Abstract

This paper examines the progress in aircraft and aircraft engines from the standpoint of the role that better materials and processing has played. Such progress includes the relatively recent transformation of the aircraft industry from purely performance driven products to products that are driven by customer value. It is demonstrated that advances in materials and processing technology and understanding has enabled much of the progress that has been made since the inception of manned, heavier than air flight. The recent constraints of cost, as determined by customer value, have changed the way new materials are introduced and these trends appear to be the new paradigm for the aircraft and aircraft engine industry.

While the focus of this paper is aircraft and aircraft engines, the broader focus is on the role of materials in creating lightweight structures. There are examples used in this paper that are relevant to automotive applications once they are adjusted for cost. This matter is briefly discussed at the end of the paper.

Abstract

In nanostructured metals with grain sizes of the order of 100 nm, dislocation mechanisms remain dominant in controlling plastic deformation. These materials, similar to their coarse-grained counterparts that have been subjected to heavy cold work, can no longer go through the several strain hardening stages of normal metals and are hence susceptible to plastic instabilities such as necking in tension. For processing and applications, it is obviously important and often necessary to control such inhomogeneous plastic deformation. Here we demonstrate three strategies to achieve relatively large stable tensile deformation in nanostructured metals, using the pure Cu processed by equal channel angular pressing as a model. The first approach uses an in situ formed composite-like microstructure, such as a bimodal grain size distribution, to impart strain hardening to the material and attain large uniform tensile strains while maintaining the majority of the strengthening brought forth by nanostructuring. In the second route, deformation is conducted at low temperatures, such as 77 K. The material regains the ability to work harden due to suppressed dynamic recovery. Uniform elongation is achieved as a result, together with an elevated strength at the cryogenic temperature. The third method takes advantage of the elevated strain rate sensitivity of the flow stress of the nanostructured Cu, especially at slow strain rates. Using the stabilizing effects of strain rate hardening on tensile deformation, nearly uniform strains can be acquired in absence of strain hardening. We also discuss the deformation mechanisms involved in these approaches to assess their applicability to nanocrystalline metals with grain sizes well below 100 nm, where normal dislocation activities become severely suppressed.

This article presents our take on the area of bulk ultrafine-grained materials produced by severe plastic deformation (SPD). Over the last decades, research activities in this area have grown enormously and have produced interesting results, which we summarise in this concise review. This paper is intended as an introduction to the field for the “uninitiated”, while at the same time highlighting some polemic issues that may be of interest to those specialising in bulk nanomaterials produced by SPD. A brief overview of the available SPD technologies is given, along with a summary of unusual mechanical, physical and other properties achievable by SPD processing. The challenges this research is facing—some of them generic and some specific to the nanoSPD area—are identified and discussed.

The evolutions of the microstructure and texture of AA1050 alloy cold rolled to large strains have been investigated using electron channeling contrast (ECC) imaging and electron backscattered diffraction (EBSD) techniques. It is found that the microstructure evolves from a cell--block structure at low strains into a lamellar structure at high strains, within which most of lamellar boundaries (LBs) are parallel to rolling direction (RD). Two mechanisms contribute to the microstructure transition, i.e., a gradual reorientation of the cell--block boundaries toward to RD due to the cold rolling deformation (Mechanism I, which is the dominant importance) and the realignment of boundaries to RD as a result of the shearing introduced by S--bands structure (Mechanism II). During this process a significant number of high angle boundaries (HABs) is created, about 47% HABs originate from deformation--induced boundaries at 90% reduction. The number of HABs increases and the spacing decreases with the increase of strains. The texture evolves into typical cold rolling deformation texture components of Brass+S+Copper, and the intensity of the texture increases with the increase of strain.

²ÉÓÃECCºÍEBSD¼¼ÊõÑо¿ÁËAA1050ÂÁºÏ½ðÀäÔþµ½´óÓ¦±äÁ¿ÏÂ΢¹Û×éÖ¯ºÍÖ¯¹¹µÄÑݱä¹æÂÉ. ½á¹û±íÃ÷, AA1050ºÏ½ðÀäÔþµ½ ´óÐαäÁ¿Ê±, ΢¹Û×éÖ¯ÓɵÍÓ¦±äϵİû¿é×éÖ¯½á¹¹×ª±ä³ÉµäÐ͵IJãƬ״½çÃæ (LBs)½á¹¹, ÆäÄÚ²¿µÄLBs»ù±¾ÓëÔþÏò (RD) ƽÐÐ; Ö÷Òª´æ ÔÚÁ½ÖÖת±ä»úÖÆ, ¼´ÓÉÓÚÔþÖƱäÐÎ (»úÖÆI) ºÍ½èÖúÓÚS--bands½á¹¹µÄ¼ôÇÐ×÷Óà (»úÖÆII),´Ó¶øµ¼ÖÂGNBsÖð½¥Ðýתµ½ÓëRDƽÐÐ, ÇÒÒÔ »úÖÆIΪÖ÷. ±äÐιý³ÌÖÐ, ÓÉÓÚ¾§Á£µÄ·ÖÁÑÐγɴóÁ¿µÄ´ó½Ç¶È½çÃæ, ËæÓ¦±äµÄÔö¼Ó, ´ó½Ç¶È½çÃæµÄ¼ä¾àÖð½¥¼õС¡¢ ÊýÄ¿Öð½¥Ôö¶à; µ±ÀäÔþµ½90%Ó¦±äÁ¿Ê±, ³ýԭʼ¾§½çÍâ, ԼΪ47%µÄ´ó½Ç¶È½çÃæÆðÔ´ÓÚ±äÐÎÓÕµ¼µÄ½çÃæ. ÀäÔþ±äÐÎÖ÷ÒªÐγɵäÐ͵ÄBrass+S+CopperÔþÖÆÖ¯¹¹, ÇÒÇ¿¶ÈËæÓ¦±äµÄÔö¼Ó¶øÖð½¥Ôö¼Ó.

Abstract

Alloys with grain or subgrain structures refined down to 1 ¦Ìm or below usually have high strength, but often inadequate tensile ductility. Past efforts in improving their ductility have usually led to a sacrifice of strength. We have developed an effective approach in achieving both high strength and high ductility in a 2024 Al alloy. The approach involves solution-treatment to partially dissolve T-phase particles, cryo-rolling to produce a fine-structure containing a high density of dislocations and submicrometer subgrains and aging to generate highly dispersed nano-precipitates. It was found that the remnant T-phase particles made it very effective in accumulating dislocations during cryo-rolling, which in turn promoted the precipitation of nanosized S′ precipitates with an interparticle spacing of only 10–20 nm. Such a high density of S′ precipitates enabled effective dislocation pinning and accumulation, leading to simultaneous increases in strength, work-hardening ability and ductility.

Although AlMgSi (6000£­series) are generally considered to have better corrosion resistance than other aluminum alloys, it may introduce susceptibility to intergranular corrosion (IGC) by some factors, for instance, improper thermo£­mechanical treatment and high Cu content. So the T78 treatment has been designed to desensitize it to IGC. In this paper, effects of T6 and T78 tempers on the microstructures and properties of an Al£­0.75Mg£­0.75Si£­0.8Cu (mass fraction, %) were investigated by hardness test, electric conductivity test,accelerated corrosion test, SEM, TEM and energy dispersive X£­ray elemental mapping. The hardness and conductivity of the T6 peak£­aged sample, which was obtained by artificial aging for 8 h at 180 ¡æ, were 128.3 HV and 27.3~10⁶ S/m, respectively. T78 tempers involved a first step aging (180 ¡æ, 5 h) followed by a second step aging at 195, 205 and 215 ¡æ, respectively.With the prolonging of second step aging, the hardness firstly decreased then increased, while the conductivity increased gradually. The optimum T78 process was(180 ¡æ, 5 h)+(195 ¡æ, 2 h), at which condition the hardness was 129.2 HV and the electric conductivity was 27.6~10⁶ S/m. TEM observation results show that it was mostly needle£­likeβ” phase in the Al matrix for the peak£­aged sample. After T78 treatment, a large amount of lath£­like precipitates formed in the matrix, while the precipitate free zones (PFZ) broadened slightly. The segregation of Cu was found at the interface between lath£­like precipitates and the matrix, so more Cu precipitated out from the Al matrix and thus reduced the electrochemical potential difference between the PFZ and the matrix. The above finding may explain why T78 temper can desensitize intergranular corrosion without sacrificing strength.

½èÖúÏÔ΢Ӳ¶È²âÊÔ¡¢µçµ¼ÂʲâÊÔ¡¢½þÅݸ¯Ê´ÊµÑé¡¢SEM, TEMºÍÔªËØÃæɨÃè, Ñо¿ÁËT6ºÍT78ʱЧ¹¤ÒÕ¶ÔAl£­0.75Mg£­0.75Si£­0.8Cu(ÖÊÁ¿·ÖÊý, %)ºÏ½ðÏÔ΢½á¹¹ºÍÐÔÄܵÄÓ°Ïì. ½á¹û±íÃ÷, ʵÑéºÏ½ðÔÚT6·åֵʱЧ(180 ¡æ, 8 h)ʱ, Ó²¶ÈΪ128.3 HV, µçµ¼ÂÊΪ27.3~10⁶ S/m. ÔÚT6(180 ¡æ, 5 h)»ù´¡ÉÏ, ·Ö±ð½øÐÐζÈΪ195, 205ºÍ215 ¡æµÄ¶þ¼¶Ê±Ð§,¼´T78ʱЧ¹¤ÒÕ, Ëæ¶þ¼¶Ê±Ð§Ê±¼äµÄÑÓ³¤, ºÏ½ðÓ²¶È×ÜÌå±ä»¯Ç÷ÊÆΪÏȽµµÍºóÉý¸ß, ¶øµçµ¼ÂÊÖµÖð½¥Ôö´ó; ×î¼ÑT78¹¤ÒÕΪ (180 ¡æ,5 h)+(195 ¡æ, 2 h), ´ËʱºÏ½ðµÄÓ²¶È×î¸ß, Ϊ129.2 HV, µçµ¼ÂÊΪ27.610⁶ S/m. TEM¹Û²ì·¢ÏÖ, T6·åֵʱЧʱ,¾§Á£ÄÚÖ÷ÒªÎö³öÕë×´µÄβ”Ïà, ¾­T78Ë«¼¶Ê±Ð§´¦Àíºó, ¾§Á£ÄÚÎö³ö´óÁ¿µÄ°åÌõ×´Îö³öÏà, ¶ø¾§½çÎÞÎö³ö´ø½öÉÔÓб仯. T78¹¤ÒÕÄÜÔÚ±£³ÖÁ¦Ñ§ÐÔÄܵÄͬʱ, ´ó·ùÌá¸ßÄ;§¼ä¸¯Ê´ÐÔÄÜ, ÆäÔ­ÒòÊǾ§Á£ÄÚÐγÉÁ˱íÃæÓÐCu¸»¼¯µÄ°åÌõ×´Îö³öÏà.¾§Á£ÄÚÎö³ö´óÁ¿°åÌõ×´Ç¿»¯Ïà, ʹµÃºÏ½ðÓ²¶ÈÓëT6·åÖµÏ൱, ͬʱCuÆ«¾ÛÔÚÎö³öÏàÓëAl»ùÌå½çÃæ´¦, ´Ùʹ¸ü¶àAl»ùÌåÖеÄCuÎö³ö,»ùÌåÓëÎÞÎö³ö´ø(PFZ)µçλ²î´ó´ó½µµÍ, ʹµÃÆ俹¾§¼ä¸¯Ê´ÐÔÄÜ´ó´óÌá¸ß.