Effect of annealing on microhardness and electrical resistivity of nanostructured SPD aluminium

Accepted Manuscript Effect of annealing on microhardness and electrical resistivity of nanostructured SPD aluminium A.M. Mavlyutov, A.S. Bondarenko, M.Yu. Murashkin, E.V. Boltynjuk, R.Z. Valiev, T.S. Orlova PII:

S0925-8388(16)34166-4

DOI:

10.1016/j.jallcom.2016.12.240

Reference:

JALCOM 40162

To appear in:

Journal of Alloys and Compounds

Received Date: 27 October 2016 Revised Date:

12 December 2016

Accepted Date: 18 December 2016

Please cite this article as: A.M. Mavlyutov, A.S. Bondarenko, M.Y. Murashkin, E.V. Boltynjuk, R.Z. Valiev, T.S. Orlova, Effect of annealing on microhardness and electrical resistivity of nanostructured SPD aluminium, Journal of Alloys and Compounds (2017), doi: 10.1016/j.jallcom.2016.12.240. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Effect of annealing on microhardness and electrical resistivity of nanostructured SPD ACCEPTED MANUSCRIPT aluminium A.M. Mavlyutova*, A.S. Bondarenkob, M.Yu. Murashkinb,c, E.V. Boltynjukb, R.Z. Valievb,c, T.S. Orlovaa,d a

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St. Petersburg National Research University of Information Technologies, Mechanics and Optics, Kronverksky Pr. 49, St. Petersburg 197101 Russia b St. Petersburg State University, Universitetskiy Pr. 28, St. Petersburg 198504 Russia c Institute of Physics of Advanced Materials, Ufa State Aviation Technical University, K. Marx str. 12, Ufa 450000 Russia d Ioffe Institute, Politekhnicheskaya ul. 26, St. Petersburg 194021 Russia

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[email protected], [email protected], [email protected], [email protected], [email protected], [email protected] *

Corresponding author, e-mail: [email protected]

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Keywords: metals and alloys, nanostructured materials, grain boundaries, electrical transport, mechanical properties, microstructure.

Abstract

The influence of microstructure evolution on microhardness and electrical resistivity

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of ultrafine grained (UFG) commercial purity Al under annealing at different temperatures within a range of 363 – 673 K was studied. The initially coarse grained Al was processed by high pressure torsion (HPT) technique for the formation of UFG structure. The microstructure was characterized by electron backscattering diffraction and X-Ray diffraction. It was shown

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that annealing of UFG Al at temperatures within a range of 363 – 473 K leads to simultaneous increase of microhardness (by 6 – 13%) and electrical conductivity (by 4 – 8% at 300 K). The correlation between microstructural features and the resulting properties were analyzed. The

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average width s of potential barriers at grain boundaries (GBs) in HPT-processed Al was estimated in the frame of a tunnel model. The obtained large value of s compared with the GB crystallographic width is associated with elastically distorted lattice near GBs. The obtained results suggest a new way to increase simultaneously strength and electrical conductivity of UFG Al alloys by an appropriate annealing.

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1. Introduction

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In recent years Al alloys in nanostructured and ultrafine grained (UFG) states are of great interest due to their unique mechanical and physical properties [1-4]. Electrical conductivity of pure Al is 62% IACS (International Annealed Copper Standard), but due to lower density, Al has about twice higher conductivity per weight unit than Cu [5]. Considering the lower density and lower cost of Al compared with Cu, Al may be regarded as

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a very promising material to develop high-strength conductors for electrical power lines.

However, the main disadvantage of Al is its low mechanical strength. Conventional strengthening methods, including alloying and ageing, lead to dramatic decrease in electrical conductivity due to its strong dependence on the induced microstructural changes [6].

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Achievement of high strength and good electrical conductivity for Al and Al-based alloys is a challenge and finding a solution for this problem will contribute to the increase of energy

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transmission efficiency and decrease of the power lines cost. Recently, a good combination of high strength and electrical conductivity has been demonstrated for Al-Mg-Si alloys processed by severe plastic deformation (SPD) in two stages: at room temperature (RT) and at elevated temperatures [7-9]. Refinement of grains down to ultrafine size by various SPD techniques at a temperature <0.4Tm (Tm is the melting temperature) makes it possible to achieve high strength [1,3,4]. SPD processing at higher (elevated) temperatures increases

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electrical conductivity in the UFG alloy through dynamic ageing, which results in clearing the Al matrix from dissolved impurity atoms by the formation of nanosized secondary phase precipitates [7,8]. However, the ways to further increase strength while keeping or even increasing the high level of electrical conductivity are still of current interest. For example,

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introducing an additional dislocation density in the UFG structure with preserving other microstructural parameters leads to enhanced mechanical strength while keeping high

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electrical conductivity of the material almost unaffected [10]. As was shown recently, strength of UFG Al can be increased by appropriate annealing. In [11,12] the influence of annealing on microstructure and strength of commercial purity (CP) Al and high-purity Al processed by accumulative roll bonding (ARB) was studied. It was demonstrated that annealing at a temperature of the range 423 – 523 K during 0.5 – 1 h leads to the anomalous increase in strength. However, there are controversies on published issues relating to the effect of annealing on the strength of UFG Al. For example, in [13] it was shown that annealing of UFG Al processed by rotary swaging does not lead to the increase in microhardness throughout the annealing temperature range from 373 to 723 K.

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The aim of this study is to investigate the possibility of simultaneous increase in

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strength and electrical conductivity of HPT-processed UFG CP Al by annealing and to identify the key microstructural parameters responsible for improvement of these functional properties.

2. Material and experimental procedures

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Commercial purity Al (99.5 wt.%) was chosen as the material of interest for this study because the effects of its solute atoms and precipitates are assumed to be negligible. Initially coarse grained (CG) disks of the diameter 20 mm and thickness 2 mm were subjected to SPD processing by high pressure torsion (HPT) [3] under a hydrostatic pressure of 6 GPa to 10

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revolutions at RT for the formation of UFG structure. After HPT processing the true strain at the distance of 5 mm from disk centre was γ~6.6 [3]. The HPT-processed samples were

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annealed at various temperatures in the range 363 – 673 K for 1 h to obtain a variety of microstructural parameters (grain size, grain boundary misorientation angle, dislocation density and others). Hereafter, the HPT-processed samples without annealing are referred to as Al_RT, the samples with subsequent annealing at Tan are referred to as Al_Tan (for example Al_363, Al_403 and so on).

Microstructural characterization was performed by electron backscattering diffraction

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(EBSD) analysis using the scanning electron microscope Zeiss Merlin. The samples for EBSD studies were prepared by conventional metallographic techniques consisting of polishing with the diamond and colloidal-silica suspensions. EBSD mapping was performed on a scan area of 32.6х24.4 µm2 with a scan step of 0.2 µm and over 1000 grains were

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analysed for each UFG sample. Seven Kikuchi bands were used for indexing the diffraction patterns. Distributions of grains on size and grain boundaries (GBs) between the adjacent

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grains on their misorientation angle θ were determined from the EBSD maps. GBs with θ≤15o were referred to the low-angle grain boundaries (LAGBs) and GBs with θ>15o were considered as the high-angle grain boundaries (HAGBs). The grain-reconstruction method was applied to determine the average grain size [14]. Vickers microhardness HV was measured using a Shimadzu HMV-G microindentation tester with a load of 1 N for 15 s. For reliable results, 3 – 4 samples were measured for each Tan and each sample was measured 15 times. For electrical resistivity measurements the samples were cut from HPT disks according to the scheme in Fig. 1. Electrical resistivity ρ was measured by standard four-probe technique. Electrical resistivity was measured at RT, at 77 K in steady-state liquid nitrogen 3

and at an intermediate temperature Tin around 200 K, which was achieved by natural

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evaporation of liquid nitrogen. The temperature of a sample was controlled by a silicon diode placed near the sample. The accuracy of temperature measurements was ±0.03 K. The obtained values of ρ77, ρTin and ρRT fitted well a linear dependence. To compare the resistivity of all studied samples at 300 K, the value ρ300 for each sample was determined from this linear approximation. Electrical resistivity measurements are described in more details in

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[15].

3. Results and discussion

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3.1. Microstructure evolution

Fig. 2 presents the typical EBSD maps of HPT-processed samples before and after

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annealing at different temperatures. The distributions of grains on size and GBs on the misorientation angle are shown in Fig. 3 and 4, respectively. Microstructural parameters obtained from the EBSD maps (Fig. 2) are shown in Table 1.The average grain size (dav) and the standard deviation (∆d) that characterizes the width of size distribution of grains are given for each sample. Amount of grains (in %) having the size below 1 µm, between 1 and 2 µm and above 2 µm (further denoted as f<1, f1–2 and f>2) and amount of GBs (in %) with the

Table 1.

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misorientation angle ≤15o and >15o (further denoted as θ≤15 and θ>15) are also shown in HPT processing transforms the initial CG structure into the UFG structure with equiaxed grains (Fig. 2a) having an average grain size dav≈810 nm, about 73% of grains

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having the size below 1µm (f<1≈73%). Most of GBs (77%) are HAGBs (Table 1). Annealing at 363 K does not lead to significant changes in the average grain size, it is equal to ~850 nm,

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but the value of f<1 slightly decreases (f<1=64%) and f1-2 increases from 26 to 34% compared to the non-annealed sample (Table 1). The portion of HAGBs does not change. The values of dav and ∆d increase gradually with the Tan growth from 403 to 473 (Table 1), the most of GBs remaining the high angle ones. Annealing at 673 K leads already to the formation of CG structure characterized by high value of dav=7 µm and very wide distribution of grains on size (Fig. 3f and Table 1). In the sample Al_673 the most of GBs are low angle. Detailed XRD analysis of the studied samples was conducted in our previous work [15] and the estimated values of dislocation density Ldis on the basis of this analysis are presented in Table 1. As was shown, the annealing at 363 К of HPT-processed samples leads to the decrease in dislocation density by ~3 times from 2·1012 to 6 – 7·1011 m-2, which 4

indicates the beginning of structure recovery process (Table 1). Annealing at a higher

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temperature of 403К only slightly reduces the dislocation density compared to that in the Al_363 sample.

3.2. Mechanical properties

HPT processing leads to a significant increase in the microhardness values of Al

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samples from 205±12 to 634±15 MPa (Fig. 5). The dependence of microhardness of HPTprocessed samples on annealing temperature is shown in Fig. 5. Annealing in a temperature range of 363 – 473 К results in the additional increase in microhardness (this part of the plot is shown in a larger scale). The maximum value of microhardness is obtained by annealing at

was observed compared to the initial CG state.

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423 К and it equals to 715±14 MPa. After annealing at 673 К, no difference in microhardness

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Microhardness of metallic materials can be determined as a superposition of the contributions of various strengthening mechanisms [12, 16]:

H V = H VAl + ∆H VGB + Σ i ∆ i H Vsol + ∆H Vdisl + ∆H Vprecip ,

(1)

where H VAl is the microhardness of non-deformed pure monocrystalline Al, ∆H VGB is the contribution of grain boundary strengthening, ∆ i H Vsol is the contribution of solid solution

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strengthening from i-th solute element, ∆H Vprecip is the contribution of precipitation strengthening, and ∆HVdisl is the contribution of dislocation strengthening. According to XRD and transmission electron microscopy (TEM) studies [15], annealing of UFG СP Al does not change the concentration of solid solution and does not

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cause the formation of second phase precipitates, but leads to the decrease in dislocation density (Table 1), hence ∆ i H Vsol and ∆H Vprecip in Eq. (1) do not change during annealing.

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Since after the annealing at Tan=363 – 473 K the dislocation density decreases, the grain size increases (or almost does not change at Tan=363 K) without the substantial change in share ratio of HAGB and LAGB (Table 1), the contributions ∆HVdisl and ∆H VGB to the strengthening of CP Al, and hence the microhardness HV, should decrease as a result of such annealing. On the contrary, the experimental data (Fig. 5) show that annealing at Tan=363 – 473 K leads to microhardness increase. It should be noted that for the first time increase in strength of the HPT-processed CP Al by its annealing is demonstrated experimentally. In [11,12] it was revealed that yield stress of ARB-processed CP Al increased after its annealing at 423 K. The authors suggested that the closely spaced high-angle boundaries in the UFG structure can serve as the sinks for dislocations during the annealing, reducing the

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number of dislocation sources. This may lead to an increase in the yield stress to activate the

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new dislocation sources during straining.

On the other hand, the influence of annealing on microhardness of rotary swaged CP Al was investigated in [13]. The value of microhardness did not change under the annealing up to 448 K. Such controversy of results in [11, 12] and [13] may be caused by specific features of UFG microstructure in these two cases. In [11, 12] the UFG microstructure of CP

microstructure is formed mainly by LAGBs (~58%).

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Al formed after ARB processing consists mainly of HAGBs, whereas in [13] the UFG

Thus, our results are in good agreement with published data in [11,12] and serve as an evidence that the strength in UFG Al can be increased by annealing, providing that the UFG structure contains mostly HAGBs regardless of a particular technique this UFG structure was

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obtained with. Appropriate annealing of such UFG structures can be an effective approach to

3.3. Electrical properties

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further strength enhancement.

Fig. 6 shows the dependences of electrical resistivities ρ77 and ρ300 of HPT-processed Al on average grain size dav achieved during annealing in ρ − d av−1 coordinates. The values of ρ77 were directly measured at 77 K. The values of ρ300 were obtained from a linear

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approximation of the measured values of ρ at 77 K, ~200 K and RT as described in Section 2. The corresponding value of Tan, at which a certain average grain size was obtained is also shown on the plots (Fig. 6). Annealing of UFG Al at 363 K leads to a substantial decrease in the electrical resistivity with an average grain size almost unchanged. As was shown in [15],

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this resistivity decrease is mainly associated with the recovery of GB state from nonequilibrium to the more equilibrium one during such annealing. Subsequent annealing at 403,

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473 and 673 K leads to further decrease in electrical resistivity. It should be noted that annealing of HPT-processed CP Al at 403 – 473 K results in the substantial increase of electrical conductivity by 6 – 9% as well as of strength (microhardness increases by 8 – 13%), which offers a new approach to simultaneously improve the strength and conductivity. To analyze the changes in resistivity during annealing, the parameter ∆exp was T an

introduced, which is equal to the change of ρ300 after annealing at Tan. Fig. 7 demonstrates the dependence of ∆exp on Tan (curve 1). On the other hand, the change in electrical resistivity can T an

be evaluated with the Matthiessen’s rule [5], according to which the microstructure parameters additively affect the value of the resistivity. As was noted above, only the average

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grain size and dislocation density change during annealing (Table 1), therefore the change of

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electrical resistivity after annealing at Tan can be estimated as: GB ∆thTan = ∆dis Tan + ∆ Tan ,

(2)

dis where ∆GB Tan and ∆Tan are contributions to the change of electrical resistivity due to decrease in

GB density (the area of grain boundaries per unit volume) and dislocation density, respectively. Considering the grains in the cubic form, the GB density can be estimated as

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S GB = 3d av−1 , then ∆TGBan = δρ GB ∆SGB , where ∆SGB is the change of the GB density of the HPTdis processed CP Al after its annealing at Tan. Similarly, ∆dis Tan = δρ ∆Ldis where ∆Ldis is the

change in the dislocation density after annealing of the HPT-processed CP Al at Tan. The values δρGB=2.6·10-16 Ωm2 and δρdis=2.7·10-25 Ωm3 which characterize the contributions from

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unit densities of GBs and dislocations, respectively, to total resistivity were earlier determined for CP Al in [17].

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The dependence of ∆thTan on Tan is also shown in Fig. 7 (curve 2). As is seen, there is a ) and theoretical estimations ( ∆thTan ). large discrepancy of experimental results ( ∆exp T an

Considering that annealing at 363 K almost does not affect dav and the decrease in resistivity occurs mainly due to the change of GB state from the non-equilibrium to the more equilibrium one [15], Eq. (2) should have extra contribution ∆noneq, associated with the non-

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equilibrium state of GBs:

GB ∆thTan* = ∆dis Tan + ∆ Tan + ∆ noneq .

(3)

For the state after annealing at 363 K we can estimate ∆noneq by following equation: (4)

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dis GB ∆ noneq = ∆exp 363 − ∆ 363 − ∆ 363 .

Assuming that annealing at higher temperatures will not change substantially the value of

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∆noneq we can estimate ∆thTan* for all states after annealing at 403, 473 and 673 K with the same value of ∆noneq (Eq. (4)). The obtained evaluations ∆thTan* according to Eq. (3) are shown in Fig. 7 (curve 3). They are in rather good agreement with the experimental results (curve 1). It should be noted that annealing at higher temperatures (403, 473 and 673 K) can slightly increase the value of ∆noneq if GBs do not reach fully the equilibrium state under annealing at 363 K. This most probably explains little difference between experimental data (curve 1) and estimated ones (curve 3) in Fig. 7. Thus, the enhancement of electrical resistivity of CP Al after HPT processing is caused not only by increased densities of GBs and dislocations, but also, to a large extent, by non-equilibrium state of the GBs.

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According to [18] GB is considered as a potential barrier with width s and height φ. If

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the energy of a moving electron is less than the height of the potential barrier, the electron can overcome the barrier via a tunneling mechanism. The width of the barrier depends on the crystallographic width of GB and also on the crystal lattice distortions and elastic strains near GB. Therefore, it is possible to judge about the grain boundary state by analyzing the potential barrier width. On the bases of the temperature dependence of electrical resistivity the potential

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barrier width can be estimated for HPT-processed Al similarly as was done for the HPTprocessed high purity Cu and Ni in [19]. Since the measured values of ρ at 77 K, ~200 K and RT fit well a linear approximation, for our estimation we took the resistivity values at three arbitrary temperatures (77, 273, 293 K) on the linear approximation (Table 2).

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The temperature dependence of electrical resistivity of GBs in the model of tunnel conductivity has the form [18]:

sin α T , αT

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0 ρ GB = ρ GB

(5)

0 where ρ GB is the GB electrical resistivity fraction independent of temperature, α is a

parameter of the model equal to:

α = 2 m e πk B s

1

h ϕ

,

(6)

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where me is the electron mass, kB is the Boltzmann constant, ћ is the Plank constant. Consequently, the temperature dependence of electrical resistivity of UFG metal has the form:

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0 ρT = ρ 0 (T ) + ρ dis + ρ GB

sin αT , αT

(7)

where, ρ dis is the contribution to the electrical resistivity from dislocation and is temperature

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independent [20] and ρ0 (T) is the electrical resistivity of the crystal lattice and is temperature dependent. This equation has three independent parameters ρGB, ρdis, α. These parameters can be determined with the experimental data given in Table 2. For the values of ρ0 at different temperatures we took the values of resistivity of the initially CG CP Al presented in Table 2 as well. Simple estimations show that since the grain size in CG CP Al is large, the contribution of the GB resistivity to the crystal lattice resistivity is negligible at RT and very small at 77 K. The value α can be used to find the product sφ-1/2 from Eq. (6). On the other hand, the relation between the values of s and φ can be determined from the Mayadas-Shatskes model [21]. The GВs in this model are represented as the potential barriers on which the conduction electrons are partially scattered during their motion in the 8

crystal. According to this model, the product of the potential barrier width s by its height φ is

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expressed by the equation:

Kh 2 vF2 (sϕ ) = , 2(1 − K ) 2

(8)

where vF is the Fermi velocity, K is the mirror reflection coefficient. In this estimation K=0.17 was used as obtained for CG CP Al in [21]. We believe that the value of K in UFG Al does not differ much from that in the CG structure, as was shown for HPT-processed high purity

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Cu [22].

The application of Eqs. (6) and (8) makes it possible to determine the average width s of the potential barrier for conduction electrons. Proceeding from the estimation, we have s≈9 nm, which is much larger than the crystallographic width of GB in metals, the latter equaling

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to 2–3 interatomic distances [23]. The obtained high value of a potential barrier width at GBs is relevant to the TEM observations, thus indicating to the presence of elastic lattice

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distortions near GBs in HPT-processed Al samples [15]. Similar evaluations of a potential barrier width at GBs were carried out for HPT-processed high purity Cu and Ni, the obtained values equaling to 2.1 and 3.7 nm, respectively [19]. According to [23], the physical width of GB in UFG Fe is 8.4 nm.

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4. Conclusions

For the first time an increase in microhardness as a result of annealing has been experimentally demonstrated and analyzed for the UFG CP Al samples processed by HPT. Annealing of UFG CP Al within a temperature range of 363 – 473 K anomalously raises the

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microhardness up to 13%, which most likely happens due to the annihilation of mobile dislocations in HAGBs leading to an increase in yield stress to activate the new dislocation

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sources.

Annealing of UFG CP Al at a temperature range of 363 – 473 K leads also to an

increase in electrical conductivity at 300 K up to ~8%. The obtained results suggest a new approach to further increase the strength and electrical conductivity of UFG Al and Al alloys processed by severe plastic deformation. In the frame of a tunnel model the width of GB potential barrier has been estimated for HPT-processed Al and it equals to ~9 nm. The large width of the GB potential barrier compared to the GB crystallographic width is associated with the elastically distorted lattice near GBs.

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Acknowledgements

EBSD measurements were performed on equipment at the Interdisciplinary Resource Centre for Nanotechnology of the Research Park of the St. Petersburg State University. M.Yu.M. and R.Z.V. acknowledge the support of the Russian Federal Ministry for

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Education and Science (RZV Grant No. 14.B25.31.0017).

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boundary width of ultrafine grained copper and nickel from the electrical resistivity measurements, Phys. Stat. Sol. A 162 (1997) 559-566.

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Figure captions

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Fig. 1 Position of sample cut from the HPT disks (annealed and non-annealed) for the electrical resistivity measurements. Arrangement of current (1, 2) and potential (3, 4) contacts on the sample (5) and position of silicon diode (6) on the holder (7).

Fig. 2 EBSD maps of CP Al after HPT processing (a) and subsequent annealing at 403 K (b),

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473 K (c) and 673 K (d).

Fig. 3 Grain size distribution in CP Al after HPT processing (a) and subsequent annealing at

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363 K (b), 403 K (c), 423 K (d), 473 K (e) and 673 K (f).

Fig. 4 GB misorientation angle distribution in CP Al after HPT processing (a) and subsequent

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annealing at 363 K (b), 403 K (c), 473 K (d) and 673 K (e).

Fig. 5 Microhardness versus annealing temperature for CP Al processed by HPT. (The insert shows the part of graph in a larger scale).

Fig. 6 Electrical resistivity at 300 K (a) and 77 K (b) versus average grain size for CP Al

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processed by HPT. The graph also shows the corresponding value of annealing temperature Tan, which a certain average grain size was obtained at. Fig. 7 Electrical resistivity change ∆ versus annealing temperature. Experimental results ( ∆exp Tan

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- curve 1), theoretical estimations without ( ∆thTan - curve 2) and with ( ∆thTan* - curve 3) allowance

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for the resistivity contribution from non-equilibrium state of GBs.

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Table 1

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Structural parameters of the HPT-processed-and-annealed samples. EBSD analysis data: dav – average grain size, ∆d – standard deviation, f<1, f1-2, f>2 - percentage of grains with size <1, 12,>2 µm, respectively; θ≤15, θ>15 – percentage of GBs with the misorientation angle ≤15 and >15 grad, respectively. Ldis – dislocation density (data of Ref. [15]).

dav, nm

∆d, nm

f<1, %

f1 – 2, %

f>2, %

θ≤15, %

θ>15, %

Ldis, m-2 [15]

Al_RT Al_363 Al_403 Al_423 Al_473 Al_673

810±9 850±11 1130±17 1160±9 1090±18 7200±525

130 150 490 425 530 4500

73 64 42 33 48 -

26 34 53 63 47 10

1 2 5 4 5 90

23 22 20 24 28 69

77 78 80 76 72 31

2·1012 7·1011 6·1011 -

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Table 2

Temperature dependence of electrical resistivity of CP Al in initial (CG) and HPT-processed (UFG) state. The values of ρ are taken from linear approximation of the measured values of ρ

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at 77 K, ~200 K and RT.

ρ77, nΩm 2.9 5.1

ρ273, nΩm 24.7 27.9

ρ293, nΩm 26.9 30.2

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State of CP Al CG UFG

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Increase of microhardness and electrical conductivity of nanostructured Al by annealing. The strengthening is associated with high angle grain boundaries in the microstructure. Contribution of grain boundary non-equilibrium state to electrical resistivity is analyzed.

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The average width of potential barriers at grain boundaries for electrons was estimated.