Microstructural evolution and mechanical properties of deep cryogenic treated Cu–Al–Si alloy fabricated by Cold Metal Transfer (CMT) process

Journal Pre-proof Microstructural evolution and mechanical properties of deep cryogenic treated Cu–Al– Si alloy fabricated by Cold Metal Transfer (CMT) process Kun Liu, Xizhang Chen, Qingkai Shen, Zengxi Pan, R. Arvind Singh, S. Jayalakshmi, Sergey Konovalov PII:

S1044-5803(19)32310-1

DOI:

https://doi.org/10.1016/j.matchar.2019.110011

Reference:

MTL 110011

To appear in:

Materials Characterization

Received Date: 24 August 2019 Revised Date:

9 November 2019

Accepted Date: 9 November 2019

Please cite this article as: K. Liu, X. Chen, Q. Shen, Z. Pan, R.A. Singh, S. Jayalakshmi, S. Konovalov, Microstructural evolution and mechanical properties of deep cryogenic treated Cu–Al–Si alloy fabricated by Cold Metal Transfer (CMT) process, Materials Characterization (2019), doi: https://doi.org/10.1016/ j.matchar.2019.110011. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2019 Published by Elsevier Inc.

Microstructural Evolution and Mechanical Properties of Deep Cryogenic Treated Cu-Al-Si alloy Fabricated by Cold Metal Transfer (CMT) Process Kun Liua, Xizhang Chenab*, Qingkai Shena, Zengxi Panc, R. Arvind Singha, S. Jayalakshmia, Sergey Konovalovab a

School of Mechanical and Electrical Engineering, Wenzhou University, Wenzhou, 325035, Zhejiang, China b

Department of Metals Technology and Aviation Materials, Samara National Research University, Samara, 443086, Russia c

School of Mechanical, Material, Mechatronic and Biomedical Engineering, University of Wollongong, NSW, 2522, Australia

*Corresponding Author, E-mail: [email protected] or [email protected]

Abstract Cryogenic treatment is an effective route to obtain fine grained microstructure in materials. In this work, the effect of deep cryogenic treatment (DCT) on the microstructural evolution and mechanical properties of a copper-aluminum-silicon (Cu-Al-Si) alloy deposited by cold metal transfer (CMT) technology was investigated. Copper-rich Cu-Al-Si alloy with ~8.3% aluminum content was deposited on T2 copper substrate using an innovative dual wire feed system based on CMT technology. The deposited alloy was subjected to DCT for varying time duration (0, 6, 12 and 24 hours). Evaluation of microhardness and tensile properties showed that the deep cryogenic treatment enhanced these properties. Microhardness increased about 10% with the increase in the DCT time duration. Under tensile loading conditions, the DCT conducted for 12 hours provided the best properties, with an increase in both the tensile yield and ultimate strengths (~20% increase). Microstructural evolution and grain orientation were investigated using Electron Back Scattered Diffraction (EBSD). Microstructural studies revealed the formation of subgrains, fine grains and deformation twins due to DCT. Significant grain refinement was achieved due to the increase in the formation of high angle grain boundaries. It is summarized that the deep cryogenic treatment induces microstructural evolution: (i) grain refinement and (ii) texture randomization, and improves the mechanical properties of Cu-Al-Si alloy. Keywords: Cu-Al alloy; Deep cryogenic treatment; Microstructure; Mechanical properties; Electron Backscatter Diffraction (EBSD); Cold Metal Transfer (CMT)

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1. Introduction With the unprecedented development of modern technology and industry, the demand for advanced engineering materials having exceptional mechanical properties is on the rise[1]. Amongst the non-ferrous alloys, copper-aluminum (Cu-Al) alloys are widely used in industries, especially as heat exchangers in thermal power plant, cryogenic tanks for energy storage, fuselage or shells for space vehicles, etc.[2, 3]. These applications require high strength, good thermal conductivity and thermal stability. Conventionally, Cu-Al alloys are fabricated using casting process and powder metallurgy[4]. In these conventional processes, grain growth cannot be controlled effectively. In a conventional process such as casting, the low rate of bulk solidification results in variation in microstructure within the casting (i.e. formation of dendrites, columnar grains and segregation). Such variation in the microstructure is undesirable as it causes variation in the properties within the cast sample, which would affect the performance adversely. In contrast, the CMT process involves layer-by-layer deposition of metal, which has high solidification rate (i.e. heat extraction rate) and thereby the microstructure of the CMT deposited samples have relatively fine and uniform grain distribution and thus is expected to enhance the performance of the material beneficially. In the present work, cold metal transfer (CMT) technology developed by Fronius [5, 6] was used to produce a Cu-Al-Si alloy. CMT is a type of wire arc additive manufacturing (WAAM) technology with low heat input, low splash and high wire deposition rate, and can improve the performance of Cu-Al-Si alloy structural parts. The difference between CMT and traditional welding method is that in the CMT process the metal droplet is short-circuited. The metal droplets fall off because of the mechanical movement of wire in a unique way of “cold transition”, which does not produce large splashes. We recently reported the fabrication of aluminum alloy [7, 8] and copper-aluminum alloy [9] by CMT technology, and identified that the fabricated alloys had better mechanical performance than the conventional cast alloy. Cong et al. [10] investigated the effect of CMT arc mode on the porosity characteristics of Al-6.3%Cu alloy and found that the porosity can be reduced with decrease of CMT heat input. During processing of metal alloys, evolution of grain size and preferred texture are the two important microstructural parameters that strongly influence the mechanical/physical properties such as tensile strength, strain hardening, ductility and thermal conductivity [11, 12]. In the present work, a Cu-Al-Si alloy was prepared by CMT technology and was subjected to low temperature treatment. Low temperature treatment can be classified into two types: (i) cold treatment (CT, wherein, samples are exposed to ~193 K) and (ii) deep cryogenic treatment (DCT, wherein, samples are exposed to liquid nitrogen temperature ~77 K) [13]. DCT is an ultra-low temperature treatment that can induce microstructural changes and thereby can alter material properties [14]. Previous studies have reported that DCT can effectively improve the mechanical properties and service life of steels [15, 16]. Liu et al. [17] reported that after DCT for 6 hours for T8A steel, the hardness, impact 2

toughness and wear performance improved significantly. DCT process has several advantages, such as: (a) it is relatively an easy process, (b) uses nitrogen (the production of which is green and environmental friendly) and (c) suitable for large-scale industrial processing. The DCT of ferrous metals has been studied in detail [15-20], while very little information is available in literature about DCT of non-ferrous metals. In the DCT process, holding time and freezing temperature are the two main factors [21]. Contradicting views have been reported earlier on DCT holding time. While some researchers consider that DCT for short duration (~1 to 2 hours) provides better results [22], others claim that longer DCT holding time results in better properties (~24 hours) [23]. In industrial applications, the overall processing time is particularly important to minimize the total cost. In the present work, a Cu-Al-Si alloy was prepared by CMT technology, and was subjected to deep cryogenic treatment in order to enhance the mechanical properties. Investigation was focused to understand the effect of varying DCT holding time on the evolution of microstructure and mechanical properties of the CMT Cu-Al-Si alloy. To the authors’ knowledge, the current work is the first of its kind on DCT of CMT fabricated materials. 2. Materials and Methods Copper wire ERCuSi28L and aluminum wire ER4043 with diameter of 1.2 mm were used as the deposition wires, and T2 copper (Cu-T2, Chinese designation of electrolytic tough pitch (ETP) Copper) of 25 mm × 20 mm × 3 mm was selected as the substrate. ERCuSi28L is the code for copper wire (as per American Welding Society, AWS), where ‘ER’ stands for ‘Electric Rod’ (i.e. ‘E’ indicates ‘welding material’ and ‘R’ indicates ‘solid’). ER4043 is the code for welding rod having aluminum with 5% Silicon. The chemical compositions of the substrate and the wires are listed in Table 1. An innovative dual wire CMT unit was built to deposit copper-rich Cu-Al alloy with ~8.3% aluminum content, as shown in Fig. 1. In recent years, research is focused towards single phase (i.e. face-centered cubic (f.c.c) α-phase) Cu-Al alloys [24]. According to the binary phase diagram, when Al content is lower than 16.8 at. %, both copper and aluminum should be α-phase f.c.c structure [25, 26]. In the present work, the concentration of aluminum in the deposition of the wall part was obtained by using the equation (2) [9]. When calculated, the aluminum content in the designed alloy was approximately 8.6 at. Pct. Metallurgically, this aluminum content will contribute to the formation of twins, which will likely improve the mechanical properties [27]. To obtain this desired aluminum content, the feed speed of ERCuSi28L and ER4043 wires were set as 5.2 m/min and 1.2 m/min, respectively. Wt ( Al ) % =

27 × 94% v ER 4043 × 100% （1） 27 × 94% v ER 4043 + 28 × 5% v ER 4043 + 64 × 95% v ERCuSi 28 L + 28 × 3% v ERCuSi 28 L

where, Wt (Al) is the concentration of aluminum in the deposition of the wall part, vER4043 and vERCuSi28L are the feed speeds of ER4043 and ERCuSi28L, respectively. 3

The atomic weight of Al, Si and Cu are 27, 28 and 64, respectively. It should be noted that while the total Si content from the wires is 7.96% (ERCuSi28L, Si 2.89%; ER4043, Si 5.07%), the actual Si content in the CMT deposited Cu-Al-Si alloy is 1.71%. This content of Si was estimated by using equation (2). Wt ( Si ) % =

28 × 5% v ER 4043 + 28 × 3% v ERCuSi 28 L × 100% （2） 27 × 94% v ER 4043 + 28 × 5% v ER 4043 + 64 × 95% v ERCuSi 28 L + 28 × 3% v ERCuSi 28 L

where, Wt (Si) is the concentration of silicon in the deposition of the wall part, vER4043 and vERCuSi28L are the feed speeds of ER4043 and ERCuSi28L, respectively. The atomic weight of Al, Si and Cu are 27, 28 and 64, respectively. During the CMT process, the key experimental parameters are the deposition current, deposition voltage, wire feed speed and travel speed of the robotic arm. The set values of these parameters are given in Table 2. The contact tip-to-work distance (CTWD) was kept at 17 mm. The sample size prepared by CMT is 65 mm x 30 mm x 20 mm. The deposited Cu-Al-Si alloy samples were subjected to DCT at different holding times: 0 (as-deposited), 6, 12 and 24 hours. Usually, the cooling rate is kept low (2 K/min) to prevent thermal cracking. In this work, the alloy samples were slowly cooled to 77 K, held at that temperature for the designated time duration, and gradually warmed to the room temperature. Vickers microhardness measurement was conducted along the deposition height (at 1 mm step interval) at the load of 4.9 N for 15 s. The tensile test of DCT samples was conducted at room temperature at the strain rate of 1 mm/min according to ASTM: E8 standard, and three specimens were tested to acquire an average value. For Electron Back Scattered Diffraction (EBSD) analyses, samples were metallurgically polished followed by electro-polishing. EBSD analysis was conducted at an acceleration voltage of 8 kV with scan step size of 1.2 µm. Analyses were made on the CMT deposited plane, containing building direction (BD) and transverse direction (TD) in the middle section of the samples.

Fig. 1. Schematic of the dual wire cold metal transfer deposition unit.

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Table 1. Chemical compositions of the substrate and the deposition wires (wt. %).

Elements

Si

Fe

Mg

Zn

Mn

Ti

Bi

Pb

S

Al

Cu

Cu-T2 ERCuSi28L ER4043

-2.89 5.07

-0.13 0.15

--0.10

--0.68

-0.82 0.02

--0.05

0.002 ---

0.005 <0.01 --

0.005 0.002 --

-0.06 Bal.

Bal. Bal. <0.01

Table 2. Key process parameters. Parameter

value

unit

Deposition current (I) Deposition voltage (U) ERCuSi28L feed speed (VERCuSi28L) ER4043 feed speed (V ER4043) Travel speed (TS) Interlayer dwell time

112 11.1 5.2 1.2 0.35 60

A V m/min m/min m/min s

3. Results and Discussion 3.1 Effect of DCT on the Mechanical Properties of CMT Deposited Cu-Al-Si alloy 3.1.1 Microhardness The microhardness values of the Cu-Al-Si alloy subjected to DCT at various holding times are shown in Fig. 2. The measurements were taken along the central line of the deposited wall. Slight fluctuation in the microhardness value occurs with the increase in the deposition height. Particularly, at the height of about 4 mm above the fusion line there is a decrease in the values. This area corresponds to the inter-layer between the first layer and the second layer. When a new layer is deposited, the previously deposited layers would undergo re-melting and re-heating. Further, intermetallic compounds (IMCs) such as Cu9Al4 and CuAl2 form in the alloy, as reported in our earlier work [9]. The fluctuations in the microhardness values are due to the presence of inter-layer and the formation of IMCs. Microhardness values of the DCT samples are higher compared to those of the as-deposited sample. When compared with the average microhardness of the as-deposited sample (220 HV), the average microhardness values after 6 and 12 hours DCT are 236.7 HV and 243.4 HV respectively, i.e. an increase by 6.72% and 9.73%. After 24 hours DCT, the microhardness value tends to be more stable (when compared to that at 12 hours), but the average value is slightly lower than that of the 12 hours DCT sample. The average microhardness value was 242.7 HV at 24 hours, which is 9.42% higher than that of the as-deposited sample. The as-deposited sample shows large fluctuations in the microhardness value when compared to the DCT samples. The DCT process: (i) stabilizes the microhardness values, and the stability increases with the increase in the 5

holding time and (ii) increases the microhardness values when compared to the untreated sample. 280 0h 12h

260

6h 24h

Hardness (HV)

240 220 Average hardness values

200

0h 6h 12 h 24 h

180 160 0

2

4

6

8

10

12

14

221.8 HV 236.7 HV 243.4 HV 242.7 HV 16

18

20

Deposition Height (mm)

Fig. 2. Microhardness values of the CMT deposited Cu-Al alloy after DCT at various holding times. 3.1.2 Tensile properties The tensile properties of the as-deposited Cu-Al-Si alloy and those after DCT at different holding time are shown in Fig.3. The DCT samples have higher values of ultimate tensile strength (UTS), yield strength (YS) and elongation (EL). Upto a certain holding time (12 hours), there is a linear relationship between the tensile properties and the DCT holding time. Similar behavior was observed earlier in austenitic stainless steel [28]. The UTS and YS of the DCT sample with 12 hours hold duration increased by 20.3% and 19.9% respectively, when compared to the as-deposited sample. For the DCT sample with 24 hours hold duration, the UTS and YS decreased slightly when compared to those of the DCT sample with 12 hours hold duration; but were higher when compared to those of the as-deposited sample (by 11.3% and 8.4%). The increase in strength after DCT is possibly due to strengthening by fine grains, and is discussed in detail in Section 3.2. Rangaraju et al. [29] reported a similar behavior during deep cooling of aluminum alloys, which promoted grain refinement effectively. Further, due to the difference in the expansion coefficients of copper, aluminum and intermetallic compounds (Cu9Al4, CuAl2 [9]), during the DCT process, the degree of shrinkage of each crystal phase is different [30]. It has been reported earlier [5,35] that, due to the different degree of shrinkage of each crystal plane, the deformation of each crystal grain (or phase) would be restricted by the surrounding crystal grains/phase, thereby some grains undergo deformation resulting 6

600

500

UTS, YS (MPa)

400 300 200 100 0 as deposited

6 12 DCT holding time (h)

20 UTS 19 YS Elongation 18 17 16 15 14 13 12 11 10 9 8 7 6 5 24

Elongation (%)

in microplasticity of the material, which contributes towards the improvement in strength and hardness. It is evident from the mechanical properties evaluation that DCT process improves the hardness and strength of the deposited alloy.

Fig. 3. Tensile properties of the CMT deposited Cu-Al-Si alloy after DCT at various holding times. 3.1.3 Tensile Fracture Morphology High magnification scanning electron microscope images from the center of the fracture surfaces of the tensile tested samples are shown in Fig. 4. It can be seen that all samples exhibit typical ductile fracture features. The fracture surfaces of DCT samples exhibit deeper dimples, which are indicative of ductile fracture with good plasticity, i.e., improved strength and ductility.

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Fig. 4. Tensile fracture surface morphologies of CMT deposited Cu-Al-Si alloy: (a) as-deposited, (b) after 6 hours DCT, (c) after 12 hours DCT and (d) after 24 hours DCT. 3.2 Effect of DCT on the Microstructural Evolution of CMT Deposited Cu-Al Alloy 3.2.1 Grain morphology and size The results from the EBSD analyses on the microstructure of both the as-deposited and DCT samples are given in Fig.5. Orientation imaging maps (OIM) were used to study the orientation of grains based on inverse pole figure (IPF). From these maps (Figs.5 (a), (c), (e), (g)), the observed microstructural features are: (i) subgrains, (ii) deformation twins and (iii) fine grains. Discussion on these observed microstructural features is presented below. (a) Formation of subgrains Subgrains appear in the microstructure of Cu-Al-Si alloy after DCT process, which contribute to grain refinement and increase the grain boundary area. Defects such as impurities, vacancies and dislocations exist at the grain boundaries [31]. During the DCT process, vacancies supersaturate at low temperature, and high-density dislocations are generated [32-34]. During the subsequent temperature recovery process, dislocations propagate, interact, tangle with each other and rearrange spatially to form various sub-structures, such as subgrains [35]. (b) Formation of deformation twins Earlier studies have reported that the formation of deformation twins in Cu-Al alloy is 8

due to Al in the solid solution, which would reduce the stacking fault energy of Cu [36, 37]. In addition, such decrease in stacking fault energy reduces mobility and energy of grain boundaries, which can suppress the migration and annihilation of grain boundaries, thereby enhancing the microstructure stability [38]. Wang et al. [39] reported that for materials with low stacking fault energy, deformation twins will play a significant role in plastic deformation and grain refinement. Further, since Cu-Al-Si alloys belong to the face-centered cubic (fcc) structure, twins occur when the deformation speed is either extremely fast or when the temperature is very low [40]. The formation of deformation twins observed in the DCT samples in this work (Fig.5) is attributed to: (i) Al as solid solution in Cu, which would reduce the stacking fault energy of Cu at low temperature and (ii) the low DCT temperature (77 K) that would provide conditions for the formation of twins and would also cause shrinkage (that generates large stresses inside the alloy) [30]. This makes atomic migration difficult at low temperatures, resulting in twins [41]. (c) Grain refinement At low temperatures, in metallic alloys, the density of defects is high due to the increase in dislocation pile-up and grain boundaries, which causes grain refinement [42]. Dislocation substructures would develop within the grains by dislocation rearrangement, accompanied by the formation of finer subgrains with high angle grain boundaries [35]. The as-deposited Cu-Al-Si alloy in this work exhibits coarse grains with size ~19.19 µm (Figs. 5 (a) and (b)). When subjected to DCT and with the increase in DCT holding time, the grains become refined. For quantitative analysis, the average grain size was estimated using line intercept method for all the specimens. Grain size was quantified by the determination of the area of each grain and the calculation of its circle-equivalent diameter. The average grain size of the DCT samples are shown in Figs. 5 (b), (d), (f) and (h). In the 12 hours DCT sample, grains become more homogenized with the grain size effectively reducing to ~8.49 µm (this sample has the smallest grain size amongst all the DCT samples). When the DCT holding time was extended to 24 hours, the grain structure did not show any significant refinement, rather a slight increase in size (~10 µm) was observed, which corroborates well with the decrease in mechanical properties observed (Figs. 2 and 3). The 12 hours DCT sample showed the best mechanical properties (Figs. 2 and 3), and this emphasizes the effect of grain size on the strengthening.

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Fig. 5. Orientation image micrographs and grain size distribution of the CMT deposited Cu-Al-Si alloys, subjected to varying DCT holding times: (a) and (b) as-deposited, (c) and (d) after 6 hours DCT, (e) and (f) after 12 hours DCT, (g) and (h) after 24 hours DCT. 10

3.2.2 Evolution of grain boundary structure during DCT In grain boundary engineering, increasing or decreasing the ratio of specific types of grain boundaries is adopted to optimize the properties of the final material. To understand the evolution of grains boundaries due to DCT of Cu-Al-Si alloy, the grain boundary and misorientation angle distribution were examined (Fig.6, the same frame of the OIM map in Fig.5.). Grain boundaries with misorientations higher than 15° are defined as high-angle grain boundaries (HAGBs) and those with misorientations between 2~15° are considered as low-angle grain boundaries (LAGBs) [43, 44]. Misorientations of less than 2° were ignored because of the scanning resolution of EBSD technique. The EBSD analyses shows significant grain boundary evolution. The new refined grains can be observed in the grain boundary map consisting of HAGBs of higher than 15° (in black line) and LAGBs of 2~5° (in blue line) and 5~15° (in green line), as shown in Figs. 6 (a), (c), (e) and (g). When compared to the as-deposited sample, in the DCT samples the refined grains form a new orientation. The grain misorientation angle distribution of the Cu-Al-Si alloys gradually shifts towards high misorientation angle in the sequence: as-deposited → 6 hours → 12 hours → 24 hours, as observed in Figs. 6 (b), (d), (f) and (h). The 12 hours DCT sample has the largest fraction of high misorientation angles. Fig.7 shows the frequency of misorientation angle of LAGBs and HAGBs at different DCT holding time of Cu-Al-Si alloy. The frequency of the misorientation angle > 15° for the samples subjected to varying DCT holding time are 95.2%, 96.6%, 98.0% and 97.7%, respectively. Further, it can be seen from this figure that both LAGBs and HAGBs are present in the as-deposited alloy. During DCT, as the temperature drops from ambient to liquid nitrogen temperature, the material experiences thermal shrinkage and volume change [45]. Due to shrinkage strain, the stored deformation energy acts as the driving force for the generation and mobility of dislocations, which eventually leads to dislocation interactions [46]. The accumulation and rearrangement of dislocations causes LAGBs to transform into HAGBs, such that with the increase in DCT holding time, HAGBs increase at the expense of LAGB fractions. To note, the deviation of the uncorrelated curve from the random curve is related to the microstructure of the refined grains. Especially, after DCT for 6 hours, there is a good similarity between the uncorrelated curve and the random curve (Fig. 6 (d)), which suggests that the microstructure of the DCT sample possess a random texture. Figs. 6(d)(f)(h) show the misorientation angle distribution peaks at 45°, which indicate that grain orientations of all DCT samples are completely random [47]. As mentioned earlier, DCT process results in the formation of deformation twins, which play a significant role in grain refinement [39]. The new grains formed due to DCT (suppression of grain growth) have different orientations and hence the texture is random.

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Fig. 6. Grain boundary and misorientation angle distribution of CMT deposited Cu-Al-Si alloy subjected to DCT at various holding times: (a) and (b) as-deposited, (c) and (d) after 6 hours DCT, (e) and (f) after 12 hours DCT, (g) and (h) after 24 hours DCT. 12

Fig. 7. Frequency of misorientation angle of LAGBs and HAGBs at different DCT holding time of Cu-Al-Si alloy. 4. Strengthening Mechanisms The main reason for the improvement in mechanical properties of the Cu-Al-Si alloy subjected to DCT is the microstructural evolution induced by the DCT process. The changes in the microstructure due to DCT are characterized by: (i) fine grains (i.e. grain refinement) and (ii) difference in the orientation of grains (i.e. misorientation angle and distribution). Fig. 8 shows a schematic illustration of grain refinement and grain orientation in the CMT deposited Cu-Al-Si alloy when subjected to DCT. In the figure, the grain sizes of the as-deposited sample are represented by d1 and d2, respectively and the grain sizes after DCT are represented by d3, d4, d5, and d6. Due to DCT: (a) the grains get significantly refined, i.e., larger grains (d1, d2) are deformed into smaller grains (d3, d4, d5, d6). Strengthening occurs due to Hall-Petch effect (i.e. finer the grains, higher is the yield strength) [40]. The grain refinement of Cu-Al-Si alloy leads to an increase in the grain boundary area. Therefore, the ability to resist deformation is enhanced [39, 48], which gives rise to improved strength. (b) The newly formed grains have different orientations (i.e. random texture). Thus, the enhancement in the strength and ductility of the samples subjected to DCT is attributed to the formation of fine grains and random texture, which would increase the resistance to crack propagation, resulting in higher fracture stress [49]. The highest strength and highest ductility at the DCT of 12 hours are due to: (i) the smallest grain size and (ii) texture randomization obtained at this DCT hold time.

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Fig. 8. Schematic illustration of grain refinement and grain orientation in the CMT deposited Cu-Al-Si alloy when subjected to DCT. 5. Conclusions In this work, cold metal transfer technology (CMT) was used to develop Cu-Al-Si alloy and subjected to deep cryogenic treatment (DCT) at different holding times. The effect of DCT on microstructural evolution and mechanical properties of the alloy were investigated. The following are the conclusions that could be drawn from the present investigation: (1) DCT increases hardness and tensile strength properties of the alloy. (2) 12 hours DCT duration was identified to be the optimal time to obtain the highest mechanical properties. (3) DCT induces microstructural changes, such as: (i) formation of subgrains, deformation twins and fine grains, and (ii) texture randomization. (4) DCT induced texture randomization is due to increased dislocation generation, interaction between dislocations and subsequent dislocation rearrangement that form new grains with different orientations. (5) DCT increases both strength and ductility due to: (i) grain refinement and (ii) texture randomization. Acknowledgements This work was sponsored by the National Natural Science Foundation of China under Grant No. 51575401 and Grant No. 51975419.

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Highlights • Cu-8.3Al alloy was fabricated using dual wire cold metal transfer (CMT) process. • The alloy was subjected to deep cryogenic treatment (DCT) at different hold times. • Microhardness and tensile properties of the alloy improved significantly after DCT. • EBSD analyses of DCT samples showed fine grains, subgrains and deformation twins. • Grain refinement and texture randomization enhanced the mechanical properties.

Conflict of Interest The paper is an original work and has not been considered for publication elsewhere. There is no conflict of interest for the manuscript submission, and the manuscript has been approved by all authors for publication.