Localized strain and heat generation during plastic deformationin nanocrystalline Ni and NiFe

T. Chan Y. Zhou I. Brooks G. Palumbo

U. Erb

Received: 29 November 2013 / Accepted: 31 January 2014 / Published online: 15 February 2014

Springer Science+Business Media New York 2014

Abstract Room temperature tensile testing was performed

on a coarse-grained polycrystalline Ni (32 lm), a nanocrys-talline Ni (23 nm) and two nanocrystalline NiFe (16 nm)

electrodeposits at two strain rates of 10-1 and 10-2/s. Strain

localizations and local temperature increases were simulta-

neously recorded during tensile testing. For all materials,

higher loads or higher strain rate generally resulted in higher

peak temperature with the highest temperatures recorded in

the fracture regions. The maximum temperature for the

nanocrystalline materials was just over 80 C, which is sig-nificantly below the reported temperatures for the onset of

thermally activated grain growth. Therefore, the previously

reported grain growth observed on similar materials after

tensile deformation is likely not thermally activated but a

stress-induced phenomenon. Despite the wide grain range

from 16 nm to 32 lm, all samples exhibited similar strainlocalization behavior. Local strain variations initiated in the

early stage of macroscopic uniform deformation, subsequent

necking and fracture took place in the region of initial strain

localization. While the coarse-grained polycrystalline Ni

exhibited little strain rate sensitivity, gradually increased

strain rate sensitivity was observed for the 23 nm Ni and the

two 16 nm NiFe samples, suggesting that both dislocation-

mediated and grain-boundary-controlled mechanisms were

operative in the deformation of the nanocrystalline Ni and

NiFe samples.

Introduction

Many past papers have reviewed the effect of grain size on

the mechanical properties of nanocrystalline metals (e.g.,

[13]). As per the HallPetch grain size hardening mech-

anism, nanometals exhibited significant increases in yield

strength and hardness (e.g., [46]). However, limited ten-

sile ductility was usually observed for nanocrystalline

materials in comparison with their conventional coarse-

grained counterparts due to their low capacity for strain

hardening [2, 7]. Low tensile ductility of nanomaterials is

almost invariably associated with early onset of localized

deformation, e.g., necking and/or shear bands [811].

While a great deal of effort has focused on the underlying

deformation mechanisms (e.g., [1217]), there is limited

experimental investigation on the phenomena of localized

strain and heat generation during plastic deformation of

these materials, which could have a significant effect on the

microstructure and mechanical properties [18].

There is considerable debate over the possibility of

thermally activated grain growth due to heat generation

associated with localized strain of specimens during defor-

mation. Previous studies have reported grain growth during

tensile testing in nanocrystalline Ni [10], NiFe [18] and

CoP [19], all produced by Integrans electrodeposition

process [2022]. For example, nanocrystalline Ni (grain

size: 20 nm) was tested in tension at strain rates between

10-5 and 10-1/s [10]. Transmission electron microscopy in

one of the shear band regions at the higher strain rates

revealed considerable grain growth. It was speculated that

grain growth in nanocrystalline Ni was due to the kinetic

energy released during high-speed deformation, which

could result in local hot spots within the shear band reaching

temperatures on the order of 300 C [10]. At this tempera-ture, significant grain growth is expected to take place in the

T. Chan Y. Zhou (&) U. ErbDepartment of Materials Science and Engineering,

University of Toronto, 184 College Street, Suite 177,

Toronto, ON M5S 3E4, Canada

e-mail: yijian.zhou@utoronto.ca

I. Brooks G. PalumboIntegran Technologies Inc, 6300 Northam Drive,

Mississauga, ON L4V 1H7, Canada

123

J Mater Sci (2014) 49:38473859

DOI 10.1007/s10853-014-8099-1

20 nm Ni electrodeposits. Calorimetric studies have shown

that peak growth temperatures for major grain growth are in

the range from 280 to 290 C for the Ni electrodeposits withaverage as-deposited grain size between 15 and 30 nm [23].

Moreover, grain growth activity was observed in a sequence

of relaxation and subgrain coalescence processes when

20 nm Ni electrodeposits were subjected to annealing at

lower temperatures, as low as 120 C [24]. Higher thermalstability was observed in nanocrystalline NiFe. For

example, the Ni-10 wt% Fe (grain size 15 nm) and

Ni-20 wt% Fe (grain size 13 nm) displayed peak tempera-

tures of grain growth around 334 and 379 C, respectively,indicating a stabilizing effect by alloying nanocrystalline Ni

electrodeposits with Fe [23]. These experimental observa-

tions provide a basis to assess the possibility of having

potential grain growth resulting from localized heating in

nanocrystalline nickel-based electrodeposits during room

temperature deformation.

The main focus of the current study was (1) to establish

a correlation between localized strain and heating, and (2)

to assess the possibility of having thermally activated grain

growth in this type of nanomaterials. Heat dissipation and

strain localization during in situ tensile testing at different

strain rates were investigated on nanocrystalline Ni and

NiFe electrodeposits produced through the same produc-

tion facility and using the same synthesis process as those

for the samples used in Refs. [10, 18, 19].

Experimental

The materials used in this study consist of four batches of

samples, coarse-grained polycrystalline nickel (commercially

available Ni 200), and three nanocrystalline electrodeposits,

Ni and two NiFe alloys with different Fe contents. All

nanocrystalline electrodeposits were prepared by Integran

Technologies Inc, which used the same process to produce

the samples for the investigations in Refs. [10, 18, 19].

Samples from the same nanocrystalline batches were already

characterized in previous studies [25, 26], using transmission

electron microscopy and X-ray line broadening measure-

ments (Scherrer approach) [27], as well as energy-dispersive

X-ray spectroscopy (EDS). For the polycrystalline Ni, the

grain size was determined using standard metallography.

Tensile coupons of 1.0 mm thickness were machined by

the electric discharge machining (EDM) method to produce

dog-bone-shape tensile testing samples with dimension

details as shown in Fig. 1. There was a minor modification in

the geometry of tensile coupons in the current study with

respect to ASTM E8 standard. The coupon tab width was

widened from 10 to 20 mm (Fig. 1) to provide increased

gripping area after slipping problems owing to the extremely

high material strength were experienced, while the rest,

especially the critical gauge section, was kept identical to that

of ASTM E8.

All materials were tested at tensile strain rates of 10-2

and 10-1/s on a MTS servo-hydraulic testing machine

combined with the in situ measurements of local strain and

temperature across the gauge length, using the digital

image correlation (DIC) technique and an infrared camera,

respectively. Prior to tensile testing, each sample was

prepared by spraying a layer of black paint on one side to

equalize emissivity for better infrared detection, thus more

accurate temperature measurement. On the opposite side,

samples were lightly abraded using 320 silicon carbide

paper for the measurement of localized strains. Localized

strain patterns during deformation were captured by a high

resolution DIC Camera (Allied Vision Technologies),

while the global strain was obtained based on the crosshead

displacement of the MTS machine frame. Samples with

abraded surfaces facing toward the CCD camera were

illuminated by a source of white light during the entire

tensile testing. The images were post-processed through a

DIC system to calculate the displacement and the localized

strain distributions. Localized strains contours were then

mapped along the sample gauge over a length of 40 mm.

Changes of temperature in the samples were measured

using a pre-calibrated high resolution infrared camera

(Deltatherm 1410) during tensile deformation. Frame rates

of the infrared detector were set to 15 frames per second to

best capture the temperature profiles along sample gauges.

Images captured during the experiment were post-pro-

cessed through a Thermoelastic Stress Analysis (TSA)

system to locate the temperature increase along the gauge

section during tensile testing. The maximum temperature

along the sample gauges was recorded for all samples.

Results and discussion

Grain size and chemical composition

Table 1 summarizes the chemical compositions and grain

sizes for the samples used in the current study. For the

coarse-grained Ni, i.e., commercial Ni 200, the average

grain size of 32 lm was obtained after analysis of a seriesof optical micrographs and measurement of over 400

grains. For the nanocrystalline materials, the nano-Ni had a

Fig. 1 Tensile coupon dimensions in mm

3848 J Mater Sci (2014) 49:38473859

123

grain size of 23 nm, while the two NiFe samples both had

a grain size of 16 nm as per XRD line broadening mea-

surements [25, 26].

Deformation at different tensile strain rates

All materials were tested at strain rates of 10-1 and 10-2/s.

Figure 2 shows the respective engineering stressstrain

curves for all samples.

At the rate of 10-2/s, the coarse-grained Ni (grain size:

32 lm) showed a yield strength (YS) of 208 MPa, ultimatetensile strength (UTS) of 442 MPa and considerable duc-

tility with total elongation (or total global strain) of 50.4 %.

With the refinement of grain sizes, there is a significant

increase in YS and UTS due to the HallPetch grain size

strengthening mechanism as indicated in Fig. 2a. The

nanocrystalline Ni (grain size: 23 nm) showed values of

YS (950 MPa) and UTS (1504 MPa), about 4 times higher

than the coarse-grained Ni. The two remaining 16 nm

NiFe samples exhibited different strength values, YS of

1010 and 1130 MPa, as well as UTS of 1579 and 1741 MPa,

for the Ni-2.6 wt% Fe and Ni-8.5 wt% Fe, respectively. The

fact that higher strength was observed at higher Fe content

suggests that there is some solution-strengthening effect due

to Fe solute for the NiFe samples. In addition, all

nanocrystalline samples showed limited ductility with

measured elongation to fracture not exceeding about 8 %.

The same trends in the tensile properties were found at

the strain rate of 10-1/s. Significant strength enhancements

were attained, whereas the total elongation or ductility was

limited for the nanocrystalline samples in comparison with

the coarse-grained Ni (Fig. 2b). These observations of high

strength and reduced ductility are in agreement with

numerous previous experimental studies summarized in

previous review papers (e.g., [2, 28]). While the HallPetch

relationship accounts for the strengthening, low tensile

ductility is generally attributed to the lack of strain hard-

ening required to resist macroscopic strain localization and

extend uniform plastic flow for nanocrystalline materials

(e.g., [9, 11]).

Table 2 summarizes tensile property values of all sam-

ples at the two strain rates, 10-2 and 10-1/s. For the coarse-

grained Ni (32 lm), strain rate sensitivity is insignificant;the key tensile properties, yield strength, ultimate tensile

strength and elongation remain virtually unchanged. When

varying the strain rate from 10-2 to 10-1/s, the nanocrys-

talline Ni (23 nm) showed YS increase from 950 to

970 MPa but essentially no change in UTS, indicating

some minor strain rate sensitivity. In comparison, the two

NiFe samples with smaller grain size of 16 nm displayed

evident sensitivity to the strain rate, showing an increase in

both YS and UTS at the higher strain rate.

Although a detailed analysis of strain rate sensitivity

was not the major focus of the current study, it should be

pointed out that the observed grain size dependent strain

rate sensitivity is in general agreement with previous

studies on nickel-based alloys (e.g., [28]). In a recent study

[29], the strain rate sensitivity, m, was examined over a

wide grain size range from 10 nm to 100 lm for Ni andNi-based alloys based upon extensive results from various

Table 1 Sample grain size and composition

Sample Fe content

(wt%)

Average

grain size

1. Coarse-grained Ni (Ni 200) 0 32 lm

2. Nanocrystalline Ni 0 23 nm

3. Nanocrystalline NiFe 1 2.6 16 nm

4. Nanocrystalline NiFe 2 8.5 16 nm

Fig. 2 Stress-strain curves of all samples deformed at a strain rate of a 10-2/s and b 10-1/s

J Mater Sci (2014) 49:38473859 3849

123

studies with different testing methods, e.g., tensile tests at

different stain rates, stress relaxation [30] and strain rate

jump tests [31]. It was observed that the m values increased

significantly and ranged from 0.011 to 0.033 for materials

with average grain sizes between 10 and 50 nm, while in

the broad 50 nm100 lm range, the m value showed lim-ited variation and remained below 0.005 [29] and within

the reported 0.0010.004 range for coarse-grained Ni [10].

This explains the insensitivity toward the strain rate for the

coarse-grained Ni (32 lm), and gradually increaseddependence on the strain rate for the 23 nm Ni and the two

16 nm NiFe alloys (Table 2).

The significant m value increase at very small grain

sizes for Ni-based materials can be attributed to deforma-

tion mechanisms other than the well-established disloca-

tion-mediated (DM) mechanisms for coarse-grained

materials, owing to the close association between strain

rate sensitivity and underlying deformation mechanisms

(e.g., [3236]). Numerous experimental and simulation

studies have provided convincing evidence that dislocation

activity is negligible and deformation is dominated by

grain-boundary-controlled (GBC) mechanisms, such as

grain boundary sliding or grain rotation for nanocrystalline

materials with average grain sizes less than 10 nm (e.g.,

[28, 37, 38]). Therefore, it is plausible that in the current

study, DM mechanisms dominated in the deformation of

the coarse-grained Ni, whereas the 23-nm Ni and the two

16-nm NiFe samples deformed through both DM and

GBC mechanisms.

Strain localization

For conventional metallic materials, plastic deformation

during uni-axial tensile testing is normally expected to be

distributed uniformly at applied stresses lower than the

UTS; macroscopic localized deformation, necking, occurs

at the stress beyond the UTS [39]. In reality, the amount of

plastic deformation can vary from one location to another

well before the onset of necking. Digital image correlation

(DIC) strain measurement [40] has been shown to be an

effective approach to examine in situ strain localization of

materials, i.e., it is capable of capturing the evolution of

localized strain during tensile testing. For example in an

earlier study, rapid evolution of shear banding was cap-

tured using the DIC approach for fully dense electrode-

posited nanocrystalline Ni (grain size: 20 nm) during an

in situ tensile testing at quasi-static strain rate of 10-4/s

[41].

Figure 3 shows the results of DIC measurements at

different stages of deformation represented by five global

strains for the coarse-grained polycrystalline Ni of 32 lmduring tensile testing at a strain rate of 10-1/s. In the

contour map (Fig. 3a), the levels of local strain from low to

high are indicated by colors in a sequence of purple, blue,

green, yellow and red and the local strain field contours of

the coarse-grained Ni are mapped along the gauge section

at selected global strains indicated at the bottom of Fig. 3a,

up to the fracture strain of just over 50 %. Although uni-

form deformation is expected for the sample within the

global strain up to UTS at 45 % (Table 2), noticeable

variations of strain are displayed within the gauge length

even at low global strain of 15 % as indicated by two

colors, purple and blue.

As the global strain increases, the colors in the gauge

section keep changing due to intensified localized defor-

mation. At the global strain of 35 %, still well within the

range of uniform deformation, the gauge section already

shows a range of different contour colors, indicating sig-

nificant levels of localized strain. The highest level of

localization occurred during necking when imposed

deformation was concentrated essentially only in the

necking region. This is demonstrated in the contour map at

the fracture strain of 50 %, whereby the red area corre-

sponds to the region of necking with very high local strain,

leading to fracture at the end of the tensile test (Fig. 3a).

Corresponding strain distributions along the gauge

length at the same global strains, i.e., 15, 25, 35, 45 and

50 %, respectively, are shown in Fig. 3b to present more

detailed information on the evolution of strain localization.

Note that the global strains do not always match with local

Table 2 Tensile properties at strain rate of 10-2 and 10-1/s

Samples Grain size Strain

rate (/s)

Yield strength

(MPa)

UTS

(MPa)

Elongation

at UTS (%)

Total

elongation (%)

Poly-Ni 32 lm 10-2 208 442 45 50.4

10-1 208 445 45 51

Nano-Ni 23 nm 10-2 950 1504 5.5 8.2

10-1 970 1508 5.5 8

Nano-Ni-2.6 % Fe 16 nm 10-2 1010 1579 5.5 7.6

10-1 1100 1763 5.5 7.5

Nano-Ni-8.5 % Fe 16 nm 10-2 1130 1741 5.5 7.7

10-1 1150 1867 5.5 7.5

3850 J Mater Sci (2014) 49:38473859

123

strain values, especially at low stress/strain levels. This

difference is somehow expected because of different

methods used in the strain measurements; global strain was

obtained indirectly from the crosshead displacement of the

tensile machine frame, while local strain values were

measured directly in the sample gauge section through the

DIC technique. It has been shown that strain values based

on crosshead displacement include contributions from

other sources, e.g., machine compliance and materials

outside the gauge length (e.g., those close to grip sections),

in addition to the strain of the sample gauge length [42]. As

a result, the measured strain based on the crosshead dis-

placement almost always exceeds the actual strain in the

sample gauge section [42]. However, a discussion of

crosshead strain vs DIC strain is beyond the scope of this

study. The focus here is instead on the strain localization

within the stage of macroscopic uniform elongation, i.e.,

up to the global strain at UTS, a characteristic deformation

feature that can be reliably identified using crosshead

displacement. Hence, the disparity between global and

local strain has little impact in addressing strain localiza-

tion for the samples in the current study.

Examination of the strain distribution curves reveals that

there are two deformation regions within the gauge section:

the region at the two ends, Region 1 and the region in the

center, Region 2. Due to its proximity to the grip section,

Region 1 is strongly influenced by its geometric confine-

ment and has to accommodate the transition from close to 0

strain at the grip section to the imposed strain on the

sample. In other words, two factors contribute to the local

strain in Region 1, geometric conformation and material

properties. It appears that the geometric contribution

dominates the local strain in Region 1 as demonstrated by

the largely linear relationship between gauge location and

local strain with a slope increasing with the respective

imposed strain or global strain (Fig. 3b). This linear rela-

tionship of Region 1 ends where the strain transition to the

imposed value is accomplished; the end locations depend

Poly Ni (0.1/s)

0

20

40

60

80

0 10 20 30 40Gauge length, mm

Loca

l str

ain,

%

50%45%35%25%15%

(a)

(b)

Region 2Region 1 Region 1

Fig. 3 Strain distribution at different global strains for the coarse-grained Ni (grain size: 32 lm), a strain field contour maps, andb strain distributions within the gauge length at five levels of global

strain. Region 2 is approximately located between the two red-dotted

lines (Color figure online)

J Mater Sci (2014) 49:38473859 3851

123

on the global strain and manifest themselves in the area

with evident slope change. For example at the UTS global

strain of 45 %, the locations are approximately at the gauge

locations of 10 and 33 mm, respectively, as shown in

Fig. 3b. Region 2 is located between these two points,

where the geometric contributions diminish and local

strains result entirely from the material properties. Hence,

the analysis of strain localization evolution should focus on

the strain distributions in Region 2, identified approxi-

mately between the two dotted lines in Fig. 3b.

For the coarse-grained Ni, strain distribution appears

fairly uniform in most locations of Region 2 at the global

strain of 15 %. Slightly different local strain was displayed

between the gauge locations of 26 and 32 mm (Fig. 3b).

This strain disparity became more visible in the same

location at the global strain of 25 %, and gradually evolved

into a broader valley, with a local strain significantly dif-

ferent from the rest in Region 2 at 35 % global strain, even

though the sample was still well within the strain range of

macroscopic uniform deformation (\45 %). When theapplied stress was above UTS, necking took place and

subsequent deformation concentrated in this area as dem-

onstrated in the local strain distribution at the global strain

of 50 % (Fig. 3b). Immediately thereafter, fracture occur-

red at the location with peak strain of about 80 %, sub-

stantially higher than the global fracture strain just over

50 % (Fig. 3a, b).

A similar development of localized strain is observed in

the nanocrystalline Ni sample with 23 nm grain size. As

shown in the strain field contour maps of Fig. 4a, deforma-

tion appears to be fairly uniform across the gauge length in

the early stage up to a global strain of 2.5 %. The sample

exhibited notable differences in local strains at a global strain

of 3.5 % as indicated by two colors, dark and light blue. The

local strain became more pronounced at the UTS strain of

5.5 % (Fig. 4a). Upon necking, imposed deformation was

concentrated in the center, and at about 8.0 % global strain,

fracture occurred in the necking region (red area) at an angle

Nano Ni (0.1/s)

0

3

6

9

12

15

0 5 10 15 20 25 30 35 40

Gauge length, mm

Loca

l str

ain,

%

7.5%5.5%4.5%3.5%2.5%

(b)

(a)

Region 2Region 1

Region 1

Fig. 4 Strain distribution at different global strains for the nanocrystalline Ni (grain size: 23 nm), a strain field contour maps, and b straindistributions within the gauge length. Region 2 is approximately located between the two red-dotted lines (Color figure online)

3852 J Mater Sci (2014) 49:38473859

123

of 62 (dashed line) as shown in Fig. 4a. In comparison withthe coarse-grained Ni, necking on the nanocrystalline Ni

sample is less pronounced. Details on strain localization

evolution can be more clearly demonstrated from strain

distributions along the gauge length at given global strains.

Figure 4b shows the distributions at five global engineering

strain levels of 2.5, 3.5, 4.5, 5.5 and 7.5 % for the nano-

crystalline Ni sample tested at a strain rate of 10-1/s.

Similar to the deformation of coarse-grained Ni, the

strain distribution in Fig. 4b can be divided into two regions

for the nanocrystalline Ni: Region 1 at the ends of the gauge

length with strong influence from sample geometric con-

finement and Region 2 (identified approximately between

the two red-dotted lines) representing intrinsic material

properties. It can be seen that the nanocrystalline Ni already

exhibited slight local strain variation somewhere between

the gauge length of 17 and 25 mm in Region 2 even at the

global strain as low as 2.5 %. As the global strain increased

from 3.5 to 5.5 %, i.e., still within the strain range of

macroscopic uniform deformation, the strain disparity

became more and more pronounced between this area and

the rest in Region 2. In the final deformation stage of

instability, the imposed deformation concentrated in this

area in the form of necking with the peak local strain

reaching about 14 % at the global strain of 7.5 %. At the

slightly higher global strain of 8.0 %, the sample fractured

at the location with the peak local strain (Table 2; Fig. 4b).

For the two 16 nm NiFe electrodeposits and all mate-

rials tested at a strain rate of 10-2/s, strain localization was

observed in a similar manner. Generally, all samples

showed not much difference in the development of strain

localization during tensile testing at both strain rates.

Despite the wide grain size range (16 nm to 32 lm) andvery different tensile strengths and measured elongation to

fracture values, strain localization invariably started during

the stage of macroscopic uniform deformation, much

earlier than the occurrence of the macroscopic strain

localization in all of them. Subsequent necking took place

close to the region where the strain localization initiated.

There is a good agreement on the strain localization

behavior observed here for the nanocrystalline samples and

a recent DIC investigation, also on electrodeposited

nanocrystalline Ni [41]. During tensile testing of a 20 nm

Ni, initiation of shear localization was observed during the

uniform deformation stage. Rapid shear-banding evolu-

tion occurred immediately after necking, the macroscopic

strain localization and was concentrated in the middle of

the gauge area, leading ultimately to sample fracture [41].

Stress-induced heat generation

When metals are plastically deformed, part of the plastic

work is dissipated as heat during the deformation process

(e.g., [43, 44]). In the current study, temperature change

due to heat release was measured and recorded by the IR

detector at different global strain levels. Figure 5a shows

selected infrared images of coarse-grained Ni at various

global strain levels, indicated at the bottom in Fig. 5a of

the individual snapshots at a strain rate of 10-1/s.

With the increase of global engineering strain from 15 to

45 %, the localized temperature along the sample gauge

increased continuously as indicated by the color changes

from light blue to yellow and orange. Temperatures are

distributed in a relatively symmetrical fashion along the

gauge with the maximum located in the central region; the

dashed line at a global strain of 50 % indicates the position

of fracture, which coincides with the area of the highest

temperature (i.e., red color) and/or most heat dissipation.

Figure 5b shows the relationship between global strain

and the instantaneous increase in the maximum tempera-

ture, DTmax, on the gauge section during tensile testing atthe strain rate of 10-1/s. During the stage of macroscopic

uniform deformation, DTmax reached approximately42 C at the UTS strain of 45 % (eu in Fig. 5b, i.e., theonset strain of necking). For the initial range of deforma-

tion (045 %), the average rate of temperature increase was

low, \1 C for each percent of global strain. Subsequentdeformation resulted in a rapid local temperature increase

in the necking region up to 70 C for the DTmax at fracturestrain of 51 %. With the addition of the room temperature

of 25 C, the recorded maximum temperature in the gaugelength, Tmax, was approximately 95 C for coarse-grainedNi at the end of room temperature tensile testing.

The measurements for the 23 nm grain size nanocrys-

talline Ni during tensile testing at the strain rate of 10-1/s

are summarized in Fig. 6. At global strains below 3.5 %,

there was a very limited amount of heat generation as

indicated by little color change in the corresponding

infrared images (Fig. 6a). The nanocrystalline Ni displayed

a discernible temperature increase starting at a global

engineering strain of 4.5 %. In the final stage of defor-

mation, a narrow region of high temperature formed around

the fracture location (the dashed line at an angle of 62along the loading axis), indicating a highly concentrated

hot spot for the nanocrystalline Ni. This is somewhat dif-

ferent from the coarse-grained Ni, for which a much

broader and more or less uniform temperature distribution

was observed in the deformation stage close to fracture

(Fig. 5a).

Regarding the evolution of the DTmax, the nanocrystal-line Ni exhibited a qualitatively similar profile as the

coarse-grained Ni. As shown in Fig. 6b, the temperature

increase was low during the uniform deformation stage.

A steep climb in temperature initiated at the onset of

necking, 5.5 % global strain (eu), and all the way to thepoint of fracture (8 %). Specifically, the net increase in

J Mater Sci (2014) 49:38473859 3853

123

temperature was approximately 12 C during the first stageof deformation from 0 to 5.5 % global strain, while a

maximum DTmax of 58 C was recorded at the strain veryclose to the fracture strain of 8.0 % (Fig. 6b). With the

addition of 25 C for room temperature, the recordedmaximum temperature, Tmax, was therefore 83 C for thenanocrystalline Ni during tensile testing.

The two NiFe samples generally exhibited heat gen-

eration behavior similar to the nanocrystalline Ni; the ini-

tial temperature increase before necking was low, and a

rapid increase was observed after necking up to fracture. In

addition, the generated heat was concentrated in the narrow

area of the gauge center, surrounding the ultimate fracture

line for the two nanocrystalline NiFe samples.

Table 3 summarizes the maximum recorded tempera-

tures, Tmax, along with the global elongation recorded at

both strain rates for all materials. It is evident that all

samples experienced local temperature increases during

tensile testing, showing Tmax values noticeably higher than

the ambient testing temperature. Furthermore for a given

sample, the attained maximum temperature, Tmax, is

reduced at lower tensile strain rate. This apparent effect of

strain rate on Tmax reveals the critical influence of heat

conduction on the process of strain-induced heat

generation.

As shown in Table 2, the yield strength, tensile strength

and elongation to fracture of the coarse-grained polycrys-

talline Ni remained virtually the same at the two strain

rates employed, namely 10-2 and 10-1/s. For the material

under scrutiny, same amount of strain-induced heat should

be generated during tensile testing at the both strain rates

due to identical sample size. Because of strain localization,

heat generation was generally not uniform and varied from

one location to another, resulting in heat flow from higher

to lower temperature regions. During tensile testing, the

sample grip sections experienced little deformation and

thus functioned mainly as heat sinks for any strain-induced

heat. In the final stage of necking, most of the heat was

generated in the necking region, i.e., the gauge center, and

was then conducted toward the grip sections at both ends of

the samples. Due to the finite heat conduction of Ni, tem-

perature distribution became a strong function of testing

0

100

200

300

400

500

0 10 20 30 40 50 60

Global strain/Elongation, %

Eng.

stre

ss, M

Pa

0

20

40

60

80

Tm

ax ,

0 C

u

(a)

(b)

600

550

500

450

400

350

300

250

200

Fig. 5 Temperature change incoarse-grained Ni, a infraredimages with the dashed line at

50 % strain indicating the final

fracture position, and

b engineering stressstraincurve and the maximum

temperature increase, DTmax, ata strain rate of 10-1/s. Note that

the numbers on the color scale

bar are temperature in Celsius

(Color figure online)

3854 J Mater Sci (2014) 49:38473859

123

duration, and thus, the recorded maximum temperature

depended on strain rate. Higher strain rate is equivalent to

shorter test duration and therefore more concentrated

in situ temperature distribution, leading to higher maxi-

mum temperature for a given material with the same input

of mechanical energy. Hence, the coarse-grained Ni

exhibited higher Tmax at the higher strain rate (Table 3).

Likewise, the different Tmax values observed at the two

strain rates for the 23 nm Ni and the two 16 nm NiFe

samples can be mainly attributed to the heat conduction. In

addition, temperature distribution/gradients in the nano-

crystalline materials (Fig. 6a) are more concentrated/stee-

per than those in the coarse-grained counterpart (Fig. 5a).

This likely resulted from the facts that for the nanocrys-

talline electrodeposits, the duration of tensile testing was

shorter due to their reduced ductility, and the thermal

conductivity is lower, e.g., up to 25 % reduction as

observed in our recent study for a series of nanocrystalline

Ni with average grain sizes of 2550 nm [45].

Implication of the measured heat generation

Returning to the main focus of this study, the measured

Tmax values observed as a result of strain-induced heat

generation do not support thermal activation as an under-

lying mechanism for the grain growth during room tem-

perature tensile testing reported in previous studies (e.g.,

[10, 18]). For instance, the highest temperature recorded in

the 23 nm grain size nanocrystalline Ni was approximately

83 C during tensile testing, well below 120 C, the

0

400

800

1200

1600

0 2 4 6 8Global strain/Elongation, %

Eng.

stre

ss, M

Pa

0

10

20

30

40

50

60

70

Tm

ax ,

0 C

u

(a)

(b)

600

550

500

450

400

350

300

250

200

Fig. 6 Temperature change inthe 23-nm nanocrystalline Ni,

a infrared images with thedashed line at 8 % strain

indicating fracture location, and

b engineering stressstraincurve versus the maximum

temperature increase at a strain

rate of 10-1/s. Note that the

numbers on the color scale bar

are temperature in Celsius

(Color figure online)

Table 3 Localized strain and maximum temperature recorded duringtensile testing

Samples Grain

size

Strain

rate (/s)

Global strain (%) Tmax(C)

Total At UTS

Poly-Ni 32 lm 10-2 50.4 45 50

10-1 51 45 95

Nano-Ni 23 nm 10-2 8.2 5.5 59

10-1 8 5.5 83

Nano-Ni-2.6 % Fe 16 nm 10-2 7.6 5.5 57

10-1 7.5 5.5 79

Nano-Ni-8.5 % Fe 16 nm 10-2 7.7 5.5 59

10-1 7.5 5.5 81

J Mater Sci (2014) 49:38473859 3855

123

reported onset temperatures for the initiation of micro-

structure evolution, e.g., by subgrain coalescence, for

20 nm grain size Ni electrodeposits [24]. Moreover, the

highest measured Tmax value was approximately 80 C forthe two nanocrystalline NiFe samples, while their grain

growth temperature was reported to be 50100 C higherthan that of the nanocrystalline Ni due to the stabilizing

effect of the alloying element Fe (e.g., [23]). Evidently

based on the measured Tmax alone, the strain-induced heat

release is likely not sufficient to activate migration of grain

boundaries.

One uncertainty regarding the measurement of Tmax is the

spatial resolution of the infrared camera. As shown in Ref.

[10], grain growth during tensile testing may be confined in

a narrow band area less than half a micrometer across. It is

unlikely that the infrared camera with a spatial resolution of

about 50 lm used in the current study can detect tempera-ture differences within a submicron area; the measured Tmaxwas actually an average value over a larger area. It is pos-

sible that a small area in the deformation core attained an

instant local temperature higher than the measured Tmaxduring the tests. The temperature difference between the

maximum and minimum value in the resolution unit area,

DTRes, may be estimated using a 1D heat conduction model,defined by Fouriers law, along the cross-section normal,

DQDt

jA DTDx

1

where j and A are, respectively, thermal conductivity andcross-sectional area; DQ/Dt is the heat flux, leading to thetemperature gradient, DT/Dx, along the heat conductionpath. The heat flux is associated with the rate of

temperature change, b = DT/Dt, through heat capacity Cpat constant pressure,

DQDt

Cp DTDt Cpb 2

Then, DTRes in the area can be assessed as follow bycombining Eqs (1) and (2), replacing Cp with specific heat

cp through Cp = mcp = qALcp, and assuming asymmetrical areal temperature distribution (Dx = L/2)and negligible ambient heat exchange,

DTRes qcpL2

2jb 3

where L and q are the feature length of the resolution area(50 lm) and material density, respectively. Some of theparameters, e.g., cp and b, can have a range of values. In such

cases, the values are chosen to maximize DTRes for thenanocrystalline Ni. Here, the upper value of cp = 0.56 J/g K

is used, based on the reported cp value range of

0.460.56 J/g K within 25200 C for nanocrystalline Ni[46]. The handbook value of q = 8.89 g/cm3 for coarse-

grained Ni [47] can be used here because the density of fully

dense nanocrystalline Ni is grain size independent [48]. The

reported value of j is 90 W/m K for coarse-grained Ni [49].Grain size reduction to about 25 nm resulted in 25 %

reduction of j or j = 67.5 W/m K for the nanocrystallineNi-based on recent observations [45]. The value of b can be

obtained from the slope in the plot of temperature vs strain.

For example in Fig. 6b, the largest slope section is located

at the top or the final stage of deformation, corresponding to

a rate of 635 K/s. Consequently, these values can be used to

calculate an upper bound value of DTRes of approximately0.1 C. This is insignificant with respect to the measuredTmax of 83 C for the nanocrystalline Ni. In other words, theuncertainty due to the spatial resolution of the infrared

camera is expected to be insignificant in the Tmax mea-

surements in the current study.

Another critical parameter is the time or duration avail-

able for thermally activated grain growth. In Ref. [10], the

observed growth from 20 to 200 nm grain size during room

temperature tensile testing at a strain rate of 103/s on the

nanocrystalline Ni occurred over a duration shorter than

10-4 s, corresponding to a rapid growth rate of more than

106 nm/s. This is in contrast to the actual growth rate

observed in an in situ TEM study on electrodeposited Ni

(as-deposited grain size: 20 nm) isothermally annealed at a

much higher temperature, 420 C [50]. In this study, it wasshown that grain boundary (GB) migration occurred as a

series of discontinuous steps followed by periods of

boundary stagnation; the GB velocity profile was repre-

sented by a series of delta functions. With a video sampling

rate of 30 frames per second, the highest recorded instan-

taneous GB migration rate was 700 nm/s, while most grain

boundaries migrated at a rate below 300 nm/s for the

nanocrystalline Ni at 420 C [50]. Much lower averagegrowth rates were observed in an ex-situ TEM study on a

series of nanocrystalline Ni-based electrodeposits, whereby

the same Ni sample as in Ref. [45] showed the fastest

average growth rate of only 33 nm/s after 5 s of isothermal

annealing at 420 C [51]. Note that the nanocrystalline Nisamples used in Refs. [10, 50, 51] were all produced by the

same Integran electrodeposition process and exhibited the

same as-deposited average grain size (20 nm), similar grain

size distributions and likely similar impurity (e.g., S and C)

contents. Therefore, not much difference is expected in the

thermal behavior and stability for the Ni samples of the

three studies. Yet huge differences in growth rates, at least 3

orders of magnitude (106 vs. 700 nm/s), were observed

among the three studies. This indicates a different activation

mechanism for the grain growth shown in Ref. [10] than the

thermally activated growth observed in Refs. [50, 51].

Overall, the measurements of Tmax in the current study

and the growth kinetics comparison between previous

studies on very similar materials suggest that thermally

3856 J Mater Sci (2014) 49:38473859

123

activated grain growth during room temperature tensile

testing for nanocrystalline nickel is highly unlikely. Hence,

the reported grain growth during tensile testing in previous

studies may be attributed to other contributing processes,

e.g., stress-induced grain growth.

In early studies on superplastic deformation, stress-

induced grain growth was reported in conventional coarse-

grained materials (e.g., [52, 53]). Given the significance of

grain size in nanostructures, this type of grain growth has

been an on-going focus of numerous experimental and

molecular dynamics simulation (MD) studies on nano-

crystalline materials in recent years, [5466]. Evidence has

been presented for stress-induced grain growth through

direct experimental observations. For instance, stress-

induced grain rotation and growth were observed during

in situ TEM tensile testing on nanocrystalline Ni electro-

deposits [54, 63]. Again, the materials used in Ref. [63]

were produced by the same Integran electrodeposition

process as the one in the current study. In another in situ

TEM study, grain growth occurred upon nanoindentation in

nanocrystalline Al films [57]. Additionally, grains were

found to grow faster at cryogenic temperature than at room

temperature in a study on strained nanocrystalline Cu [58].

Further, it was recently concluded, based on the results of

tensile testing coupled with transmission electron micros-

copy on nanocrystalline Al films that grain growth was a

shear stress-driven process [64]. Generally, the experi-

mental observations were in agreement with the results

from MD studies, e.g., on nanocrystalline Pd [54], Cu [56]

and Ni [62, 66].

Summary

Temperature increases were measured during tensile test-

ing of both coarse-grained Ni and nanocrystalline Ni and

NiFe electrodeposits due to strain-induced heat genera-

tion. Generally, higher strain values and/or higher strain

rates resulted in higher observed peak temperature for each

material, with the highest temperatures recorded in the

regions of subsequent fracture. For the coarse-grained

polycrystalline Ni, the maximum attained mechanically

induced temperature was approximately 95 C, whereasthe highest temperature for the nanocrystalline electrode-

posits was approximately 83 C. This is significantly belowthe reported threshold temperatures for the onset of ther-

mally activated grain growth in nanocrystalline Ni and

NiFe alloys. Therefore, thermally activated grain growth

is highly unlikely for nanocrystalline Ni and NiFe during

room temperature tensile testing. Previously reported grain

growth events observed in similar materials are likely due

to stress-driven grain boundary migration.

As per DIC measurements, the samples with grain sizes

ranging from 16 nm to 32 lm exhibited similar strainlocalization behavior. Local strain variations initiated in

the early stages of macroscopic uniform deformation

and became more pronounced close to the point of peak

load or UTS. Subsequent necking took place in the region

where significant strain localization initiated. Material

grain size showed little influence on strain localization

behavior. Some grain size dependent strain rate sensitivity

was observed during the tensile testing at two strain rates of

10-1 and 10-2/s. While the coarse-grained Ni exhibited

insensitivity to strain rate, gradually increased strain rate

sensitivity was demonstrated in the 23 nm grain size Ni

and the two 16 nm grain size NiFe samples. This sensi-

tivity suggests that both dislocation-mediated and grain-

boundary-controlled mechanisms were operative in the

deformation of the nanocrystalline Ni and NiFe samples.

Acknowledgements The authors would like to thank Dr. DavidBackman, Mr. Richard Bos and Mr. Thomas Sears from the National

Research Council, Institute of Aerospace Research for their help and

many valuable suggestions during the DIC and infrared experiments.

YZ would like to thank Mr. Cho for the fruitful discussion on thermal

and electric transport in nanocrystalline metals. Highly appreciated is

the financial support by the Natural Sciences and Engineering

Research Council of Canada (NSERC) and the Ontario Research

Fund (ORF).

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Localized strain and heat generation during plastic deformation in nanocrystalline Ni and Ni--FeAbstractIntroductionExperimentalResults and discussionGrain size and chemical compositionDeformation at different tensile strain ratesStrain localizationStress-induced heat generationImplication of the measured heat generation

SummaryAcknowledgementsReferences