《金属材料强韧化与组织调控》教学资源（参考文献）Rapid hardening induced by electric pulse annealing in nanostructured pure aluminum

Available online at www.sciencedirect.com SciVerse ScienceDirect Scripta MATERIALIA ELSEVIER Scripta Materialia 66(2012)147-150 www.elsevier.com/locate/scriptamat Rapid hardening induced by electric pulse annealing in nanostructured pure aluminum Wei Zeng,a Yao Shen,a*Ning Zhang,a Xiaoxu Huang,*Jeff Wang, Guoyi Tanga and Aidang Shana State Key Lab.of Metal Matrix Composite,School of Materials Science and Engineering. Shanghai Jiao Tong University,Shanghai 200240,People's Republic of China Danish-Chinese Center for Nanometals,Materials Research Division,Riso National Laboratory for Sustainable Energy. Technical University of Denmark.DK-4000 Roskilde,Denmark General Motors China Science Laboratory,Shanghai 201206,People's Republic of China dAdvanced Materials Institute,Graduate School at Shenzhen,Tsinghua University,Shenzhen 518055. People's Republic of China Received 14 July 2011;revised 2 October 2011;accepted 13 October 2011 Available online 20 October 2011 Nanostructured pure aluminum was fabricated by heavy cold-rolling and then subjected to recovery annealing either by applying electric pulse annealing or by traditional air furnace annealing.Both annealing treatments resulted in an increase in yield strength due to the occurrence of a "dislocation source-limited hardening"mechanism.However,the hardening kinetics was substantially faster for the electric pulse annealed material.Detailed microstructural characterization suggested that the rapid hardening during electric pulse annealing is related to an enhanced rate of recovery of dislocation structures. 2011 Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved. Keywords:Annealing:Hardening:Rolling:Aluminum A characteristic structural feature of nanostruc- have been observed,such as annealing-induced increases tured metals produced by plastic deformation to high in yield strength[1,12-16]and decreases in tensile elonga- strains is the presence of dislocations lying inside cells, tion (annealing-induced hardening)[1,14,15,17]and which can have a density of the order of 1014 m-2 for deformation-induced softening [1,6,18].Analysis of the face-centered cubic and body-centered cubic nanostruc- correlation between the change in the dislocation struc- tured metals with medium to high stacking fault energy ture and the unexpected annealing-induced hardening [1-6].These dislocations are present in the form of indi- and deformation-induced softening has led to a sugges- vidual dislocations or as loose dislocation tangles,and tion [1,5]that a new strengthening mechanism,referred their density and arrangement are expected to affect the to as dislocation source-limited hardening [1,5],occurs mechanical behavior of the nanostructured metals.For in nanostructured metals if they lack mobile dislocations example,enhanced strain-rate sensitivity as compared and easy dislocation sources.From this mechanism,one with coarse-grained materials has been observed in nano- may expect that in addition to the occurrence of anneal- structured metals produced by equal-channel angular ing-induced hardening,the hardening kinetics can be pressing (ECAP)[7-10]and accumulative roll bonding manipulated if the rate of decrease of dislocation density (ARB)[11],which indicates [8]a reduced activation vol- can be changed,e.g.by different annealing treatments. ume resulting from the fine-scale structure and the higher It has been reported in many studies that the mobility mobile dislocation density in the nanostructured metals. of dislocations can be increased by application of electric Post-process annealing or plastic deformation has shown pulses during plastic deformation or during annealing, strong effects on the mechanical behavior of nanostruc- although the exact mechanisms underlying this phenom- tured materials.Most interestingly,unexpected behaviors enon is still open to debate [19-27].Therefore an enhance- ment of dislocation annihilation is expected during *Corresponding authors.E-mail addresses:yaoshen@sjtu.edu.cn; electric pulse annealing(EPA)as compared to conven- xihu@risoe.dtu.dk tional annealing at a similar temperature.If this is the 1359-6462/S-see front matter2011 Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved. doi:10.1016/j.scriptamat.2011.10.023

Rapid hardening induced by electric pulse annealing in nanostructured pure aluminum Wei Zeng,a Yao Shen,a,⇑ Ning Zhang,a Xiaoxu Huang,b,⇑ Jeff Wang,c Guoyi Tangd and Aidang Shana a State Key Lab. of Metal Matrix Composite, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, People’s Republic of China b Danish–Chinese Center for Nanometals, Materials Research Division, Risø National Laboratory for Sustainable Energy, Technical University of Denmark, DK-4000 Roskilde, Denmark c General Motors China Science Laboratory, Shanghai 201206, People’s Republic of China d Advanced Materials Institute, Graduate School at Shenzhen, Tsinghua University, Shenzhen 518055, People’s Republic of China Received 14 July 2011; revised 2 October 2011; accepted 13 October 2011 Available online 20 October 2011 Nanostructured pure aluminum was fabricated by heavy cold-rolling and then subjected to recovery annealing either by applying electric pulse annealing or by traditional air furnace annealing. Both annealing treatments resulted in an increase in yield strength due to the occurrence of a “dislocation source-limited hardening” mechanism. However, the hardening kinetics was substantially faster for the electric pulse annealed material. Detailed microstructural characterization suggested that the rapid hardening during electric pulse annealing is related to an enhanced rate of recovery of dislocation structures. 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Annealing; Hardening; Rolling; Aluminum A characteristic structural feature of nanostruc￾tured metals produced by plastic deformation to high strains is the presence of dislocations lying inside cells, which can have a density of the order of 1014 m2 for face-centered cubic and body-centered cubic nanostruc￾tured metals with medium to high stacking fault energy [1–6]. These dislocations are present in the form of indi￾vidual dislocations or as loose dislocation tangles, and their density and arrangement are expected to affect the mechanical behavior of the nanostructured metals. For example, enhanced strain-rate sensitivity as compared with coarse-grained materials has been observed in nano￾structured metals produced by equal-channel angular pressing (ECAP) [7–10] and accumulative roll bonding (ARB) [11], which indicates [8] a reduced activation vol￾ume resulting from the fine-scale structure and the higher mobile dislocation density in the nanostructured metals. Post-process annealing or plastic deformation has shown strong effects on the mechanical behavior of nanostruc￾tured materials. Most interestingly, unexpected behaviors have been observed, such as annealing-induced increases in yield strength [1,12–16] and decreases in tensile elonga￾tion (annealing-induced hardening) [1,14,15,17] and deformation-induced softening [1,6,18]. Analysis of the correlation between the change in the dislocation struc￾ture and the unexpected annealing-induced hardening and deformation-induced softening has led to a sugges￾tion [1,5] that a new strengthening mechanism, referred to as dislocation source-limited hardening [1,5], occurs in nanostructured metals if they lack mobile dislocations and easy dislocation sources. From this mechanism, one may expect that in addition to the occurrence of anneal￾ing-induced hardening, the hardening kinetics can be manipulated if the rate of decrease of dislocation density can be changed, e.g. by different annealing treatments. It has been reported in many studies that the mobility of dislocations can be increased by application of electric pulses during plastic deformation or during annealing, although the exact mechanisms underlying this phenom￾enon is still open to debate [19–27]. Therefore an enhance￾ment of dislocation annihilation is expected during electric pulse annealing (EPA) as compared to conven￾tional annealing at a similar temperature. If this is the 1359-6462/$ - see front matter 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2011.10.023 ⇑ Corresponding authors. E-mail addresses: yaoshen@sjtu.edu.cn; xihu@risoe.dtu.dk Available online at www.sciencedirect.com Scripta Materialia 66 (2012) 147–150 www.elsevier.com/locate/scriptamat

148 W.Zeng et al.Scripta Materialia 66 (2012)147-150 case,an enhanced hardening kinetics can also be ex- Figure I shows the tensile curves of nanostructured pected.In this study,we apply EPA to nanostructured samples in the as-cold-rolled state and after annealing Al produced by heavy cold-rolling to explore the effect by AFA at 150C for 30 min or by EPA for 0.5,1,2 of this treatment on the kinetics of annealing-induced and 5 min.Compared with the as-cold-rolled sample. hardening.The results will also shed light on the occur- all the annealed samples show an increase in the yield rence of the dislocation source-limited hardening mecha- strength and a decrease in elongation,similar to previ- nism in nanostructured pure Al. ous observations in annealed nanostructured Al pro- To prepare a fine-scale structure,a recrystallized Al cessed by ARB [1,5]or ECAP [12-14]or cold rolling plate (see Table 1 for the chemical composition) [16].The strength and elongation for the peak hardening 40 mm thick was cold-rolled to 0.6 mm (corresponding conditions are given in Table 1.In the case of EPA,peak to an equivalent von Mises strain of svM=4.85)by 12 hardening is observed after 2 min EPA treatment.After rolling passes.The rolling reduction applied in each pass 5 min EPA treatment the stress-strain curve is shifted was to keep the ratio between the contact length and the down,although the yield strength is still higher than that sample thickness in the range of 1-4,which ensures that for the as-cold-rolled sample.Figure 2 shows the change the deformation penetrates through the sample thick- in the yield strength as a function of annealing time up ness.The sample was cooled in liquid nitrogen after each to the peak hardening condition.In the plot,the results successive rolling pass to suppress the effect of heat gen- from a previous study [16]on AFA of the same material erated during rolling. cold-rolled to a similar strain are included.Similarities Tensile specimens with gauge dimensions of 18 mm in both the yield strength of the as-cold-rolled state long x 5 mm wide were machined using electrodischarge and in the hardening behavior during AFA are observed machining from the cold-rolled sheets with the tension between the present and previous [16]nanostructured Al direction parallel to the transverse direction of the rolled samples,which validates the effect of AFA treatment on sheet.Some of the machined tensile specimens were sub- the yield strength.When comparing the hardening jected to the EPA and air furnace annealing (AFA) behavior between the samples treated by EPA and treatments,which were done within a month of the spec- AFA.two remarkable differences are found.The first imens being prepared.EPA was conducted using 50 Hz is the substantially enhanced hardening rate by EPA electric pulses with an average current density of compared to AFA:the annealing time required to reach 500 A cm.The reason for using such a relatively low the peak hardening is 2 min for the EPA treatment, current density is to limit the temperature rise generated whereas it is 30 min for the AFA treatment.The second by Joule heating.The electric pulses were directly ap- is the higher maximum yield strength achieved by EPA plied to a tensile coupon by clamping electrodes at the than by AFA (190.2 vs.179.7 MPa). two sample ends.The duration of each single pulse A comparison of annealing parameters (temperature was 2 ms,while the pulse period,including one pulse and heating rate)was made between the two annealing and the following waiting time,was 20 ms (i.e.a fre- procedures to understand whether the observed differ- quency of 50 Hz).The total EPA treatment duration ences in the mechanical behavior are caused by the dif- was varied from 0.5 to 5 min.The temperature rise ferences in the thermal history.The highest temperature due to Joule heating was monitored by a thermocouple for the EPA treatment was measured to be in the range attached to the surface of the sample.For comparison, of 120-150C,with the average temperature well below AFA was conducted at 150C for 30 min.which corre- the temperature of 150C for AFA.In this case,the sponds to the peak hardening condition established in a maximum heating rate during EPA can be calculated previous study on the effect of annealing condition(tem- to be 2.8C s by assuming adiabatic heating.For perature and time)on the yield strength of the same AFA,because of the small thickness of the tensile spec- material cold-rolled to a similar strain(EvM=4.51)[16]. imen(0.6 mm),it takes less than 2 min for the tempera- The tensile tests were performed at room temperature ture of the specimen to reach the preset annealing at a nominal strain rate of 4 x 10-4s-on a Zwick/Roel temperature of 150C after the sample is put into the testing machine.The yield strength was determined as furnace.This corresponds to an average heating rate the 0.2%proof stress.To evaluate the variation of yield of above ICs-,which is similar to the heating rate strength within a given rolled sheet.18 tensile specimens estimated for EPA.The above analysis suggests that taken from a single rolled sheet were tested.It was found the more rapid hardening kinetics and the larger maxi- that the yield strengths measured from these 18 speci- mum yield strength induced by the EPA treatment are mens were rather constant with a standard deviation not directly related to the annealing parameters. of +2 MPa around the average.To eliminate the effect Detailed microstructural analysis was conducted to of the strength variation from sheet to sheet,the data re- understand the underlying mechanisms responsible for ported in this paper were obtained from tensile speci- the rapid hardening induced by EPA.The microstruc- mens that were cut from a single sheet and tested in ture was characterized by transmission electron micros- the as-rolled condition and after being subjected to var- copy (TEM)with a JEOL 2000FX operated at 200 kV. ious annealing treatments. Samples for TEM analysis were made from the longitu- Table 1.Chemical composition of the pure Al sample used. Element Fe Ni Si Ti Zn Content (wt.%) 0.0019 0.1420 0.0211 0.0043 0.0306 0.0039 0.0445 Balance

case, an enhanced hardening kinetics can also be ex￾pected. In this study, we apply EPA to nanostructured Al produced by heavy cold-rolling to explore the effect of this treatment on the kinetics of annealing-induced hardening. The results will also shed light on the occur￾rence of the dislocation source-limited hardening mecha￾nism in nanostructured pure Al. To prepare a fine-scale structure, a recrystallized Al plate (see Table 1 for the chemical composition) 40 mm thick was cold-rolled to 0.6 mm (corresponding to an equivalent von Mises strain of evM = 4.85) by 12 rolling passes. The rolling reduction applied in each pass was to keep the ratio between the contact length and the sample thickness in the range of 1–4, which ensures that the deformation penetrates through the sample thick￾ness. The sample was cooled in liquid nitrogen after each successive rolling pass to suppress the effect of heat gen￾erated during rolling. Tensile specimens with gauge dimensions of 18 mm long 5 mm wide were machined using electrodischarge machining from the cold-rolled sheets with the tension direction parallel to the transverse direction of the rolled sheet. Some of the machined tensile specimens were sub￾jected to the EPA and air furnace annealing (AFA) treatments, which were done within a month of the spec￾imens being prepared. EPA was conducted using 50 Hz electric pulses with an average current density of 500 A cm2 . The reason for using such a relatively low current density is to limit the temperature rise generated by Joule heating. The electric pulses were directly ap￾plied to a tensile coupon by clamping electrodes at the two sample ends. The duration of each single pulse was 2 ms, while the pulse period, including one pulse and the following waiting time, was 20 ms (i.e. a fre￾quency of 50 Hz). The total EPA treatment duration was varied from 0.5 to 5 min. The temperature rise due to Joule heating was monitored by a thermocouple attached to the surface of the sample. For comparison, AFA was conducted at 150 C for 30 min, which corre￾sponds to the peak hardening condition established in a previous study on the effect of annealing condition (tem￾perature and time) on the yield strength of the same material cold-rolled to a similar strain (evM = 4.51) [16]. The tensile tests were performed at room temperature at a nominal strain rate of 4 104 s 1 on a Zwick/Roel testing machine. The yield strength was determined as the 0.2% proof stress. To evaluate the variation of yield strength within a given rolled sheet, 18 tensile specimens taken from a single rolled sheet were tested. It was found that the yield strengths measured from these 18 speci￾mens were rather constant with a standard deviation of ±2 MPa around the average. To eliminate the effect of the strength variation from sheet to sheet, the data re￾ported in this paper were obtained from tensile speci￾mens that were cut from a single sheet and tested in the as-rolled condition and after being subjected to var￾ious annealing treatments. Figure 1 shows the tensile curves of nanostructured samples in the as-cold-rolled state and after annealing by AFA at 150 C for 30 min or by EPA for 0.5, 1, 2 and 5 min. Compared with the as-cold-rolled sample, all the annealed samples show an increase in the yield strength and a decrease in elongation, similar to previ￾ous observations in annealed nanostructured Al pro￾cessed by ARB [1,5] or ECAP [12–14] or cold rolling [16]. The strength and elongation for the peak hardening conditions are given in Table 1. In the case of EPA, peak hardening is observed after 2 min EPA treatment. After 5 min EPA treatment the stress–strain curve is shifted down, although the yield strength is still higher than that for the as-cold-rolled sample. Figure 2 shows the change in the yield strength as a function of annealing time up to the peak hardening condition. In the plot, the results from a previous study [16] on AFA of the same material cold-rolled to a similar strain are included. Similarities in both the yield strength of the as-cold-rolled state and in the hardening behavior during AFA are observed between the present and previous [16] nanostructured Al samples, which validates the effect of AFA treatment on the yield strength. When comparing the hardening behavior between the samples treated by EPA and AFA, two remarkable differences are found. The first is the substantially enhanced hardening rate by EPA compared to AFA: the annealing time required to reach the peak hardening is 2 min for the EPA treatment, whereas it is 30 min for the AFA treatment. The second is the higher maximum yield strength achieved by EPA than by AFA (190.2 vs. 179.7 MPa). A comparison of annealing parameters (temperature and heating rate) was made between the two annealing procedures to understand whether the observed differ￾ences in the mechanical behavior are caused by the dif￾ferences in the thermal history. The highest temperature for the EPA treatment was measured to be in the range of 120–150 C, with the average temperature well below the temperature of 150 C for AFA. In this case, the maximum heating rate during EPA can be calculated to be 2.8 C s1 by assuming adiabatic heating. For AFA, because of the small thickness of the tensile spec￾imen (0.6 mm), it takes less than 2 min for the tempera￾ture of the specimen to reach the preset annealing temperature of 150 C after the sample is put into the furnace. This corresponds to an average heating rate of above 1 C s1 , which is similar to the heating rate estimated for EPA. The above analysis suggests that the more rapid hardening kinetics and the larger maxi￾mum yield strength induced by the EPA treatment are not directly related to the annealing parameters. Detailed microstructural analysis was conducted to understand the underlying mechanisms responsible for the rapid hardening induced by EPA. The microstruc￾ture was characterized by transmission electron micros￾copy (TEM) with a JEOL 2000FX operated at 200 kV. Samples for TEM analysis were made from the longitu￾Table 1. Chemical composition of the pure Al sample used. Element B Fe V Ni Si Ti Zn Al Content (wt.%) 0.0019 0.1420 0.0211 0.0043 0.0306 0.0039 0.0445 Balance 148 W. Zeng et al. / Scripta Materialia 66 (2012) 147–150

W.Zeng et al.ISeripta Materialia 66(2012)147-150 149 EPA,2 min EPA,1min (b) EPA,0.5min Cold rolled tou■4.85 200 AFA,30 min/150-C 150 As cold rolled EPA,5min Buu 100 p止 50 10 2030405060 70 Msorientasion ange(deg) Figure 3.(a)TEM image showing a lamellar structural morphology 2 345 and dislocation configurations in nanostructured Al processed by cold Engieering strain (% rolling to M=4.85.(b)Histogram showing the distribution of Figure 1.Stress-strain curves of 99.7%pure Al in the as-cold-rolled boundary misorientation angles measured by Kikuchi diffraction. state(EM=4.85)and after EPA and AFA treatments. around 50.The fraction of high-angle boundaries 200 (>15)is about 50%.The misorientation characteristics ▲6w=4.85.EPA measured in this cold-rolled sample are therefore very E=4.85,AFA similar to those observed in nanostructured Al samples 190 7vw-4.51,AFA[16l processed by ARB [1,5,14]or ECAP [30]. Figure 4 shows the microstructural observations for samples that had experienced annealing under peak hard- ening conditions by AFA and EPA.The lamellar struc- tural features remain unchanged after either AFA or EPA treatment,but the decrease after annealing in dislo- 160 cation density is obvious(compare,e.g.,Fig.3a).Mea- 0 10 20 surements of boundary spacings and dislocation Annealing time(min) densities are summarized in Table 1.It is seen that a mod- Figure 2.The yield strength as a function of annealing time of 99.7% erate structural coarsening occurred in the AFA sample, pure Al cold-rolled to M=4.85 by EPA or AFA. with the lamellar boundary spacing increasing from 280 nm in the as-cold-rolled state to 356 nm after anneal- dinal section,which contains the normal direction(ND) ing.However,the structural coarsening is negligible in the and the rolling direction(RD)of deformed sheet,and is EPA sample(Table 1).The dislocation density measure- the best section to reveal the detailed structural features ments show a similar value,5 x 103m2,in the AFA in cold-rolled samples [28].A semiautomatic Kikuchi and the EPA samples,showing a clear decrease as com- pattern analysis technique [29]was used for measure- pared with the as-cold-rolled state.The decrease in the ments of orientation and misorientation.To resolve dislocation density is in general caused by annihilation the dislocations present in a grain(subgrain)for disloca- of dislocations of opposite signs and by the removal of tion density measurement,TEM images were recorded dislocations into sinks at high-angle boundaries.Bound- when the zone axis of either [001]or [01 1]of the grain ary migration associated with the structural coarsening is approximately parallel to the beam direction.The dis- provides an additional mechanism for dislocation recov- location density was measured using the intercept meth- ery.Because of the negligible structural coarsening in the od.For each condition,about 150-250 dislocations were EPA sample,the former two mechanisms must be respon- counted from the TEM images of about 20 grains.The sible for the dislocation recovery.Furthermore,because foil thickness was determined by means of convergent of the lower temperature and the much shorter treatment beam electron diffraction. time during EPA,the similar decrease in the dislocation The as-cold-rolled sample exhibited a typical fine- density suggests that the recovery process is accelerated, scale lamellar structure [1,6]characterized by extended leading to an enhanced rate of dislocation density de- lamellar boundaries parallel to the rolling plane and crease during EPA.Coupling the microstructure with short transverse boundaries interconnecting the lamellar the mechanical behavior suggests that the rapid anneal- boundaries,as shown in Figure 3a.The boundary spac- ing-induced hardening during EPA has something to do ings were measured from a number of TEM images and with the accelerated decrease in the dislocation density. about 100-200 spacings were counted for each measure- It has been shown that the annealing-induced hardening ment;the average values and standard deviations are is insensitive to the impurity up to levels of 99.2-99.99% shown in Table 2.The presence of individual disloca- [1,5].However,whether the EPA treatment can enhance tions or loose dislocation tangles was observed in indi- the impurity effect on the hardening kinetics by,for exam- vidual lamellae.The dislocation density was measured ple,accelerating the segregation of impurity elements to be 8x 1013m-2(Table 2),which is comparable to into high-angle boundaries needs further clarification. the values of dislocation density measured in nanostruc- The second difference in the annealing-induced hard- tured Al produced by ARB [1,4,5].The combined distri- ening between EPA and AFA is the larger maximum bution of boundary misorientation angles,for the yield strength(peak hardening)achieved by EPA com- lamellar boundaries and the interconnecting boundaries, pared with AFA.As the dislocation density change is is shown in Figure 3b.A bimodal distribution is seen, similar after annealing at the peak hardening condition with one peak located below 2 and the other located (EPA 2 min,AFA 30 min/150C),a similar contribution

dinal section, which contains the normal direction (ND) and the rolling direction (RD) of deformed sheet, and is the best section to reveal the detailed structural features in cold-rolled samples [28]. A semiautomatic Kikuchi pattern analysis technique [29] was used for measure￾ments of orientation and misorientation. To resolve the dislocations present in a grain (subgrain) for disloca￾tion density measurement, TEM images were recorded when the zone axis of either [0 0 1] or [0 1 1] of the grain is approximately parallel to the beam direction. The dis￾location density was measured using the intercept meth￾od. For each condition, about 150–250 dislocations were counted from the TEM images of about 20 grains. The foil thickness was determined by means of convergent beam electron diffraction. The as-cold-rolled sample exhibited a typical fine￾scale lamellar structure [1,6] characterized by extended lamellar boundaries parallel to the rolling plane and short transverse boundaries interconnecting the lamellar boundaries, as shown in Figure 3a. The boundary spac￾ings were measured from a number of TEM images and about 100–200 spacings were counted for each measure￾ment; the average values and standard deviations are shown in Table 2. The presence of individual disloca￾tions or loose dislocation tangles was observed in indi￾vidual lamellae. The dislocation density was measured to be 8 1013 m2 (Table 2), which is comparable to the values of dislocation density measured in nanostruc￾tured Al produced by ARB [1,4,5]. The combined distri￾bution of boundary misorientation angles, for the lamellar boundaries and the interconnecting boundaries, is shown in Figure 3b. A bimodal distribution is seen, with one peak located below 2 and the other located around 50. The fraction of high-angle boundaries (>15) is about 50%. The misorientation characteristics measured in this cold-rolled sample are therefore very similar to those observed in nanostructured Al samples processed by ARB [1,5,14] or ECAP [30]. Figure 4 shows the microstructural observations for samples that had experienced annealing under peak hard￾ening conditions by AFA and EPA. The lamellar struc￾tural features remain unchanged after either AFA or EPA treatment, but the decrease after annealing in dislo￾cation density is obvious (compare, e.g., Fig. 3a). Mea￾surements of boundary spacings and dislocation densities are summarized in Table 1. It is seen that a mod￾erate structural coarsening occurred in the AFA sample, with the lamellar boundary spacing increasing from 280 nm in the as-cold-rolled state to 356 nm after anneal￾ing. However, the structural coarsening is negligible in the EPA sample (Table 1). The dislocation density measure￾ments show a similar value, 5 1013 m2 , in the AFA and the EPA samples, showing a clear decrease as com￾pared with the as-cold-rolled state. The decrease in the dislocation density is in general caused by annihilation of dislocations of opposite signs and by the removal of dislocations into sinks at high-angle boundaries. Bound￾ary migration associated with the structural coarsening provides an additional mechanism for dislocation recov￾ery. Because of the negligible structural coarsening in the EPA sample, the former two mechanisms must be respon￾sible for the dislocation recovery. Furthermore, because of the lower temperature and the much shorter treatment time during EPA, the similar decrease in the dislocation density suggests that the recovery process is accelerated, leading to an enhanced rate of dislocation density de￾crease during EPA. Coupling the microstructure with the mechanical behavior suggests that the rapid anneal￾ing-induced hardening during EPA has something to do with the accelerated decrease in the dislocation density. It has been shown that the annealing-induced hardening is insensitive to the impurity up to levels of 99.2–99.99% [1,5]. However, whether the EPA treatment can enhance the impurity effect on the hardening kinetics by, for exam￾ple, accelerating the segregation of impurity elements into high-angle boundaries needs further clarification. The second difference in the annealing-induced hard￾ening between EPA and AFA is the larger maximum yield strength (peak hardening) achieved by EPA com￾pared with AFA. As the dislocation density change is similar after annealing at the peak hardening condition (EPA 2 min, AFA 30 min/150 C), a similar contribution Figure 1. Stress–strain curves of 99.7% pure Al in the as-cold-rolled state (evM = 4.85) and after EPA and AFA treatments. Figure 3. (a) TEM image showing a lamellar structural morphology and dislocation configurations in nanostructured Al processed by cold rolling to evM = 4.85. (b) Histogram showing the distribution of boundary misorientation angles measured by Kikuchi diffraction. Figure 2. The yield strength as a function of annealing time of 99.7% pure Al cold-rolled to evM = 4.85 by EPA or AFA. W. Zeng et al. / Scripta Materialia 66 (2012) 147–150 149

150 W.Zeng et al./Scripta Materialia 66(2012)147-150 Table 2.Boundary spacing,dislocation density and mechanical properties of the as-cold-rolled state and after EPA or AFA under peak hardening conditions. Sample GYs (MPa)GUTs (MPa)euniform (%etotal (% p%(103m-3 Dzs(nm）Dcs(nm) Dg (nm) As-cold-rolled 167.1 188.2 2.0 5.9 8 280 600 345 AFA for 30 min/150C 179.7 194.4 1.5 4.5 5 356 680 449 EPA for2min/120-150C)190.2 206.2 1.5 48 310 607 351 Note:DL and Dic are the distances between the lamellar boundaries and interconnecting boundaries,Dg is the average grain size along a random test line.The standard deviations of these spacing measurements are 5-15 nm.Po is the dislocation density in the volumes between boundaries (standard deviation is 0.6-1.0 x 103m). (b) [1]X.Huang,N.Hansen,N.Tsuji,Science 312 (2006)249. [2]Y.H.Zhao,Y.T.Zhu,X.Z.Liao,Z.Horita,T.G Langdon,Appl.Phys.Lett.89(2006)121906. [3]Y.H.Zhao,J.F.Bingert,Y.T.Zhu,X.Z.Liao,R.Z. Valiev,Z.Horita,T.G.Langdon,Y.Z.Zhou,E.J. Lavernia,Appl.Phys.Lett.92(2008)081903. [4]Y.Miyajima,M.Mitsuhara,S.Hata,H.Nakashima,N. 500nm 500nm Tsuji,Mater.Sci.Eng.A 528(2010)776 [5]N.Kamikawa,X.Huang,N.Tsuji,N.Hansen,Acta Figure 4.TEM images showing the microstructural features and Mater.57(2009)4198-4208. dislocation configurations in nanostructured Al processed by cold [6]X.Huang,N.Kamikawa Tsuji,N.Hansen,J.Mater.Sci. rolling to EyM =4.85 followed by (a)AFA at 150 C for 30 min and (b) 45(2011)4761 EPA for 2 min. [7]G.T.Gray III,T.C.Lowe,C.M.Cady,R.Z.Valiev,I.V. Aleksandrov.Nanostuct.Mater.9 (1997)477 to the hardening effect from the change in dislocation [8]Q.Wei,S.Cheng,K.T.Ramesh,E.Ma,Mater.Sci.Eng. density is expected.The lower peak yield strength A381(200471. (Fig.2)in the case of AFA could therefore be caused [9]Y.M.Wang,E.Ma,Acta Mater.52 (2004)1699 by the structural coarsening.Using the average bound- [10]J.May,H.W.Hoppel,M.Goken,Scr.Mater.53(2005)189. ary spacing,Dg (Table 2),and the well-established [11]W.H.Hoppel,J.May,M.Goken,Adv.Eng.Mater.6 Hall-Petch slope (0.04 MPa mP)for Al of similar pur- (2004)781. ity [31],the difference in the Hall-Petch contribution to [12]J.R.Bowen,P.B.Prangnell,D.Juul Jensen,N.Hansen, Mater.Sci.Eng.A 387 (2004)235. the yield strength was estimated to be 9 MPa between [13]R.Z.Valiev,A.V.Sergueeva,A.K.Mukherjee,Scr.Mater the samples treated by EPA and AFA to peak harden- 49(2003)669-674. ing.This calculated difference accounts for the experi- [14]C.Y.Yu,P.W.Kao,C.P.Chang,Acta Mater.53(2005)4019. mentally observed difference (Table 2). [15]X.Huang,N.Kamikawa,N.Hansen,Mater.Sci.Eng.A In summary,substantially enhanced hardening kinetics 493(2008)184. were found during EPA of a nanostructured Al produced [16]N.Zhang,Master thesis,Shanghai Jiao Tong University, by heavy cold-rolling.Microstructural characterizations 2009). suggest that this behavior may be attributed to an en- [17]N.Tsuji,Y.Ito,Y.Saito,Y.Minamino,Scr.Mater.47 hanced rate of dislocation recovery caused by the greatly 2002)893. increased mobility of dislocations.This finding shows the [18]X.Huang,Scr.Mater.60(2009)1078. [19]K.Okazaki,M.Kagawa,H.Conrad,Scr.Metall.12 potential of using EPA to manipulate the microstructure (1978)1063-1068. of materials over a wider spectrum,by substantially [20]V.L.A.Silveira,M.F.S.Porto,W.A.Mannheimer,Scr. increasing the recovery rate of dislocation structures with Metall.15(1981)945. little coarsening in the structural scale.This potential facil- 21]H.Conrad,N.Karam,S.Mannan.Scr.Metall.17(1983)411. itates not only the quantitative separation of the contribu- [22]A.F.Sprecher,S.L.Mannan,H.Conrad,Acta Metall.34 tions of different strengthening mechanisms to the (1986)1145. mechanical strength,but also the development of new [23]M.Molotskii,V.Fleurov,Phys.Rev.Lett.78 (1997) strategies for optimization of the structure and properties 2779. of nanostructured materials. [24]Y.Z.Zhou,S.Xiao,J.D.Guo,Mater.Lett.58 (2004) 1948. 25]H.Conrad,Mater.Sci.Eng.A 287(2000)276. This work was supported by General Motors 26]F.Ding,G.Y.Tang,Z.H.Xu.S.Q.Tian,J.Mater.Sci. China Science Lab.Y.S.also acknowledges financial Technol.23(2007)273. support from the National Science Foundation of China [27]R.Delville,B.Malard,J.Pilch,P.Sittner,D.Schryvers, through projects 50601018,50971090 and 50890174. Acta Mater.58 (2010)4503. X.H.was supported by the Danish National Research [28]G.Winther,X.Huang,A.Godfrey,N.Hansen,Acta Foundation through the Danish-Chinese Center for Mater.52(20044437. Nanometals.The authors thank A.Godfrey and N. [29]Q.Liu,Ultramicroscopy 60(1995)81. Hansen for helpful discussions.Y.S.thank A.H.W. [30]M.Cabibbo,W.Blum,E.Evangelista,M.E.Kassner, M.A.Meyers.Metall.Mater.Trans.A 39(2008)181. Ngan for a visiting position in the University of Hong [31]N.Hansen,Acta Metall.25(1977)863. Kong and helpful discussions

to the hardening effect from the change in dislocation density is expected. The lower peak yield strength (Fig. 2) in the case of AFA could therefore be caused by the structural coarsening. Using the average bound￾ary spacing, DR (Table 2), and the well-established Hall–Petch slope (0.04 MPa m1/2) for Al of similar pur￾ity [31], the difference in the Hall–Petch contribution to the yield strength was estimated to be 9 MPa between the samples treated by EPA and AFA to peak harden￾ing. This calculated difference accounts for the experi￾mentally observed difference (Table 2). In summary, substantially enhanced hardening kinetics were found during EPA of a nanostructured Al produced by heavy cold-rolling. Microstructural characterizations suggest that this behavior may be attributed to an en￾hanced rate of dislocation recovery caused by the greatly increased mobility of dislocations. This finding shows the potential of using EPA to manipulate the microstructure of materials over a wider spectrum, by substantially increasing the recovery rate of dislocation structures with little coarsening in the structural scale. This potential facil￾itates not only the quantitative separation of the contribu￾tions of different strengthening mechanisms to the mechanical strength, but also the development of new strategies for optimization of the structure and properties of nanostructured materials. This work was supported by General Motors China Science Lab. Y.S. also acknowledges financial support from the National Science Foundation of China through projects 50601018, 50971090 and 50890174. X.H. was supported by the Danish National Research Foundation through the Danish–Chinese Center for Nanometals. The authors thank A. Godfrey and N. Hansen for helpful discussions. Y.S. thank A.H.W. Ngan for a visiting position in the University of Hong Kong and helpful discussions. [1] X. 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Malard, J. Pilch, P. Sittner, D. Schryvers, Acta Mater. 58 (2010) 4503. [28] G. Winther, X. Huang, A. Godfrey, N. Hansen, Acta Mater. 52 (2004) 4437. [29] Q. Liu, Ultramicroscopy 60 (1995) 81. [30] M. Cabibbo, W. Blum, E. Evangelista, M.E. Kassner, M.A. Meyers, Metall. Mater. Trans. A 39 (2008) 181. [31] N. Hansen, Acta Metall. 25 (1977) 863. Table 2. Boundary spacing, dislocation density and mechanical properties of the as-cold-rolled state and after EPA or AFA under peak hardening conditions. Sample rYS (MPa) rUTS (MPa) euniform (%) etotal (%) q0 (1013 m2 ) DLB (nm) DICB (nm) DR (nm) As-cold-rolled 167.1 188.2 2.0 5.9 8 280 600 345 AFA for 30 min/150 C 179.7 194.4 1.5 4.5 5 356 680 449 EPA for 2 min/(120–150 C) 190.2 206.2 1.5 4.8 5 310 607 351 Note: DLB and DICB are the distances between the lamellar boundaries and interconnecting boundaries, DR is the average grain size along a random test line. The standard deviations of these spacing measurements are 5–15 nm. q0 is the dislocation density in the volumes between boundaries (standard deviation is 0.6–1.0 1013 m2 ). Figure 4. TEM images showing the microstructural features and dislocation configurations in nanostructured Al processed by cold rolling to evM = 4.85 followed by (a) AFA at 150 C for 30 min and (b) EPA for 2 min. 150 W. Zeng et al. / Scripta Materialia 66 (2012) 147–150