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text:
Online Streaming Hasn't Made DVDs Obsolete—Yet
By Ryan Bort On 06/08/16 at 7:10 AM EDT
Harry Campbell
Culture Streaming
Long ago, before Netflix, Hulu and Amazon Prime, gadgets called VCRs roamed the earth. Do you remember? They were rectangular boxes that plugged into televisions, and if a videocassette (a smaller rectangular box that contained long spools of film) was inserted into a slot on the player's front, a film would begin playing on your set—that is, as long as whoever watched it last was kind enough to rewind. Most VCRs now sit on the dusty bottom shelves of thrift stores or are buried in garages, because someone figured out how to put movies onto shining, reflective discs called DVDs, and then we all bought flat, rectangular boxes for playing them.
Thanks to online streaming, DVD players may soon join VCRs on those thrift store shelves, according to a new report from research firm Strategy Analytics. The amount of money consumers have been spending on DVDs has been declining steadily for the past 10 years. In 2016, consumers are expected to spend 7 and 10 percent less on DVD sales and rentals, respectively, than they did in 2015. It's kind of like when young people notice that their aging parents are acting like their grandparents.
But just because Dad loses track of his keys a little more often than he used to doesn't mean you should reserve a room in the retirement home just yet. American consumers are still projected to spend $8.42 billion on DVD sales and rentals in 2016, versus $6.62 billion on video-streaming services. Tech-savvy Americans who cut the cord long ago may have a hard time believing DVDs still reign supreme, but there are plenty of Americans who haven't embraced on-demand streaming. For those who weren't weaned on the internet, the suddenly rapid progression of entertainment technology can be hard to grasp. But despite DVDs' current lead, projected 2016 streaming revenue is up 22 percent just from 2015.
While that revenue increase is less than the streaming industry saw in 2015—the first year that growth has slowed—it's mostly a testament to how staggering the growth was in the first place. The market for streaming video may have reached a saturation level, but it will continue to expand and sales will continue to increase. At least for 2016, though, plenty of American families are still dropping those shining, reflective discs into their DVD players. Now, to figure out which of their TV's AUX channels will play the movie...
Your Old VHS Tapes Could Be Worth Loads of Money
Hello From the Safe Side
Netflix, Amazon Could Face EU Quotas
Turning Tidal: What's Exclusive in the Streaming Age?
Downtime Culture
Periscope World
Horizons Tech & Science
Periscope U.S.
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Features Sports
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text:
Get Your Ashes on Main Street
February 26 is Ash Wednesday on the Christian calendar, a day people throughout the world will go to church, to mark the start of the season of Lent, the 46 days before Easter, often a time for penitence and reflection. The day is named for the practice of imposing ashes, a tradition that many Christians have found to be a meaningful part of the Ash Wednesday liturgy.
If, however, you can't make it to church on Ash Wednesday, the ashes can come to you! Saint Stephen Lutheran Church offers this creative alternative. Pastor Joseph Graumann and assisting minister Ann Gibson will be available in Marlborough near the corner of Main and Florence streets, from noon to 1:00 pm on Ash Wednesday, and will impose ashes and offer a blessing to anyone who asks. Those who prefer not to receive ashes are also welcome to come for prayer or conversation.
"Ash has been a symbol of death and mourning since ancient times," Pastor Graumann says. "Christians mark their foreheads with ashes to remember that 'we are dust, and to dust we shall return.' We remember our mortality, but also in the sign of the cross, how Jesus triumphed over death," he adds.
Lenten Services February 26-April 1
In addition to noontime ashes and prayer on Main Street, St. Stephen Lutheran Church will offer traditional Ash Wednesday worship at 7:00 p.m. on February 26 at the church at 537 Bolton St., Marlborough.
On subsequent Wednesdays in Lent, March 4, 11, 18, 25 and April 1, people are invited to gather at St. Stephen for a simple soup supper (one soup is vegetarian) at 6:30 PM, followed by a choral Vespers—or evening prayer—service. "Lent is a time of discipline, but it is also a time of celebration and beauty," Pastor Graumann says. "On March 25, our friends from Temple Emanuel in Marlborough will be joining us," he noted.
Vespers will feature Marty Haugen's Holden Evening Prayer, along with a reading and reflection from the Lenten stories from the Hebrew Bible. Holden Evening Prayer "is easy to sing and focuses on Christ as the light of the world. During evening prayer, we hear the story of Mary and her response to God's news that she would bear for us a savior. We pray for a world in need, and we receive God's blessing," Pastor Graumann said.
Written in 1985-86 by Marty Haugen during a musical residency at Holden Village, a scenic Lutheran camp in the mountains of Washington State, Holden Evening Prayer is a simple Lutheran vespers service that follows traditional form while using contemporary and inclusive language. Since then, the service has been adopted around the world. Haugen is a composer of liturgical music for Lutheran, Roman Catholic and other denominations, with over 400 compositions.


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It was the third Sunday of the month, which meant the day belonged to the bicycling community in Bangalore and specifically the Bangalore Bicycle Championships, which execute some of the best bicycle races in the country! Bhartiya City share a similar love for bicycling with proper bicycle infrastructure to promote bicycling and eco friendly living. This time around the race was a Cyclocross event, a much awaited fun endurance event. A combination of intense riding in dirt, tarmac, grass, slush, basically anything the cycle can be ridden on. This race location was perfect; the terrain had a hard gravel surface with a mixture of grass and slush due to the previous days rain. All of which led to the participants and their bikes turning out to be clay models by the end of the race.


text:
Stegostoma fasciatum е вид акула от семейство Stegostomatidae. Възникнал е преди около 47,8 млн. години по времето на периода палеоген. Видът е уязвим и сериозно застрашен от изчезване.

Разпространение и местообитание
Разпространен е в Австралия (Западна Австралия, Куинсланд, Нов Южен Уелс и Северна територия), Бангладеш, Бахрейн, Виетнам, Йемен, Индия, Индонезия, Иран, Камбоджа, Катар, Китай, Кувейт, Мадагаскар, Малайзия (Саравак), Малдиви, Мозамбик, Нова Каледония, Обединени арабски емирства, Оман, Пакистан, Палау, Папуа Нова Гвинея, Провинции в КНР, Саудитска Арабия, Сингапур, Тайван, Тайланд, Танзания, Филипини, Шри Ланка, Южна Африка и Япония.

Обитава крайбрежията и пясъчните дъна на сладководни и полусолени басейни, океани, морета, заливи и рифове в райони с тропически климат. Среща се на дълбочина от 10 до 144 m, при температура на водата от 20,4 до 28,4 °C и соленост 34,4 – 35,5 ‰.

Описание
На дължина достигат до 3,5 m.

Популацията на вида е намаляваща.

Галерия

Източници

Литература

Външни препратки

Stegostoma

redpajama_set_name: RedPajamaWikipedia


text:
You are here: Parliament home page > Parliamentary business > Publications and Records > Committee Publications > All Select Committee Publications > Commons Select Committees > Public Accounts > Public Accounts
Select Committee on Public Accounts Minutes of Evidence
Examination of Witnesses (Questions 40 - 59)
MR ROBIN YOUNG, KCB AND DR ALAN BORG, CBE
40. When?
(Dr Borg) In the current year.
41. Really?
(Dr Borg) We are 13 per cent up.
42. 13 per cent up? Whereabouts does that put you? We are nearly at the end of the year, where does that put you? I make it back to something like the level of 1997-98 looking at the diagram.
(Dr Borg) We are forecasting just short of 1.4 million this year.
43. That is almost back to 1998-99. But you are still well below 1994-95. I do not doubt your ambitions but they seem unduly modest.
(Dr Borg) We also believe that the figures will continue to go up because we are opening later this year the entirely refurbished British Galleries, which account for ten per cent of the Museum's display space and we also have in train the proposal, as I mentioned, to build the Spiral Building which I think will contribute when it is there in a dramatic way to our audience size.
44. Look at Table 8 on the same page and that shows the three branch museums. The interesting thing is the aberration. The other two seem to be following in the example of South Ken, but then you have the Theatre Museum which has been on an almost continuous growth chart. Why have they been able to achieve growth but the other four units not been able to achieve growth?
(Dr Borg) I think there are a number of reasons for the success of the Theatre Museum. One, perhaps most importantly, is its location. It is set in the heart of Covent Garden in the heart of theatre land in a very much visited part of London. Therefore the possibility of drop-in visitors is much higher than it is at South Kensington or certainly Bethnal Green, for example, but it is also true to say that the Theatre Museum has had a good and imaginative educational programme which has also brought visitors in.
45. It is a highly competitive field, is it not, trying to draw visitors in? There are so many visitor attractions and youngsters, particularly nowadays, are used to being pandered, and quite rightly, with hands-on experience, they learn a lot more in that way rather than inanimate, stodgy exhibits. What scope—and I recognise it varies from sector of market to sector—have you discovered, if any, for hands-on approaches in your own operations?
(Dr Borg) Quite a number and an increasing number. I have already referred to the educational programmes we use, the backpacks, the activity cart in which children make things related to the collections. We also have tours for, for example, the visually impaired where people can handle sculpture and can pick up objects. I think all of those sort of experiences are very important within the Museum generally.
46. Turning to Page 33 and looking at Figure 23, activity backpacks, to which you have just referred. These are aimed at 5-to-11 year olds, a good age to capture youngsters and get their imagination, and are designed to help families move around the Museum and to explore themes connected with the collections. There are 36 packs in total. I do not know these packs but is it one pack per person. Is a pack something that is individual to one visitor?
(Dr Borg) That refers to the programme that the pack contains. If you come you are given a particular programme to follow, and there are 36 of those.
47. That is not limited.
(Dr Borg) I am informed there are now 50, and they are increasing all of the time.
48. 50 programmes?
(Dr Borg) Yes.
49. How many people at one time would be able to go round with the aid of these backpacks?
(Dr Borg) Probably about 50.
50. That is not particularly market oriented, is it?
(Dr Borg) This is dealing on a one-to-one basis.
51. That is what I said in the first place, and you said that I was misunderstanding you. If you are talking about attracting youngsters in and you only have three dozen of these visitor aids it does not look as if you will have a highly optimistic view of the numbers you are likely to attract, does it?
(Dr Borg) That is only one of the activities we provide for children, we provide many others.
Mr Williams: I am sure they are equally exciting. My time is up, thank you very much.
Mr Steinberg
52. If you look at Figure 7, page 17 it is quite clear, as has been mentioned by previous members of the Committee, that the number has declined since 1994. It has not been a constant drop, has it? If you look at the chart it tends to have gone up one year and gone down the next year, gone up the next year and then gone down the next year, until eventually it has reached its low in 1999-2000. Can you explain this peculiar pattern?
(Dr Borg) The pattern is due, almost exclusively, to the varied number that attend temporary exhibitions. Table 22 on page 32 demonstrates it visually.
53. What you are saying is in the years it goes up it is because you have put something on and in the years it has gone down you have not bothered to do anything extra.
(Dr Borg) In all years there are temporary exhibitions, but some are better attended than others.
54. I see. What you are saying is in the years that actually improved there were popular exhibitions taking place and in the years that it declined there were not. Why does it have a pattern like that? Do you plan specifically to have a popular exhibition one year and not a popular exhibition the year after?
(Dr Borg) We plan to have good exhibitions all of the time. We are aware that some are going to appeal to a wider audience and some to a narrower audience. We see it as part of the remit of the V&A to appeal to all of our audiences, including putting on exhibitions, sometimes directed at what we know is quite a small audience.
55. Okay. In the years that it seems to rally does that give you a false hope?
(Dr Borg) I do not think it is gives us a false hope because we are, on the whole, reasonably good, and occasionally not, at forecasting what we think attendance at exhibitions is going to be.
56. Even so, the figures have declined something like, half a million people over the last 6 years. Whether the exhibitions have been good, bad or indifferent, the fact of the matter is that half a million people less are coming into the Museum.
(Dr Borg) That trend is reversed now and the figures are going back up. We believe that the opening of the British galleries and the opening of the Spiral building, when that is constructed, will boost our figures very dramatically.
57. If you look at paragraph 2.7, that does not give you a lot of confidence in the management, does it? It says that in the years 1999-2000 there is a marked drop in visitor numbers. This was quite apparent six months into the year and yet nothing was done to find out why. Why was that? If you saw in that year that your numbers were dramatically falling and it took six months to take some action, why did you take so long?
(Dr Borg) On the whole we are aware of the reasons why an exhibition does not attract visitors. Taking remedial action is something that is done in the longer term. Museums are, if you like, like oil tankers at sea, they cannot be turned round in a matter of moments.
58. With great respect, it has been in decline, if you like, in terms of numbers for six years, that is quite a considerable length of time to try and turnaround the oil tanker, is it not?
(Dr Borg) Absolutely. That trend is now reversed and the longer term measures which we put in place, the refurbishment of the British galleries and the building of the Spiral building, are what will contribute to a permanent reversal of that trend.
59. Is this borne out in research that you have done?


text:
\section{Introduction}

Zirconium alloys are used in the nuclear industry as fuel cladding tubes and structural components in light and heavy water reactors \cite{Motta2017}.
Like in other hexagonal close-packed (hcp) metals,
plastic activity strongly varies between the different possible slip systems in zirconium.
The principal slip system is controlled by glide of \hkl<a> dislocations, \textit{i.e.}{} dislocations with $1/3\,\hkl<1-210>$ Burgers vectors, in the \hkl{10-10} prismatic planes
\cite{Caillard2003,Clouet2015}.
These \hkl<a> dislocations can also cross-slip above room temperatures to glide in the \hkl(0001) basal and \hkl{10-11} first-order pyramidal planes \cite{Caillard2018}.
However all these deformation modes only accommodate deformation along \hkl<a> directions
and other mechanisms, either twinning or glide of dislocations with a \hkl<c> component,
are needed to allow for a deformation along the $\hkl<c>=\hkl[0001]$ axis of the hcp crystal.

Both twinning and glide of \hkl<c+a> dislocations, \textit{i.e.}{} dislocations with $1/3\,\hkl<11-23>$ Burgers vectors, are active in zirconium
\cite{Tenckhoff1972,Jensen1972,Akhtar1973,Woo1979,Holt1987,Merle1987,Numakura1991,McCabe2006,Long2015a,Long2015b,Long2017},
with $\hkl<c+a>$ slip becoming more active when the temperature increases.
Depending of the hcp metals, these \hkl<c+a> dislocations can glide in the first-order \hkl{10-11}
or second-order \hkl{-2112} pyramidal planes.
Almost all transmission electron microscopy (TEM) observations have shown, by slip traces analysis \cite{Akhtar1973}
or stereography \cite{Woo1979,Merle1987,Numakura1991,Long2015a,Long2015b,Long2017,Long2018},
that \hkl<c+a> dislocations glide in first-order pyramidal planes in Zr alloys,
with possible cross-slip between two pyramidal planes sharing the same \hkl<c+a> direction \cite{Long2015b,Long2017,Long2018}.
Scanning electron microscopy of slip traces on bended micro-cantilever well oriented to activate only \hkl<c+a> slip
also concluded that \hkl<c+a> dislocation glide in first-order pyramidal planes,
without any trace of slip in second-order pyramidal planes,
both for the compressive and tensile parts of the cantilevers \cite{Gong2015}.
Only Long \textit{et al}.{} \cite{Long2018} have recently evidenced that a minority of \hkl<c+a> can also be found in the second order pyramidal plane in a Zr-2.5Nb alloy.
Besides, some of these TEM observations indicate that \hkl<c+a> dislocations have a tendency to straighten
in the direction corresponding to the intersection of the pyramidal glide plane with the basal plane \cite{Woo1979,Numakura1991,Long2017,Long2018},
a feature which is not specific to zirconium and can be found in other hcp metals, like Mg \cite{Stohr1972,Geng2014,Geng2015,Wu2015,Wu2018,Zhang2019}
or Ti \cite{Minonishi1982,Numakura1986}, regardless of the pyramidal glide plane.

The critical resolved shear stress (CRSS) necessary to activate glide of \hkl<c+a> dislocations strongly depends on the temperature
but is always much higher than the CRSS of \hkl<a> dislocation glide \cite{Akhtar1973,Gong2015},
thus explaining the strong plastic anisotropy of zirconium single crystals.
An increased activity of \hkl<c+a> slip is observed in irradiated zirconium alloys,
as the hardening induced by irradiation defects appears stronger on \hkl<a> slip systems than on \hkl<c+a> pyramidal slip \cite{Long2015b,Long2016},
thus partly compensating for the friction difference between the two types of dislocations.

Although the mobility of \hkl<c+a> dislocations is a key ingredient of the plastic anisotropy of zirconium, not much is known on this mobility
besides the slip plane, the dislocation ability to cross-slip and their tendency to align in a specific direction as mentioned above.
\textit{In situ}{} TEM straining experiments are a valuable tool to study dislocation mobility.
Such experiments in zirconium have already revealed some key features of \hkl<a> dislocations
\cite{Clouet2015,Long2015b,Caillard2015,Drouet2016,Caillard2018,Gaume2018}.
But no such \textit{in situ}{} observations have been performed for the glide of \hkl<c+a> dislocations in zirconium.
The difficulty arises from the necessity to have grains well oriented to activate \hkl<c+a> slip.
The aim of the present article is to present such experiments.
\textit{In situ}{} TEM straining experiments have been performed in two different zirconium alloys, Zircaloy-4
and pure zirconium, as previous experimental studies on polycrystals \cite{Jensen1972} have shown
a difference in the strain accommodation along the \hkl<c> axis between Zircaloy-4 and pure zirconium.
The results of these \textit{in situ}{} experiments are presented below.

\section{Experimental procedure}

\subsection{Materials}

Two different zirconium materials have been chosen for this study: zirconium sponge, which is essentially pure zirconium,
and Zircaloy-4 in the recrystallized metallurgical state, which is often considered as a model material for zirconium alloys.
Their chemical compositions are given in Table \ref{tab:Zr_composition}.

\begin{table}[!tbh]
	\caption{Chemical composition (wt.\,\%) of zirconium materials used in this study.}
	\label{tab:Zr_composition}
	\centering
	\begin{tabular}{cccccc}
		\hline
			& Sn & Fe & Cr & O & Zr \\
		\hline
		Pure Zr & - & 0.019 & - & 0.024 & bal. \\
  		Zircaloy-4 & 1.32 & 0.215 & 0.108 & 0.125 & bal. \\
		\hline
	\end{tabular}
\end{table}

\subsubsection{Pure zirconium}

Pure zirconium results from the Kroll process in the form of a 3\,mm thick plate.
The raw zirconium sponge has first been arc-melted and cold-rolled with intermediate recrystallization treatments during 3 hours at 580\,$^{\circ}$C. After the last cold-rolling step, the same recrystallization treatment has been applied followed by a tensile test up to a strain of 1.5$\%$. The purpose of this straining was to reach critical strain annealing phenomenon in order to obtain large grains \cite{Chaubet2001,Zhu2005}. This straining was followed by an annealing at 840\,$^{\circ}$C during 24 hours. After this treatment, the mean grain size was 70\,$\mu$m with a large dispersion around this value. The material texture exhibited \hkl<c> axis mainly along the normal direction (ND) of the plate. Small samples, 3\,mm long and 1\,mm wide, were taken out of the plate with the long direction along ND, to activate \hkl<c+a> slip.

\subsubsection{Zircaloy-4}

Recrystallized Zircaloy-4 is used for light water reactor applications as guide tubes.
These tubes have a thickness smaller than 1\,mm and exhibit a strong texture with the \hkl<c> axis of the grains close to the radial direction (RD) of the tube, in the AD-RD (AD: axial direction) plane.
Such products are not suitable to study easily the activation of pyramidal \hkl<c+a> slip, as one has to pull the grains along directions close to the \hkl<c> axis which are mainly along the radial direction of the thin tube. Because of the small thickness of the tube along the radial direction, it is not possible to machine small tensile test samples in this direction. This is the reason why an intermediate product of the fabrication process, called TREX for Tube Reduced Extrusion, has been used. This is a thick tube, with diameter of the order of 55\,mm and thickness of the order of 10\,mm. This material exhibits equi-axed grains with mean size of 6\,$\mu$m and a low initial dislocation density, only of the order of $10^{11}$\,m$^{-2}$.
The \hkl<c> axis of the grains are mainly oriented in the RD-TD (TD: transverse direction) plane with the highest pole density along TD \cite{Tournadre2012,Gharbi2015}. Small dog-bone specimens, 11.5\,mm long and 2.3\,mm wide, were machined by electrical discharge, with the tensile direction along TD. Therefore, thanks to the texture, a large number of grains are expected to have their \hkl<c> axis along the tensile direction.

\subsection[In situ TEM straining experiments]{\textit{In situ}{} TEM straining experiments}

Pure zirconium and Zircaloy-4 samples have been mechanically polished down to 0.1\,mm thick and electropolished to make a small hole, using a mixture of 90$\%$ of ethanol and 10$\%$ of perchlorate acid at $-10$\,$^{\circ}${}C. The thin area around the hole is only few hundreds nanometer thick and is therefore suitable for TEM observations.

\textit{In situ}{} straining experiments were performed at room temperature on both materials using a GATAN  cooling and straining TEM sample holder on a JEOL 2010 operating at 200\,kV.
The thin foils were pasted on two ceramic rings or on copper grids at each ends to be pulled by the TEM holder.
For Zircaloy-4, experiments have been also conducted at room temperature on a FEI Tecnai operating at 200\,kV, using the commercial room temperature GATAN straining specimen-holder and the dog-bone TEM samples.
On these two specimen holders, one cross-head of the sample is fixed and the other cross-head is attached to a bar which displacement is controlled by a motor.
Typical displacements are 100\,$\mu$m.

The experimental procedure is as follows. First, the orientation of various grains is analyzed until finding a grain with an orientation that satisfies two criteria: i) the \hkl<c> axis must be close enough to the tensile direction
(typically less than 10\,$^{\circ}${}) so that \hkl<c+a> slip will be activated, ii) the grain can be tilted, in the limit of the tilt range and with the restriction of single-tilt allowed by the TEM sample holder, so that a two beam diffraction condition with $\vec{g}=0002$ diffraction vector can be obtained. In this condition, the \hkl<a> dislocations disappear and only \hkl<c+a> dislocations are observed. The number of grains that meet these two conditions appears to be rather limited, despite the effort to select materials with adequate texture.

Once the grains have been selected, the displacement of the cross-head is progressively increased until \hkl<c+a> dislocation motion is observed. The displacement is then maintained constant or slightly decreased and the dislocation motion is recorded using an ORIUS wide angle GATAN camera operating at ten images per second or a megaview III camera operating with 50 images per second.
Supplementary video related to this article can be found at [INSERT DOI].

For some experiments, it has been possible to achieve extinction conditions for the \hkl<c+a> dislocations, by tilting the sample. This allows the determination of the exact Burgers vector of the dislocation. For most of experiments only the glide plane of the dislocation was deduced, without ambiguity, using two informations : i) the orientation of the slip traces and ii) the evolution of the distance between the two slip traces when tilting the thin foil. From this last measurement, the thickness of the thin foil can also be estimated. Furthermore, when cross-slip was observed, the Burgers vector of the dislocation has been deduced from the knowledge of the two gliding planes.

\section{Results in recrystallized Zircaloy-4}

Besides the \textit{in situ}{} TEM straining experiments described below,
some \textit{post mortem}{} observations on the same recrystallized Zircaloy-4 have been conducted
after tensile tests to check that the behavior of \hkl<c+a> dislocations characterized \textit{in situ}{}
in thin foils concurs with the dislocation microstructure obtained after straining macroscopic samples.
These \textit{post mortem}{} observations are described in \ref{sec:postmortem}.

During the \textit{in situ}{} experiments, the observation of \hkl<c+a> dislocation motion was always associated with the observation of twinning
which often occurred after \hkl<c+a> dislocation glide.
The twinning system was \hkl[10-11]\hkl{10-1-2}
as expected in tension in zirconium alloys at room temperature \cite{Rapperport1960,Yoo1991,Viltange1985}.
\ref{sec:twinning} describes an analysis of one twin event where it has been possible
to measure the twin propagation and thickening velocity.
In the following, we only focus on the description of \hkl<c+a> dislocation glide.

\subsection{Glide planes and dislocation orientation}

\begin{figure}[!bth]
	\centering

	\includegraphics[width=\linewidth]{fig1.jpg}
	\caption{Zircaloy-4 strained in situ at room temperature. The dislocations are observed using a diffraction vector (a) $\vec{g}=0002$, (b) $\vec{g}=\bar{1}011$ and (c) $\vec{g}=\bar{1}01\bar{1}$. The tensile axis is vertical on the picture and close to the normal of the basal plane noted B. Two dislocations denoted as 1 and 2 glide in a \hkl(0-11-1) plane noted $\pi$. From the two possible Burgers vectors, the extinction shown in (c) gives $1/3\,\hkl[-1-123]$ as the Burgers vector. It is observed that the dislocations tend to be aligned along the \hkl[-1-120] direction noted $\pi \cap \textrm{B}$, which is the intersection of the first-order pyramidal glide plane and the basal plane.
	}
	\label{fig:extinction}
\end{figure}

\begin{figure}[!bth]
	\centering

	\includegraphics[width=\linewidth]{fig2.pdf}
	\caption{In situ TEM observations of \hkl<c+a> dislocations gliding in two different pyramidal planes.
	The corresponding slip traces are shown with dashed white lines.
	For both slip systems, the dislocations are aligned along a \hkl<-2110> direction:
	the \hkl[2-1-10] direction for the two short horizontal dislocations gliding in \hkl(01-11)
	(slip plane in red in the stereographic projection)
	and the \hkl[-12-10] direction for the long horizontal dislocation gliding in \hkl(10-1-1) (slip plane in blue).
	Possible $1/3\,\hkl<1-213>$ Burgers vectors are indicated by filled circles on the stereographic projection.}
	\label{fig:traces_py1}
\end{figure}

Gliding \hkl<c+a> dislocations in recrystallized Zircaloy-4 are shown in Figures \ref{fig:extinction} and \ref{fig:traces_py1}.
In Figure \ref{fig:extinction}, it has been possible to determine exactly the Burgers vector $\vec{b}$ using different tilts:
the dislocations are visible with diffraction vectors $\vec{g}=0002$ and $\vec{g}=\bar{1}011$, but invisible with $\vec{g}=\bar{1}01\bar{1}$,
thus showing unambiguously that $\vec{b}=1/3\,\hkl[-1-123]$.
The dislocation glide plane could be determined from the direction of the slip traces on the surface and from the variations of the projected size of the curved dislocation and of its slip traces with the tilt angle:
these $1/3\,\hkl[-1-123]$ dislocations
glide in a \hkl(0-11-1) first-order pyramidal planes (Fig. \ref{fig:extinction}).
All the slip traces analyzed during our \textit{in situ}{} tensile tests lead to a first-order pyramidal plane.
When several \hkl<c+a> slip systems were activated in the same grain (Fig. \ref{fig:traces_py1}),
they all correspond to the first-order pyramidal slip plane.
In all grains where \hkl<c+a> slip was activated, the corresponding Schmid factors were higher than or equal to 0.36,
while the Schmid factors of \hkl<a> prismatic slip systems was smaller than 0.01.
As a consequence, no \hkl<a> dislocations was observed at this stage in grains where \hkl<c+a> slip is activated,
but \hkl<a> dislocations could be observed in the surrounding grains.
As plastic strain increases, some \hkl<a> dislocations could sometimes also be observed in the same grains as \hkl<c+a> dislocations,
probably because of the complex mechanical load coming from the deforming surrounding grains,
from the appearance of some cracks and from twinning, and also because of crystal rotation.
But these \hkl<a> dislocations clearly appear latter than \hkl<c+a> dislocations.
We could see some \hkl<c+a> dislocation nucleated on grain boundaries or on twins.
But most of the time, \hkl<c+a> dislocations are coming from thick areas of the thin foil which could not be imaged.

In all these \textit{in situ}{} straining experiments (Figs. \ref{fig:extinction} and \ref{fig:traces_py1}),
it can be noticed that the dislocation lines exhibit straight segments perpendicular to the $\vec{g}=0002$ diffraction vector.
The \hkl<c+a> preferentially align in a direction defined by the intersection of their first-order pyramidal glide plane with the basal plane,
that is, along an \hkl<a> direction.
This dislocation orientation is nearly edge with a character $\theta$ defined by
$\cos{(\theta)}=a/2\sqrt{a^2+c^2}$.
For Zr ($c/a=1.593$), this leads to $\theta=75^{\circ}$.

\subsection{Dislocation mobility}

\begin{figure}[!b]
	\centering

	\includegraphics[width=\linewidth]{fig3.png}
	\caption{Glide of a \hkl<c+a> dislocation in the \hkl(01-1-1) pyramidal plane.
	The same dislocation is observed at three different times:
	the dislocation is almost straight in (a) and (b) with a line direction close to \hkl[2-1-10],
	whereas macro-kinks can be seen in (c).
	Figures (d) and (e) are obtained by subtraction respectively of images (a) and (b)
	and of images (b) and (c).
	In these figures, the initial and final positions of the dislocation appear
	with black and white contrast respectively.
	The dislocation Burgers vector has been determined to be $1/3\,\hkl[11-23]$
	among the two possible ones shown on the stereographic projection (f).
	See video as a supplementary material.
	}
	\label{fig:glissement_coin}
\end{figure}

\begin{figure}[!bth]
	\centering

	\includegraphics[width=\linewidth]{fig4.pdf}
	\caption{Jump of the screw part of the \hkl<c+a> dislocation shown in Fig. \ref{fig:glissement_coin}.
	The image subtraction (a) corresponds to a time interval of 0.1\,s during which the screw part jumps a distance
	$\Delta L\simeq 123$\,nm whereas the part aligned in the \hkl[2-1-10] direction moves slowly.
	The initial and final positions of the dislocation, as well as the associated slip traces,
	are sketched respectively in black and orange in (b).}
	\label{fig:glissement_vis}
\end{figure}

\begin{figure}[!bth]
	\centering

	\includegraphics[width=0.8\linewidth]{fig5.png}
	\caption{Jump length distribution for the \hkl<c+a> dislocation shown in Fig. \ref{fig:glissement_coin}.
	The jumps have been measured on a time interval $\Delta t=0.1$\,s in a viscous motion regime and the whole analyzed sequence lasts 49.2 s.
	The blue histograms correspond to the jumps of dislocation segments along the \hkl[-2110] direction
	and the purple ones to all other directions.}
	\label{fig:histogram}
\end{figure}

\begin{figure}[!bth]
	\begin{center}

		\includegraphics[width=\linewidth]{fig6.png}
	\end{center}
	\caption{Glide of \hkl<c+a> dislocations observed in the same zone as Fig. \ref{fig:extinction}.
	The dislocation labeled 1 arrived in the observation zone between images (a) and (b)
	in a time interval $\Delta t=0.18$\,s.
	Its straight portion aligned in the \hkl<a> direction noted $\pi \cap \textrm{B}$
	is then moving slowly between images (b) and (c) in $\Delta t=1.7$\,s.
	See video as a supplementary material.}
	\label{fig:glissement_rapide_coin}
\end{figure}

\textit{In situ}{} observations of the dislocation motion in recrystallized Zircaloy-4 show that
the \hkl<c+a> dislocations aligned along the \hkl<a> direction glide slowly in a rigid manner
with occasionally the formation of macro-kinks.
This can be seen on Fig. \ref{fig:glissement_coin},
where the \hkl<c+a> dislocation aligned in the \hkl<a> direction is nearly horizontal
and the slip traces are the diagonal lines.
Dislocation glide is better analysed on Figs. \ref{fig:glissement_coin}d and \ref{fig:glissement_coin}e which are obtained as subtraction
of the same area observed at two different times so as to reveal the differences between the two images,
and thus the dislocation motion in the time interval.
Between the initial time $t=0$ and $t=10$\,s, the nearly edge part of the \hkl<c+a> dislocation
has kept its straight \hkl<a> orientation while gliding viscously a distance $\sim100$\,nm.
At this time, this dislocation gets pinned in its middle with only one half keeping gliding
in the following 7.4\,s.  This leads to the formation of a macro-kink.
After the last frame shown in Fig. \ref{fig:glissement_coin}, the dislocation suddenly escapes from the observation zone,
with a much faster motion.
Such a fast motion of the nearly edge part of the \hkl<c+a> dislocation has been captured
in another \textit{in situ}{} straining experiment.
In Fig. \ref{fig:glissement_rapide_coin}, one can see that the dislocation labeled 1
has arrived suddenly in the observation zone in a time interval $\Delta t=0.18$\,s
with a long part of the line aligned in the \hkl<a> direction
noted $\pi \cap \textrm{B}$ on Fig. \ref{fig:glissement_rapide_coin}d.
In the following 1.7\,s, the dislocation moves much more slowly (Fig. \ref{fig:glissement_rapide_coin}e).
Two mechanisms, leading either to a slow viscous glide or to a fast sudden motion,
seem to control the mobility of this nearly edge part of the \hkl<c+a> dislocation.

Other orientations of the \hkl<c+a> dislocation, including the screw part, glide more rapidly with larger jumps than the viscous motion of the \hkl<a> orientation,
as can be seen for instance on Fig. \ref{fig:glissement_vis}.
The shape of these dislocations segments with other orientations is much more rounded and appears to be driven by minimization of the line tension.
Nevertheless, a lot of pinning points can be seen (Figs \ref{fig:extinction}, \ref{fig:glissement_vis} and \ref{fig:glissement_rapide_coin}),
showing that these segments are anchored on numerous localized point obstacles.

We have tried to evaluate the difference of mobility in the viscous glide regime between the part of the \hkl<c+a> dislocation
aligned along the \hkl<a> direction (nearly edge) and the other mixed-screw orientations for the sequence shown in Fig. \ref{fig:glissement_coin}.
The glide distance $\Delta L$ between two frames ($\Delta t=0.1$\,s) is measured for each part of the dislocation (Fig. \ref{fig:glissement_vis}b).
If the dislocation does not move in this time interval, this is added to the waiting time of the dislocation.
The overall waiting time accounts for 70$\%$ of the sequence duration ($\Delta t=49.2$\,s).
Only the time intervals corresponding to a dislocation motion are then analyzed,
grouping measured glide distances by bins of width 10\,nm so as to build a histogram
showing jump frequency as a function of glide distances
both for the nearly edge and the other mixed orientations of the \hkl<c+a> dislocation (Fig. \ref{fig:histogram}).
It then clearly appears that there is a higher frequency of short jumps, smaller than 10\,nm, for dislocation lines along the \hkl<a> direction than for the other mixed-screw parts,
thus leading to a lower glide velocity of this near edge orientation.
The large jump of 100\,nm long of the mixed-screw part shown in Fig. \ref{fig:glissement_vis} corresponds to the tail of the histogram.

\subsection{Cross-slip}

\begin{figure}[!tb]
	\centering
 	\includegraphics[width=\linewidth]{fig7.pdf}
	\caption{Dislocation cross-slip.
	(a,b) The dislocation with a Burgers vector $1/3\,\hkl[2-1-13]$ is initially gliding in the \hkl(10-11) plane.
	Its interaction with an unknown defect makes it cross-slip.
	(c) The dislocation pinned by this defect continues gliding in the initial \hkl(10-11) plane
	and in the cross-slipped \hkl(1-101) plane.
	On the right sketches, the dislocation is drawn with a red and blue line when gliding
	respectively in its initial and crossed-slip plane.
	Bold black lines are the slip traces on the thin foil surface.
	See video as a supplementary material.
	}
	\label{fig:cross_slip}
\end{figure}

During \textit{in situ}{} straining experiments, cross-slip of  \hkl<c+a> dislocations has been observed. One of these cross-slip events is illustrated on Fig. \ref{fig:cross_slip},
where the dislocation exhibits a light contrast and the slip traces a strong contrast. New slip traces appear with another orientation, coming from the initial slip traces.
Both slip planes correspond to first-order pyramidal planes. When analyzing carefully the sequence to understand the geometry of the cross-slip event, it is found out that there must be an obstacle along the track of the dislocation.
This obstacle could be a nano-hydride, a small precipitate or even a forest dislocation.
Because initially the dislocation is along the \hkl<a> direction, it is nearly edge and it cannot cross-slip.
When the gliding dislocation meets the obstacle, which seems to be between the middle and the bottom of the thin foil, it is pinned. The part on the left continues to glide and bends towards the screw direction. When aligned along the screw direction, the dislocation cross-slips in the other first-order pyramidal plane.

Four additional slip traces emerge between $t=0$ (Fig. \ref{fig:cross_slip}a) and $t=95$\,s (Fig. \ref{fig:cross_slip}b).
These traces correspond to two \hkl<c+a> dislocations which have also slipped in two other first-order pyramidal planes,
leading to two slip traces on the top surface of the thin foil and two on the bottom surface.
The slip traces on the top and bottom surfaces are not parallel because of the wedge shape of the thin foil.
We have checked that these additional \hkl<c+a> dislocations did not interact with the ones which have cross-slipped.

The process detailed here points out that \hkl<c+a> dislocations can cross-slip when they meet an obstacle
and that the cross-slipped plane is also a first-order pyramidal plane.

\subsection{Relaxation}

\begin{figure}[!tb]
	\centering

	\includegraphics[width=0.7\linewidth]{fig8.png}
	\caption{Relaxation of a \hkl<c+a> dislocation when the displacement
	of the sample holder is slightly decreased.
	Between (a) the initial and (b) the final times, the dislocation on the left glides back
	in its slip plane and partially removes the slip trace on one of the thin foil surface.
	This is better seen on the zoom-in differential image (c).
	The stereographic projection of the grain indicates that the gliding dislocation
	is close to a screw orientation,
	its Burgers vector being either $1/3\,\hkl[11-2-3]$ or $1/3\,\hkl[2-1-1-3]$.
	See video as a supplementary material.
	}
	\label{fig:glide_back}
\end{figure}

When decreasing the stress applied on the specimen, by reducing the displacement of the cross-head from 120 to 100\,$\mu$m,
a dislocation has been found to move back on its track (Fig. \ref{fig:glide_back}).
This dislocation relaxation has only been observed once in Zircaloy-4, but has also been observed in pure Zr.
In the case of Zircaloy-4 shown on Fig. \ref{fig:glide_back},
it happens for a \hkl<c+a> dislocation which is far from the specific \hkl<a> orientation.
When decreasing the applied stress, the dislocation moves back to align in a direction closer to the screw orientation,
since screw dislocations have a lower line energy than dislocations with an edge character.
This relaxation of the \hkl<c+a> is therefore probably driven by the line tension,
thus further showing that the lattice friction is not as strong for these orientations than for the \hkl<a> direction.
One also notices that the back motion of this \hkl<c+a> dislocation erases the slip traces on the thin foil,
proving that, at least in this particular case, dislocation glide is well confined in a single crystallographic plane.

\section{Results in pure Zr}

The same \textit{in situ}{} TEM straining experiments have been performed at room temperature on pure zirconium.
Comparing these observations with the previous ones in Zircaloy-4 gives some insights on alloying effects
on the mobility of \hkl<c+a> dislocations.

\subsection{Dislocation motion}

\begin{figure}[!b]
	\centering

	\includegraphics[width=\linewidth]{fig9.jpg}
	\caption{Pure Zr strained at room temperature.
	The \hkl<c+a> dislocations are gliding in a pyramidal plane and are preferentially aligned in the \hkl<a> direction
	defined by the intersection of the pyramidal glide plane with the basal plane.
	Two dislocations denoted as 1 and 2 appear between $t=0$\, and 2.7\,s:
	the dislocation 1 pulls two straight \hkl<a> segments and ends with a curved line,
	while the dislocation 2 has slipped to the surface where it has left a slip trace  tr.~$\pi$.
	Between $t= 2.7$ and 3.4\,s, the upper and right parts of dislocation 1 unlock and jump to a new position
	while its lower straight section extends to the surface without moving.
	Between $t=3.4$ and 28\,s the upper straight part of the dislocation slightly glides
	but without any lateral motion of macro-kinks present on this dislocation.
	See video as a supplementary material.
	}
	\label{fig:zr_alignement}
\end{figure}

Like in Zircaloy-4, the slip planes of \hkl<c+a> dislocations have always been found
in pure zirconium to correspond to the first-order pyramidal planes (Fig. \ref{fig:zr_alignement}).
The dislocations also show some long rectilinear segments in the \hkl<a> direction defined by the intersection of the pyramidal glide plane with the basal plane.
Apart from this \hkl<a> orientation, the dislocations are curved, with a shape driven by the line tension.
One notices that in pure zirconium the straight portions are longer and better defined than in Zircaloy-4.
This shape difference between both zirconium samples is a signature that the difference of lattice friction between this \hkl<a> direction
and the other orientations of the \hkl<c+a> dislocations is stronger in pure zirconium than in Zircaloy-4.
The presence of numerous solute elements in Zircaloy-4 slow-downs all \hkl<c+a> segments,
thus reducing the mobility anisotropy of \hkl<c+a> dislocation.
As pure zirconium contains very few impurities, this mobility anisotropy appears more important,
with an enhanced relative lattice friction acting against the glide of the straight \hkl<a> orientation.

\begin{figure}[!bt]
	\centering

	\includegraphics[width=\linewidth]{fig10.png}
	\caption{
	Detailed glide motion of the straight \hkl<a> part of the \hkl<c+a> dislocation in pure Zr (same zone as Fig. \ref{fig:zr_alignement}).
	The dislocation ``d'' is composed of segments lying along the intersection of the pyramidal glide plane and the basal plane.
	Between (a) and (b), a lateral motion of a macro-kink is observed (see inclined arrows)
	whereas the whole dislocation is gliding between (b) and (c)
	without any motion of the macro-kink (see vertical arrows in).
	See video as a supplementary material.
	}
	\label{fig:zr_mouvement}
\end{figure}

Like in Zircaloy-4, glide of the straight \hkl<a> orientation of \hkl<c+a> dislocations leads to the creation
of macro-kinks (Fig. \ref{fig:zr_alignement}) in pure zirconium.
These macro-kinks can either glide along the dislocation line (Fig. \ref{fig:zr_mouvement}d)
or remain fixed while the dislocation moves slowly (Fig. \ref{fig:zr_mouvement}e).
Besides this viscous glide of the \hkl<a> orientation, one also observes some large jumps,
with therefore the same combination of slow and rapid motion as seen in Zircaloy-4.

A few cross-slip events have been also observed in pure zirconium.
They appear, at first glance, less numerous than in Zircaloy-4,
but a detailed statistical analysis would be needed to be able to really compare
the occurrence of cross-slip in both materials.

\subsection{Friction stress}

Some relaxation observations have been also performed in pure zirconium.
These experiments have been used to estimate the friction stress acting
against the glide of the \hkl<a> orientation and of the other orientations,
thus characterizing the mobility anisotropy.
The friction stress originates from the dislocation core structure,
like a possible non planar dissociation, and also from the dislocation interaction
with the solute atoms present in the zirconium matrix.

\begin{figure}[!bt]
	\centering

	\includegraphics[width=\linewidth]{fig11.png}
	\caption{
	Relaxation of \hkl<c+a> dislocation in pure zirconium.
	The applied strain has been released between (a) and (b), leading to a back motion of the dislocations
	and an increase of their curvature radius.
	}
	\label{fig:zr_relaxation}
\end{figure}

When the applied stress $\tau_{\rm a}$ is removed, one observed that the curved parts of the dislocations
are gliding back and that their curvature radius increases (Fig. \ref{fig:zr_relaxation}).
This variation of the curvature radius can be used to estimate the applied stress
and the friction stress $\tau_{\rm f}$ acting against the glide motion of these curved segments.
Using the program \textsc{Disdi} based on line tension calculations in anisotropic elasticity \cite{Douin1986},
one can deduce from the dislocation curvature the line tension stress $\tau_{\rm l}$.
The shape of the loops 1 to 9 under an applied stress $\tau_{\rm a}$ shown on Fig. \ref{fig:zr_relaxation}a
corresponds to a line tension stress $\tau_{\rm l}$ between 100 and 150\,MPa.
When the stress is relieved ($\tau_{\rm a}=0$ on Fig. \ref{fig:zr_relaxation}b),
we obtain a line tension stress $\tau_{\rm l}$ between 25 and 50\,MPa.
The applied stress is always in equilibrium with the line tension and the friction stresses,
$\tau_{\rm a} = \tau_{\rm l} + \tau_{\rm f}$.
We therefore deduce a friction stress $\tau_{\rm f}=37\pm13$\,MPa acting again the glide of curved dislocation
and an applied stress $\tau_{\rm a}=160\pm40$\,MPa.
This applied stress corresponds to the friction stress acting against the motion of the straight \hkl<a> parts of the dislocation, because the  line tension is zero for a straight dislocation,
thus confirming that this friction stress is much higher for this particular orientation.

\section{Discussion}

\subsection{Glide planes and cross-slip}

Only glide in first-order pyramidal planes has been observed in our experiments,
both in pure zirconium and in Zircaloy-4.
This agrees with previous TEM observations \cite{Akhtar1973,Woo1979,Merle1987,Numakura1991,Long2015a,Long2015b,Long2017,Long2018}
which invariably report the same glide plane for \hkl<c+a> dislocations.
Only Long \textit{et al}.{} \cite{Long2018} have shown that cross-slip in second order pyramidal plane is also possible,
with nevertheless first-order pyramidal plane being the main glide plane.
As the zirconium alloy studied by Long \textit{et al}.{} contains 2.5\,wt.\% Nb, it is possible that
niobium addition promotes the activation of this secondary slip system for \hkl<c+a> dislocations.

Cross-slip events between the two first-order pyramidal planes shared by a single \hkl<c+a> direction
have been   observed in Zircaloy-4 and in pure zirconium in our experiments,
and also in Zr-2.5Nb alloy by Long \textit{et al}.{} \cite{Long2015b,Long2017,Long2018}.
Cross-slip activation at room temperature may look surprising as the dislocation microstructure shows dislocations mostly aligned
in a direction far from the screw orientation, which should prevent cross-slip.
Nevertheless, cross-slip has been observed in our experiments only when the dislocation meets an obstacle,
this obstacle forcing the dislocation to adopt locally a screw orientation
which can then cross-slip.
As \hkl<c+a> dislocations gliding in first-order pyramidal planes are thought to be dissociated in two partials separated by a stable stacking fault \cite{Rodney2017}, the obstacle should also help the dislocation constriction
necessary for its propagation then in the cross-slipped plane,
in the same way as the Friedel-Escaig mechanism operating in face-centered cubic metals \cite{Friedel1964,Escaig1968}.
As less obstacles are present in pure zirconium than in Zircaloy-4, one should expect cross-slip
being less active in pure zirconium.  This looks in agreement with our observations,
although, as mentioned earlier, a detailed statistical analysis would be needed to truly compare
cross-slip activity in the two materials.

\subsection[Straightening along <a> direction]{Straightening along \hkl<a> direction}

One key feature of the glide motion of \hkl<c+a> dislocations is the strong lattice friction
existing for the orientation corresponding to the intersection of the pyramidal glide plane
with the basal plane, leading to long straight dislocations aligned along this \hkl<a> direction.
This important friction exists both in pure zirconium and in Zircaloy-4,
thus showing that it certainly derives from an intrinsic core property of this specific orientation.
As mentioned in the introduction, this behavior is not specific to zirconium, but can be found in other hcp metals.
In Mg, it has been observed that \hkl<c+a> dislocations gliding in second-order pyramidal planes
may lock themselves by a non-conservative dissociation in the basal plane when their line direction belongs to this plane \cite{Stohr1972},
a mechanism which has been well described by elasticity \cite{Agnew2015} and atomistic simulations \cite{Wu2015}.
For \hkl<c+a> dislocations gliding in first-order pyramidal plane,
atomic simulations \cite{Numakura1990a,Numakura1990b,Wu2016a}
predict that the same locking mechanism operates,
with a non conservative dissociation in the basal plane of the \hkl<c+a> in two partial dislocations separated by a basal $I_1$ stacking fault:
$1/3\,\hkl<1-213> \to 1/6\,\hkl<2-203> + 1/6\,\hkl<0-223>$.
Elastic calculations of the energy variation induced by this dissociation predict that
it is also favorable in Zr \cite{Wu2016b}, thus potentially explaining the locking of the \hkl<c+a>
when they meet the basal plane.

However our \textit{in situ}{} experiments provide some clues that the lattice friction
of this \hkl<a> orientation may have a different origin in zirconium.
Locking of edge \hkl<c+a> dislocations in Mg by a non conservative dissociation in the basal plane
has been correlated with the emission of basal prismatic loops with a $I_1$ stacking fault \cite{Stohr1972,Sandlobes2012}.
No such basal $I_1$ loops were detected in our \textit{in situ}{} TEM straining experiments,
neither in pure zirconium nor in Zircaloy-4.
As the emission of these loops is intimately related to the non conservative dissociation
of \hkl<c+a> dislocations in the basal planes, this may be an indication that such a non conservative dissociation
does not operate in zirconium at room temperature.
Besides, contrary to magnesium \cite{Stohr1972},
\hkl<c+a> dislocations intersecting the basal plane are not completely sessile in zirconium
but only suffer a reduced mobility compared to other orientations.
Finally, if \hkl<c+a> dislocations lock themselves in the \hkl<a> direction
by a non conservative dissociation,
\textit{i.e.}{} by climb of partial dislocations in the basal plane,
one would expect that the locking efficiency would depend on the arrest time.
This would lead to dynamic strain ageing and to plastic instabilities
which we did not observe in our \textit{in situ}{} experiments.
The large lattice friction of this \hkl<a> orientation of the \hkl<c+a> dislocation
should therefore result from a different mechanism in zirconium.

This non conservative dissociation is not the solely possible mechanism leading to a lattice friction.
It has been also observed in Mg
that \hkl<c+a> dislocation may emit an \hkl<a> dislocation gliding in the basal plane
to leave behind a sessile \hkl<c> dislocation \cite{Stohr1972}.
Nevertheless, this will lead once again to the permanent locking of the \hkl<c+a> dislocation
and not only to a reduced mobility.
Besides, the extinction tests performed during our \textit{in situ}{} experiments (Fig. \ref{fig:extinction})
clearly indicate that the dislocations aligned along the \hkl<a> direction have a \hkl<c+a> and not a \hkl<c> Burgers vector.
The possible dissociation $\hkl<c+a> \to \hkl<c> +\hkl<a>$ appears therefore not compatible with our observations,
unless the \hkl<c> and \hkl<a> are so tightly bound that they could not separate and should be considered as a single dislocation \cite{Wu2016a}.

Without going to the emission of a perfect basal \hkl<a> dislocation,
Numakura \textit{et al}. \cite{Numakura1990a,Numakura1990b} have shown that \hkl<c+a> dislocations
intersecting the basal plane can emit a single Shockley partial trailing a basal $I_2$ stacking fault.
The corresponding dissociation, $1/3\,\hkl<1-213> \to 1/3\,\hkl<0-110> + 1/3\,\hkl<1-103>$,
does not require climb of the partial dislocations and
leads to a non planar core with a stacking fault in the basal and in the pyramidal planes.
Atomistic simulations relying on a generic interatomic potential for hcp metals \cite{Numakura1990a,Numakura1990b}
show that this core is more stable than the planar core dissociated in the pyramidal plane,
but that it recombines in this planar core under an applied stress
to easily glide then in the pyramidal plane.
This non planar core could therefore explain the lattice friction of the \hkl<a> orientation.
Its ability to transform in a planar glissile core is also compatible with the jerky glide motion
sometimes seen for this \hkl<a> orientation.

\subsection{Macro-kinks}

One characteristic of the gliding \hkl<c+a> dislocations when they are aligned in the \hkl<a> direction
is the presence of macro-kinks on these otherwise straight segments.
Recent \textit{in situ}{} TEM compression experiments performed in Mg \cite{Zhang2019} have evidenced the same macro-kinks
on the straight portions of \hkl<c+a> dislocations, leading to ``stair-like'' shapes.
Although \hkl<c+a> dislocations are not gliding on the same pyramidal plane in Zr and Mg,
their glide motion induces the formation of macro-kinks in both metals.
We can reasonably exclude that these macro-kinks are a by-product of a reaction with \hkl<a> dislocations
as extinction tests in TEM observations showed that they have the same \hkl<c+a> Burgers vector
as the whole dislocation lines.
They could arise from a local change of configurations of the \hkl<c+a> dislocation,
between an almost sessile core corresponding to the straight segments
and a glissile core for the macro-kinks.
These macro-kinks are less numerous in pure zirconium than in Zircaloy-4,
leading to longer straight dislocations in the former case.
Their interaction with the different solute atoms present in solid solution hinder their motion,
thus leading to a stronger pinning of these macro-kinks in Zircaloy-4 than in pure zirconium.

\subsection{Lattice friction}

Our line tension measurements indicate a friction stress of $160\pm40$\,MPa acting against glide of the straight part
of \hkl<c+a> dislocations in pure zirconium.
This friction has both an intrinsic origin probably due to a complex core of the \hkl<c+a> dislocation
for the \hkl<a> orientation as discussed before and an extrinsic one arising from interaction with solute atoms.
This friction stress of the less mobile segments should correspond to the yield stress
of \hkl<c+a> pyramidal slip at room temperature.
Modeling with crystal plasticity their micro-cantilevers bending experiments,
Gong \textit{et al}.{} \cite{Gong2015} obtained for the same slip system a much higher value, $532\pm58$\,MPa,
in a commercial pure zirconium.
With a higher level of solute elements in commercial pure zirconium than in the pure zirconium
used in our experiments, in particular oxygen concentration which is four times higher,
this yield stress difference indicates a strong influence of alloying elements.
This is further supported by the strongest pinning of gliding dislocations
observed in Zircaloy-4 than in pure zirconium.

This reduced mobility of \hkl<c+a> dislocations in alloyed zirconium
may appear at first glance in contradiction with the study of Jensen and Backofen \cite{Jensen1972}
who observed that plastic strain along the $\hkl<c>$ axis was mostly accommodated by twinning below 300\,{$^{\circ}$}C in pure zirconium
whereas only $\hkl<c+a>$ slip was activated in Zircaloy-4 with very few compressive twins.
Although solute elements further hinder glide of $\hkl<c+a>$ dislocations, this hardening contribution appears less important
than the impact of the same solute elements on twinning with the suppression of twinning by impurities in zirconium \cite{Garde1973},
in particular oxygen which addition strongly decreases twin activity \cite{Viltange1985}.

\section{Conclusions}

\textit{In situ}{} TEM straining experiments performed both in pure zirconium and in Zircaloy-4
have shown that \hkl<c+a> dislocations glide exclusively in first-order pyramidal planes
at room temperatures, with some cross-slip events between two pyramidal planes
sharing the same \hkl<c+a> direction.
A characteristic feature of the microstructure is that \hkl<c+a> dislocations align in the \hkl<a> direction
defined by the intersection of their glide plane with the basal plane,
leading to long straight dislocations.
Dislocations segments aligned along this orientation experience a higher lattice friction than others,
with a friction stress estimated to $160\pm40$\,MPa for the \hkl<a> orientation
instead of $37\pm13$\,MPa for others in pure zirconium.
This \hkl<a> orientation exhibits two different glide motions, either viscous or jerky,
and its frequent pinning in localized points leads to the creation of macro-kinks.
The same dynamic behavior of \hkl<c+a> dislocations has been observed
in pure zirconium and in Zircaloy-4, except for a highest friction
acting against the motion of all characters in Zircaloy-4 because
of the numerous alloying elements in solid solution.
The mechanisms controlling the mobility of \hkl<c+a> dislocations in zirconium
show strong similarities with the ones operating in other hcp metals,
in particular magnesium despite a different glide plane.

\vspace{0.5cm}
\linespread{1}
\small

\textbf{Acknowledgments} -
The authors thank Framatome for providing the raw Zr sponge and the Zircaloy-4 TREX tube,
D. Nunes and S. Urvoy (SRMA/CEA) for the preparation of the pure Zr samples
and B. Arnal (SRMA/CEA) for thin foils preparation.
This work is funded by the French Tripartite Institute (CEA-EDF-Framatome) through the GAINE project.

redpajama_set_name: RedPajamaArXiv


text:
posted Jan 15, 2012, 6:29 PM by Bill Duncan
Public money, private schools
Bill puts local education in peril
By Bill Duncan / For the Monitor
You might think that Beaver Meadow and the Kimball-Walker schools are Concord's local public schools - precious assets to be nurtured, debated over, improved.
But our Free State/Tea Party Legislature, the one we elected in 2010, says, no, these are "government-run" schools, from which parents and children need to be protected. You know, as in, "We need less government!" Make no mistake about it. They talk about "choice," but their mission is to shut down our public schools and replace them with religious and home schools.
Republican Sen. James Forsythe of Strafford has proposed an education tax credit program does just that. It is essentially a voucher program in which the state would spend $100 million in the first five years, sending children to private schools. It would take students and money out of the public schools and aggravate already declining enrollment. It would do real damage to local public schools immediately.
And Forsythe wants us to help pay for it with our property taxes. Here's how the program would work.
New Hampshire businesses would make scholarship money available to new "scholarship organizations." Then those businesses could subtract 75 percent of what they donated to the program from their New Hampshire business profits tax.
The state would spend $15 million on those tax credits in the first year, but the program could grow every year to as much as $36 million in the fifth year.
The scholarship organizations would give scholarships averaging $2,500 to students. Scholarships of up to $1,500 could go for home-schooling costs.
The scholarships would not be for low-income families who might need the money, but instead would be available to anyone at any income level - including private school students who have never attended public schools.
The scholarships would not be targeted to getting children out of "failing" schools, as many voucher programs are. Supporters argue that our whole New Hampshire education system is failing.
There would be no academic accountability for the schools receiving the money, most of which would likely go to religious schools.
The program starts out huge, funding 7,000 to 8,000 students in the first year, and could grow to 17,000 students in the fifth year. There are only 190,000 students in the whole state.
Where does the money come from? Our pockets. It's a shell game. The business funds the scholarships. Then the state repays the businesses with the tax credits. Then property taxpayers repay the state!
For instance, Mary Jones lives in Concord and attends Beaver Meadow School. Currently, the state gives the Concord School District an average of $3,000 in state aid toward her education each year. We pay another $1,000 in statewide education property tax to make up the $4,000 that's called the "cost of an adequate education." Then we add more local property tax money and that's the school's budget. (Leave the donor town question aside for this discussion.)
Under Forsythe's bill, Mary could get a $2,500 scholarship and leave to go to a private school. The state would repay the business 75 percent of that, a little less than $2,000. Then the state would take back from us not just the $3,000 state aid, but the whole $4,000 "cost of adequacy" grant, including the $1,000 that comes from our property tax.
Two alternatives
Just to be clear, there's real complexity behind all this, and advocates of privatizing education will try to confuse the issue. But there are only two alternatives. One is that the $100 million in tax credits the state spends on vouchers comes straight out of New Hampshire tax revenue.
The other is that money comes from state aid and property taxes that currently fund our schools, so the sponsors can assert that the program is "revenue neutral" at the state level.
Either way, in these times when the Legislature is cutting the University of New Hampshire, the Department of Health and Human Services and everything in sight, the citizens of New Hampshire have no need for a large, complex program to fund scholarships for families that do not need them.
This is an expensive new public program with no legitimate public purpose. We already have great teachers and a great education system, among the best in the country. We don't need to shut it down and replace it with a privatized patchwork of religious and home schools.
If you don't like this idea, tell your legislator. Tell the governor. Track the issue at the website Defending New Hampshire Public Education, dnhpe.org.
(Bill Duncan of New Castle is a retired software entrepreneur.)


0:00 Matthew Escareno at goalie for Oxnard.
0:00 Bailey Brogan at goalie for Fresno.
9:01 Shot by Oxnard Raul Gonzalez, High.
20:31 Shot by Fresno Jose Ramirez, Blocked.
21:20 Shot by Fresno Jose Torres, Wide.
22:13 Shot by Oxnard Jaime Mendez, Blocked.
28:57 Substitution!Oxnard substitution: Miguel Coronel for Damian Lopez.
35:31 Shot by Oxnard Kiefer Cooksey, High.
36:32 Shot by Oxnard Jorge Jimenez, Save (by goalie) Bailey Brogan.
38:12 Shot by Oxnard Luis Carbajal, Blocked.
39:12 Substitution!Oxnard substitution: Miguel Coronel for Damian Lopez.
41:40 Shot by Oxnard Luis Carbajal, Save (by goalie) Bailey Brogan.
42:03 Shot by Oxnard Kiefer Cooksey, Wide.
44:34 Corner kick by Oxnard Saul Higuera [44:34].
45:00 Corner kick by Fresno Jose Tovar [45:00].
47:15 Shot by Fresno Noel Prieto, Wide.
58:31 Shot by Fresno Jose Torres, Save (by goalie) Matthew Escareno.
58:49 Corner kick by Fresno Juan Lazcano [58:49].
59:08 Shot by Fresno Andres Rios, Blocked.
70:30 Shot by Oxnard Raul Gonzalez, Save (by goalie) Bailey Brogan.
80:51 Corner kick by Fresno Noel Prieto [80:51].
81:04 Shot by Fresno Andres Rios, Blocked.
81:56 Shot by Fresno Jesse Camarena, Save (by goalie) Matthew Escareno.
87:31 Shot by Oxnard Luis Cabrera, Save (by goalie) Bailey Brogan.
87:37 Corner kick by Oxnard Team [87:37].
88:51 Shot by Oxnard Luis Cabrera, Blocked.
89:36 Shot by Oxnard Miguel Coronel, Wide.