaces in the fractured specimens with different grain boundary microstructures. First,
it should be mentioned briefly how to produce different
grain boundary microstructures in nickel specimens.
Type A specimens were prepared by cold rolling to
80 % reduction ratio and subsequent annealing at 923 K
for 1.2 ks. These annealed specimens were subjected to123Author's personal copy
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Fig. 5 SEM micrographs oftypical fracture surfaces
produced by static bending test
and schematic illustrations ofpossible crack propagation
processes: a and b stress strain
curves for the Type A, and TypeB specimens, with different
grain boundary microstructuresin sulfur-doped nickel
specimens during cyclicloading, c and d schematic ofcrack propagation processes,
e and f SEM micrographs offracture surfaces, respectively

further cold rolling to 5 % reduction ratio and subsequentannealing at 1173 K for 3.6 ks. On the other hand, Type B
specimens were obtained by cold rolling to 80 % reduction
ratio and subsequent annealing at 923 K for 3.6 ks. These
specimens were subjected to further cold rolling to 5 %reduction ratio and subsequent annealing at 1223 K for
1.8 ks. Both Type A and Type B specimens were subjected
to sulfur-doping heat treatment. These specimens wereenclosed with sulfur powders (0.04 g) in a Pyrex glass tube

123with 8 mm internal diameter and 70 mm length in a vacuum of 3 9 10-3 Pa and heated at 873 K for 72 h. The
average value of sulfur content in these specimens was
290 ppm slightly different between the two types: 270 ppm
in Type A specimen, and 310 ppm in Type B specimen.
According to Doherty et al. [3], sulfur-induced embrittlement in polycrystalline nickel takes place around room
temperature when the sulfur content is increased up to
140 ppm. Moreover, Loier and Boos [4] have reported that

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embrittlement of polycrystalline occurred when the bulk
sulfur concentration ranges from 50 to 200 ppm. Therefore,
the sulfur content of both Type A and Type B specimens
was considered as being enough for occurrence of sulfurinduced embrittlement in nickel.
Type A specimens had a homogeneous fine-grained
structure, a higher fraction of low-R CSL boundaries(FR = 53 %) and a lower fractal dimension of randomboundary network with maximum connectivity (maximum
random boundary connectivity (MRBC, DR = 1.23) in thespecimen. On the other hand, Type B specimens had aheterogeneous grain structure, a lower fraction of low-R
CSL boundaries (FR = 40 %) and a higher fractal dimension of MRBC, DR = 1.52. The higher fractal dimensionof MRBC can indicate the higher random boundary
connectivity.
(2) The stress–strain curves obtained by tensile fracture
tests [26] for five different specimens received the sameheat treatment are shown in Fig. 5a and b on the top. The
arrow heads indicate the position at which a crack seems tonucleate at the bottom of the notch, because the stressirregularity arose on the stress–strain curves. Most of theType A specimens did not fracture in the range of small
bending strain of e = 0.14, while all of the Type B specimens fractured, when five bending tests were carried out
for each type of the studied specimens [26]. The Type A
specimens were subjected to excess bending strain until
they fractured in order to obtain SEM micrographs offracture surface.
The crack propagation processes observed in the Type A
and Type B specimens are schematically shown Fig. 5cand d. The intergranular crack propagation is inhibited by
the frequent interaction with fracture-resistant low-energy/
low-R CSL boundary in the Type A specimens which had ahigher fraction of low-R CSL boundaries and a lower
fractal dimension of MRBC. As a result, the Type A
specimen tends to show a higher fracture resistance. On theother hand, in the case of Type B specimen, the main crack
propagates at a higher crack propagation rate without crack
branching, due to a lower fraction of low-R CSL boundaries and a higher fractal dimension of MRBC, leading toless hindrance of the crack propagation originating fromthe interaction with low-R CSL boundaries.
From SEM micrographs of the fracture surface shown in
Fig. 5e and f on the bottom, it is evident that the Type A
specimen shows a typical ductile fracture surface accompanying with extensive formation of dimples, while theType B specimen shows a typical intergranular brittle
fracture surface. Accordingly, it is concluded that thefracture mode can drastically change from intergranularbrittle to transgranular ductile fracture, resulting in an
improvement in the fracture toughness, by increasing thefraction of low-R CSL boundaries and decreasing the4015

fractal dimension of MRBC (random boundary connectivity) recently proposed by the present authors [26]. Insummary, sulfur-induced intergranular embrittlement ofpolycrystalline nickel is controlled by the strategy of GBE,
based on the introduction of the lower level of randomboundary connectivity realized by the introduction of ahigher fraction of low-R CSL boundaries.
Grain boundary engineering for control of segregationinduced embrittlement
The basic concept of the GBE for controlling segregationinduced embrittlement in polycrystalline material systems
is discussed.
The present work based on in situ observations on crack
propagation processes has revealed the important role ofgrain boundary microstructure in fracturing of polycrystalline nickel embrittled by sulfur. Figure 6 schematicallyshows the crack propagation processes observed in sulfurdoped nickel specimen during cyclic loading. Randomboundaries were proved to be the most preferential path ofcrack propagation. The higher connectivity of randomboundaries was found to bring about the higher crack
propagation rate in grain boundary network in embrittled

Fig. 6 Schematic illustration of the crack propagation processes inpolycrystalline metallic materials with different grain boundary
microstructures, leading to the grain boundary engineering for control
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polycrystalline nickel. The crack branching occurs at the
partly cracked triple junction when the propagating crack
along random boundaries is arrested by segregation andfracture-resistant low-R CSL boundaries with low energy.
When the main propagating crack is arrested, the crack
propagation rate tends to locally slow down, due to crack
branching resulting in a kind of stress relaxation. The lowR boundaries likely fracture only when the grain boundary
plane is aligned parallels to the slip direction in adjacentgrains. It is concluded that the crack propagation along
low-R boundaries tends to proceed slowly, resulting in a
low crack propagation rate. Accordingly, it is reasonable toexpect that the introduction of a higher fraction of low-R
boundaries and of the lower connectivity of randomboundaries can effectively control the dominant occurrenceof intergranular fracture along high-energy randomboundaries in polycrystalline materials embrittled by detrimental elements.

Conclusions
In situ observations of crack propagation in sulfur-doped
polycrystalline nickel specimens were performed in order toclarify possible roles of grain boundary microstructure incrack propagation processes in polycrystalline nickel
embrittled by grain boundary segregation of sulfur. The
importance of the grain boundary microstructure for controlling segregation-induced intergranular fracture in polycrystalline nickel is discussed based on in situ observations.
The main results and conclusions are the following:
(1)

(2)

(3)

(4)

The main crack propagates predominantly along
random boundaries. The crack propagation rate canbe locally accelerated in the region of grain boundary
network where random boundaries are more frequently interconnected to each other.
The crack propagation tends to be arrested when the
propagating crack reached the triple junction composed of more low-R CSL boundaries in the grainboundary network.
The branching of the main propagating crack occurs
at the triple junction partly cracked, when the crack
propagation is arrested by connecting low-R CSL
boundaries. The crack propagation rate can be locallyaffected by occurrence of crack branching, dependingon the character of triple junction.
The low-R boundaries are generally difficult to
fracture, but can fracture only when the trace of thegrain boundary is aligned parallel to the slip directionin the adjoining grains. The crack propagation along
low-R boundaries results in lowering the crack
propagation rate or the arrest of crack propagation.

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(5)

The fracture mode can drastically change fromintergranular brittle to transgranular ductile fracture
or vice versa, depending on the fraction of low-R CSL
boundaries and the connectivity of random boundaries. It has been revealed that sulfur-induced embrittlement of polycrystalline nickel can be effectively
controlled by GBE through lowering the connectivityof random boundaries by introducing the high fractionof low-R CSL boundaries.

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